Joining Technology and Application of Advanced Materials 9811996881, 9789811996887

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Advanced and Intelligent Manufacturing in China

Yajiang Li

Joining Technology and Application of Advanced Materials

Advanced and Intelligent Manufacturing in China Series Editor Jie Chen, Tongji University, Shanghai, Shanghai, China

This is a set of high-level and original academic monographs. This series focuses on the two fields of intelligent manufacturing and equipment, control and information technology, covering a range of core technologies such as Internet of Things, big data, 3D printing, robotics, intelligent equipment, industrial network security, and artificial intelligence, and epitomizing the achievements of technological development in China’s manufacturing sector. With Prof. Jie Chen, a member of the Chinese Academy of Engineering and a control engineering expert in China, as the Editorial in Chief, this series is organized and written by more than 80 young experts and scholars from more than 40 universities and institutes. It typically embodies the technological development achievements of China’s manufacturing industry. It will promote the research and development and innovation of advanced intelligent manufacturing technologies, and promote the technological transformation and upgrading of the equipment manufacturing industry.

Yajiang Li

Joining Technology and Application of Advanced Materials

Yajiang Li Shandong University Jinan, Shandong, China

ISSN 2731-5983 ISSN 2731-5991 (electronic) Advanced and Intelligent Manufacturing in China ISBN 978-981-19-9688-7 ISBN 978-981-19-9689-4 (eBook) https://doi.org/10.1007/978-981-19-9689-4 Jointly published with Chemical Industry Press The print edition is not for sale in China (Mainland). Customers from China (Mainland) please order the print book from: Chemical Industry Press. The translation was done with the help of artificial intelligence (machine translation by the service DeepL.com). A subsequent human revision was done primarily in terms of content. © Chemical Industry Press 2023 This work is subject to copyright. All rights are solely and exclusively licensed by the Publisher, whether the whole or part of the material is concerned, specifically the rights of reprinting, reuse of illustrations, recitation, broadcasting, reproduction on microfilms or in any other physical way, and transmission or information storage and retrieval, electronic adaptation, computer software, or by similar or dissimilar methodology now known or hereafter developed. The use of general descriptive names, registered names, trademarks, service marks, etc. in this publication does not imply, even in the absence of a specific statement, that such names are exempt from the relevant protective laws and regulations and therefore free for general use. The publishers, the authors, and the editors are safe to assume that the advice and information in this book are believed to be true and accurate at the date of publication. Neither the publishers nor the authors or the editors give a warranty, expressed or implied, with respect to the material contained herein or for any errors or omissions that may have been made. The publishers remain neutral with regard to jurisdictional claims in published maps and institutional affiliations. This Springer imprint is published by the registered company Springer Nature Singapore Pte Ltd. The registered company address is: 152 Beach Road, #21-01/04 Gateway East, Singapore 189721, Singapore

Preface

The application of advanced materials joining technology has been producing significant economic and social benefits and is worthy of great promotion. This book provides a systematic description of the joining mechanism, weldability, technical points and applications of the advanced materials, e.g., high-tech ceramics, intermetallic compounds, composite materials, functional materials, etc. Some application examples of typical engineering welding structure are shown to guide the development of new products. The content of this book reflects the development of advanced material joining technology in recent years, especially some high-tech development, which is of great significance to promote the welding application of advanced materials. This book is intended for engineers and technicians engaged in careers related to materials development and welding technology, and also for teachers or students in higher education institutions, and scientific researchers in enterprises and institutions. The manufacturing industry is the mainstay of the national economy and is the foundation of the country, the instrument of its prosperity and the basis of its strength. In the past decade, China’s manufacturing industry has continued to develop rapidly. Its comprehensive strength and international status has been greatly enhanced and it has become the largest manufacturing country in the world. However, China is still in the process of industrialization, and the problem of large but not strong is prominent. There is still a large gap compared with advanced countries. In order to solve the problems of large but not strong manufacturing industry, weak independent innovation capability, high dependence on key core technologies and high-end equipment, the State Council released the national plan of “Made in China 2025” on May 8, 2015. Subsequently, the Ministry of Industry and Information Technology released the “Made in China 2025” plan, which put forward the “three-step” development strategy of China’s manufacturing industry and the goals, guidelines and strategic lines for 2025, and formulated nine strategic tasks as well as ten key development areas. On August 19, 2016, the Ministry of Industry and Information Technology, the Development and Reform Commission, the Ministry of Science and Technology and the Ministry of Finance jointly released five project implementation guidelines, including the “Made in China 2025” manufacturing innovation center, industrial v

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foundation, green manufacturing, intelligent manufacturing and high-end equipment innovation. In response to the major strategic deployment made by the Party Central Committee and the State Council to build a strong manufacturing country, governments, enterprises and scientific research departments all over the world are actively exploring and deploying. Accelerating the development of the integration of new generation information technology and manufacturing technology, promoting the transformation of China’s manufacturing mode from “Made in China” to “Intellectually Manufactured in China”, and accelerating the transformation of China’s manufacturing industry from big to strong, are becoming our new historical missions. At present, the process of information revolution continues to evolve rapidly, and technologies such as Internet of Things, cloud computing, big data and artificial intelligence are widely penetrated in various fields of economy and society, and the prosperity of information economy has become an important symbol of national strength. Technologies in the fields of additive manufacturing (3D printing), robotics and intelligent manufacturing, control and information technology, artificial intelligence and other areas continue to make major breakthroughs, pushing the traditional industrial system to differentiate and change, and will reshape the international division of labor in manufacturing. The integration and development of manufacturing technology with the Internet and other information technology has become a major trend and main feature of the new round of scientific and technological revolution and industrial change. Against this background of great development and change in China’s manufacturing industry, Chemical Industry Press has taken the initiative to follow the development trend of technology and industry and organized the publication of the “Made in China 2025” publication project series, which can be regarded as courageous and leading at the right time. The “Made in China 2025” publication project is a series of books closely focused on the “Made in China 2025”, the action program of the first decade of the implementation of the manufacturing power strategy issued by the State Council. It is a set of high-level and original academic monographs. Based on the two major fields of intelligent manufacturing and equipment, control and information technology, the series covers a series of core technologies such as Internet of Things, big data, 3D printing, robotics, intelligent equipment, industrial cyber security, knowledge automation and artificial intelligence. The selection plan of the series is closely integrated with the “Made in China 2025” plan and 11 supporting implementation guides, action plans or special plans, and the content of each fascicle is organized for some core technologies in each field, focusing on the technological development achievements in the domestic manufacturing industry, aiming to strengthen the research and development, promotion and application of advanced technologies, and provide a platform for targeted directional guidance and systematic technical reference for the implementation of the “Made in China 2025” action program. The set focuses on the following key features. First of all, the contents of the series are striving for originality, with networked and intelligent technology as the core, bringing together many cutting-edge technologies, reflecting some of the latest technological achievements at home and abroad,

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especially reflecting the relevant original technological achievements in China. These books contain many new technologies which have won many national and provincial science and technology awards, and the publication of the books is very helpful to the promotion and application of new technologies! Thee contents of books not only solve practical problems for technical personnel, but also provide new directions and expand new ideas for research. Secondly, while introducing new technologies, theories and methods in the corresponding professional fields, each fascicle of the series gives priority to new technologies with application prospects and examples of their promotion and application, so as to promote the transformation of outstanding scientific research results into industry. The series is led by Sun Youxian, an academician of the Chinese Academy of Engineering, who is the director of the editorial committee. Wu Cheng, Wang Tianran, Zheng Nanning and many other academicians participate in the planning and organizing work. Many middle-aged and young scholars, also participate in the specific writing work, promoting academic level and writing quality. It is believed that the publication of this series is of positive significance to the implementation of the national strategic plan of “Made in China 2025”, which can effectively promote the R&D and innovation of China’s intelligent manufacturing technology, the technological transformation and upgrading of the equipment manufacturing industry, and improve the design capability and technical level of products. Thus, the core competitiveness of Chinese manufacturing industry can be enhanced from multiple angles. The emergence of each new material in history has been accompanied by the emergence of new joining processes, which has also driven the development of science and technology. The research and development of advanced materials is the result of multidisciplinary interpenetration and joining technology plays a crucial role in the promotion and application of advanced materials in fields of electronics, energy, automotive, aerospace and nuclear industries. The development of advanced materials is an important material basis for developing high technology. The welding and joining of advanced materials in engineering structures is frequently encountered. Many problems arise in practice, and sometimes even hinder the progress of the entire project. In particular, it is difficult to achieve the joining of many advanced materials using conventional welding methods. Thus, the superiority of advanced welding technology is increasingly prominent. In line with the National Manufacturing Power Strategy of “Made in China 2025” and the development of advanced materials, this book combines theory and practice to address the problems of joining advanced materials (such as high-tech ceramics, intermetallic compounds, composite materials and functional materials) that have attracted attention in recent years. The book aims to highlight the characteristics of science, advancement and novelty. The content reflects the development of the joining technology of advanced materials in recent years, especially some high-tech developments, which are of great significance in promoting the welding applications of advanced materials. Some application examples of structural joining of advanced materials are given in the book, which can guide the development of new products.

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This book is not only for engineers and technicians engaged in materials development and welding technology, but also for the reference of teachers and students of higher education institutions, research institutes and scientific researchers of enterprises and institutions. Other participants in the writing of this book include Wang Juan, Xia Chunzhi, Shen Xiaoqin, Ma Qunshuang, Liu Kun, Wu Na, Ma Haijun, Wei Shouzheng and Li Jianing, et al. Due to the limitations of the author’s level, the shortcomings in the book are inevitable, and we are grateful for the reader’s criticism and correction. Jinan, China

Yajiang Li

Contents

1 Overview . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1.1 Classification and Performance Characteristics of Advanced Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1.1.1 Classification of Advanced Materials . . . . . . . . . . . . . . . . . . . 1.1.2 Performance Characteristics of Advanced Materials . . . . . . . 1.2 Applications and Development Prospects of Advanced Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1.2.1 Advanced Ceramics . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1.2.2 Intermetallic Compounds . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1.2.3 Laminated Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1.2.4 Composite Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1.2.5 Functional Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Bibliography . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2 Welding of Advanced Ceramic Materials . . . . . . . . . . . . . . . . . . . . . . . . . 2.1 Performance Characteristics and Joining Problems of Ceramic Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.1.1 Performance Characteristics of Structural Ceramics . . . . . . . 2.1.2 Basic Requirements for Ceramic-to-Metal Joining . . . . . . . . 2.1.3 Problems with Ceramic-to-Metal Joining . . . . . . . . . . . . . . . . 2.1.4 Joining Methods of Ceramics and Metal . . . . . . . . . . . . . . . . . 2.2 Weldability Analysis of Ceramic Materials . . . . . . . . . . . . . . . . . . . . . 2.2.1 Welding Stress and Cracks . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.2.2 Interfacial Reactions and Interface Formation Processes . . . 2.2.3 Bond Strength at the Diffusion Interface . . . . . . . . . . . . . . . . . 2.3 Brazing of Ceramic to Metal Joints . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.3.1 Characteristics of Ceramic-to-Metal Brazed Joints . . . . . . . . 2.3.2 Surface Metallization Brazing of Ceramics and Metals . . . . 2.3.3 Active Metallization Brazing of Ceramics to Metals . . . . . . . 2.3.4 Examples of Ceramic-to-Metal Brazing . . . . . . . . . . . . . . . . .

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2.4 Diffusion Bonding of Ceramics to Metals . . . . . . . . . . . . . . . . . . . . . . 2.4.1 Characteristics of Ceramic-to-Metal Diffusion Bonding . . . . 2.4.2 Process Parameters for Diffusion Bonding . . . . . . . . . . . . . . . 2.4.3 Characteristics of the Al2 O3 Composite Ceramic/metal Diffusion Interface . . . . . . . . . . . . . . . . . . . . . . 2.4.4 Diffusion Bonding of SiC/Ti/SiC Ceramics . . . . . . . . . . . . . . 2.5 Electron Beam Welding of Ceramics to Metals . . . . . . . . . . . . . . . . . . 2.5.1 Characteristics of Electron Beam Welding of Ceramics and Metals . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.5.2 Processes for Electron Beam Welding of Ceramics to Metals . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.5.3 Example of Electron Beam Welding of Ceramics to Metals . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Bibliography . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3 Diffusion Welding of Composite Ceramics to Steel . . . . . . . . . . . . . . . . . 3.1 Diffusion Welding Process of Composite Ceramics to Steel . . . . . . . 3.1.1 Basic Properties of Al2 O3 –TiC Composite Ceramics . . . . . . 3.1.2 Process Characteristics of Composite Ceramic to Steel Diffusion Welding . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.1.3 Specimen Preparation and Test Methods for Diffusion Joints . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2 Diffusion Welding of Al2 O3 –TiC Composite Ceramics and Q235 Steel . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2.1 Interfacial Characteristics and Microhardness of Al2 O3 –TiC/Q235 Diffusion Welded Joint . . . . . . . . . . . . . 3.2.2 Shear Strength of Al2 O3 –TiC/Q235 Diffusion Joint . . . . . . . 3.2.3 Microstructure of Al2 O3 –TiC/Q235 Diffusion Welded Joint . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2.4 Analysis of Precipitated Phases in the Interface Transition Zone . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2.5 Effect of Process Parameters on the Microstructure of Al2 O3 –TiC/Q235 Diffusion Interface . . . . . . . . . . . . . . . . . 3.3 Diffusion Welding of Al2 O3 –TiC Composite Ceramics with 18-8 Austenitic Steel . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.3.1 Interfacial Characteristics and Microhardness of the Al2 O3 –TiC/18-8 Diffusion Welding Joint . . . . . . . . . . 3.3.2 Shear Strength of Al2 O3 –TiC/18-8 Diffusion Joint . . . . . . . . 3.3.3 Microstructure of Al2 O3 –TiC/18-8 Diffusion Welded Joint . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.3.4 Analysis of Precipitated Phases in the Interface Transition Zone . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.3.5 Effect of Process Parameters on the Microstructure of the Al2 O3 –TiC/18-8 Diffusion Interface . . . . . . . . . . . . . . .

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3.4 Diffusion Welding of Al2 O3 –TiC Composite Ceramics and W18Cr4V High-Speed Steel . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.4.1 Diffusion Process Characteristics and Specimen Preparation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.4.2 Interfacial Characteristics of Al2 O3 –TiC/W18Cr4V Diffusion Welded Joint . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.4.3 Shear Strength of the Al2 O3 –TiC/W18Cr4V Diffusion Welded Interface . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.4.4 Effect of Process Parameters on the Microstructure of the Interface Transition Zone . . . . . . . . . . . . . . . . . . . . . . . . 3.4.5 Crack Extension and Fracture Characteristics at the Al2 O3 –TiC/W18Cr4V Diffusion Welded Interface . . . Bibliography . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4 Joining of Nickel–Aluminium and Titanium–Aluminium Intermetallic Compounds . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.1 Development and Properties of Intermetallic Compounds . . . . . . . . 4.1.1 Development of Intermetallic Compounds for Structures . . . 4.1.2 Basic Properties of Intermetallic Compounds . . . . . . . . . . . . 4.1.3 Three Promising Intermetallic Compounds . . . . . . . . . . . . . . 4.1.4 Superplasticity of Ni–Al and Ti–Al Intermetallic Compounds . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.2 Welding of Ni–Al Intermetallic Compounds . . . . . . . . . . . . . . . . . . . . 4.2.1 Diffusion Bonding of NiAl Alloys . . . . . . . . . . . . . . . . . . . . . . 4.2.2 Fusion Welding of Ni3 Al Intermetallic Compounds . . . . . . . 4.2.3 Diffusion Bonding of Ni3 Al to Carbon Steel (or Stainless Steel) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.2.4 Diffusion Bonding and Vacuum Brazing of Ni3 Al (IC10) Alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.3 Welding of Ti–Al Intermetallic Compounds . . . . . . . . . . . . . . . . . . . . 4.3.1 Welding Characteristics of Ti–Al Intermetallic Compounds . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.3.2 Arc Welding of TiAl Intermetallic Compounds . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.3.3 Electron Beam Welding of TiAl Intermetallic Compounds . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.3.4 Diffusion Welding of TiAl and Ti3 Al Alloys . . . . . . . . . . . . . 4.3.5 Diffusion Bonding of TiAl Dissimilar Materials . . . . . . . . . . Bibliography . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5 Joining of Iron-Aluminium Intermetallic Compounds . . . . . . . . . . . . . 5.1 Iron-Aluminium Intermetallic Compounds and Its Weldability . . . . 5.1.1 Characteristics of Iron-Aluminium Intermetallic Compounds . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.1.2 Weldability Characteristics of Iron-Aluminium Intermetallic Compounds . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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5.1.3 Cracking in the Fe3 Al Welded Joint Area . . . . . . . . . . . . . . . . 5.2 Wire-Filled Tungsten Arc Welding of Fe3 Al and Steel (Q235, 18–8 Steel) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.2.1 Characteristics of the Tungsten Arc Welding Process of Fe3 Al and Steel . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.2.2 Microstructure Characteristics of the Fe3 Al/Steel Joint Zone of Filled Wire GTAW . . . . . . . . . . . . . . . . . . . . . . . 5.2.3 Microhardness of the Fe3 Al/Steel Filled Wire GTAW Joint . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.2.4 Shear Strength and Fracture Morphology of Fe3 Al/Steel GTAW Joints . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.3 Vacuum Diffusion Welding of Fe3 Al to Steel (Q235, 18-8 Steel) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.3.1 Process Characteristics of Fe3 Al/Steel Vacuum Diffusion Welding . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.3.2 Shear Strength of the Fe3 Al/Steel Diffusion Weld Interface . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.3.3 Microstructural Characteristics of the Fe3 Al/Steel Diffusion Weld Interface . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.3.4 Microhardness of Fe3 Al/Steel Diffusion Welded Joints . . . . 5.3.5 Element Diffusion Near the Interface and Transition Zone Width . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.3.6 Effect of Process Parameters on the Interface Characteristics of Diffusion Welding . . . . . . . . . . . . . . . . . . . . 5.4 Other Welding Methods of Fe3 Al Intermetallic Compounds . . . . . . 5.4.1 Electron Beam Welding of Fe3 Al Intermetallic Compounds . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.4.2 Electrode Arc Welding of Fe3 Al . . . . . . . . . . . . . . . . . . . . . . . . 5.4.3 Argon Arc Overlay Welding and Characteristics of Fe3 Al . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Bibliography . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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6 Welding of Laminated Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.1 Characteristics and Weldability of Laminated Materials . . . . . . . . . . 6.1.1 Characteristics of Laminated Materials . . . . . . . . . . . . . . . . . . 6.1.2 Weldability Analysis of the of Laminated Materials . . . . . . . 6.1.3 Research Status of Laminated Materials Welding . . . . . . . . . 6.2 Wire-Filled GTAW of Laminated Materials . . . . . . . . . . . . . . . . . . . . 6.2.1 Process Characteristics of Wire-Filled GTAW of Laminated Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.2.2 Fusion State of Welding Zone of Laminated Materials . . . . . 6.2.3 Microstructure and Properties of Joint Between Laminated Material and 18–8 Steel . . . . . . . . . . . . . . . . . . . . . 6.2.4 Process Characteristics of Diffusion Brazing of Laminated Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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6.2.5 Bonding Interface of Laminated Composite/18–8 Steel Diffusion Brazed Joint . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.2.6 Microhardness of Laminated Composite/18–8 Steel Diffusion Brazed Joint . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.2.7 Shear Strength of Laminated Composite/18–8 Steel Diffusion Brazed Joint . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Bibliography . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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7 Welding of Advanced Composites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7.1 Classification, Characteristics and Properties of Composite Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7.1.1 Classification and Characteristics of Composite Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7.1.2 Reinforcement of Composite Materials . . . . . . . . . . . . . . . . . . 7.1.3 Performance Characteristics of Metal Matrix Composites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7.2 Analysis of the Weldability of Composite Materials . . . . . . . . . . . . . 7.2.1 Weldability of Metal Matrix Composites . . . . . . . . . . . . . . . . 7.2.2 Weldability of Resin Matrix Composites . . . . . . . . . . . . . . . . 7.2.3 Weldability of C/C Composites . . . . . . . . . . . . . . . . . . . . . . . . 7.2.4 Weldability of Ceramic Matrix Composites . . . . . . . . . . . . . . 7.3 Welding of Continuous Fiber Reinforced Metal Matrix Composites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7.3.1 Problems in Welding of Continuous Fiber Reinforced MMC . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7.3.2 Joint Form Design for Continuous Fiber Reinforced MMC . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7.3.3 Characteristics of Welding Process for Fiber Reinforced MMC . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7.4 Welding of Discontinuously Reinforced Metal Matrix Composites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7.4.1 Welding Problems of Discontinuously Reinforced MMC . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7.4.2 Welding Process Characteristics of Discontinuously Reinforced MMC . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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8 Connection of Functional Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8.1 Connection of Superconducting Materials to Metals . . . . . . . . . . . . . 8.1.1 Performance Characteristics and Applications of Superconducting Materials . . . . . . . . . . . . . . . . . . . . . . . . . . 8.1.2 Connection Methods for Superconducting Materials . . . . . . 8.1.3 Characteristics of the Joining Process for Superconducting Materials . . . . . . . . . . . . . . . . . . . . . . . . . 8.1.4 Welding of Oxide Ceramic Superconducting Materials . . . .

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8.2 Shape Memory Alloy to Metal Connection . . . . . . . . . . . . . . . . . . . . . 8.2.1 Characteristics and Applications of Shape Memory Alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8.2.2 Advances in the Welding of Shape Memory Alloys . . . . . . . 8.2.3 Resistance Brazing of TiNi Shape Memory Alloys . . . . . . . . 8.2.4 Transition Liquid Phase Diffusion Welding of TiNi Alloy to Stainless Steel . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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Chapter 1

Overview

Advanced materials are a class of engineering materials that have superior properties to traditional steel and non-ferrous materials. They can meet the needs of hightech development, such as high-tech ceramics, intermetallic compounds, laminated materials, composites, etc. Welding of advanced materials is frequently encountered more problems, sometimes even hindering the progress of the entire development and engineering application (welded structures). The main characteristics of advanced materials are good performance, high hardness and difficulty in welding.

1.1 Classification and Performance Characteristics of Advanced Materials With the development of modern science and technology, the requirement for quality of welded joint and structural performance is increasingly high. The welding process of steel and conventional non-ferrous metals has been difficult to meet the requirements for high-tech development. A variety of advanced and special materials has been emerging and been interested in recent years, which can greatly promote the progress of science and technology and the development of society including electronics, energy, automotive, aerospace, nuclear industry and so on.

1.1.1 Classification of Advanced Materials Advanced materials are engineering materials with special properties and application that have been developed or are being developed, such as high-tech ceramics, intermetallic compounds, composite materials, etc. Advanced materials with more excellent performance than traditional materials and they are closely related to the © Chemical Industry Press 2023 Y. Li, Joining Technology and Application of Advanced Materials, Advanced and Intelligent Manufacturing in China, https://doi.org/10.1007/978-981-19-9689-4_1

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development of high technology. Developing advanced materials is the technology of creating new materials that can meet various needs through a series of research progress such as improving of physical chemistry, material design or processing, and test or evaluation. Advanced materials are divided into four categories according to the properties of materials: advanced metallic materials, inorganic non-metallic materials (such as ceramic, etc.), organic polymer, and advanced composite materials. According to performance, advanced materials are involved in structural materials and functional materials. Structural materials mainly use their mechanical, physical and chemical properties to meet the requirements of high strength, good toughness, high hardness and excellent resistance to high temperature, wear, corrosion or irradiation etc. Functional materials mainly use their electrical, magnetic, acoustic, optical, thermal and other effects to achieve a certain function, such as superconducting materials, magnetic materials, photosensitive materials, thermo sensitive materials, stealth materials and nuclear materials. Advanced materials play a significant role in national defense construction. For example, the successful development of ultra-pure silicon and gallium arsenide has promoted the advent of large-scale and ultra-largescale integrated circuits, which increased the speed of computer operations from hundreds of thousands of times per second to more than 10 billion times per second. The operating temperature of aero-engine can make the thrust increased by 24% in the enhancement of 100 °C. Stealth materials can absorb electromagnetic waves or reduce the infrared radiation of weapons and equipment, making detection systems from enemy difficult to work. The development and application of advanced materials is an important part of the development of modern science and technology. With the development of aerospace, new energy, electric power and other industries, demand on the performance of materials is becoming higher and higher. The development of advanced materials for application in some special conditions is also one of the trends in the development of science and technology, in which, structural materials is an essential part of it. Advanced materials are wide-ranging and in a constant state of development and application, mainly involved in engineering including: high-tech ceramics, intermetallic compounds, laminated materials, composite materials, functional materials, etc. A prominent feature of these materials is their high hardness and strength, poor plasticity and toughness, which makes them difficult to be welded with conventional fusion welding technology. The development and application of advanced materials is closely related to the improvement of high technology and it has a unique and irreplaceable role. For example, the development and application of advanced ceramic materials, intermetallic compounds and refractory materials will provide an important basis for the development of energy, the development of space and the oceans, and the exploration of aerospace and other fields. Advanced materials are the necessary material basis on the development of high technology and the precursors of new technological revolutions.

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1.1.2 Performance Characteristics of Advanced Materials From the perspective of the synthesis and manufacturing process of advanced materials, advanced ceramics, intermetallic compounds, laminated materials and composite materials are obtained by high-tech means (such as ultra-high pressure, ultra-high temperature, ultra-high speed cooling rate, etc.) under some extreme conditions. At the same time, the development of advanced ceramics, intermetallic compounds and composite materials is closely related to the development and application of computer and automatic control technology that require very strict quality control of the materials. Advanced materials are developing with more excellent properties and special applications such as high strength, resistance to high temperature, corrosion and oxidation so on. (1) Advanced ceramic Also known as high technology ceramics, newly developed ceramics are a new generation of ceramics with excellent performance, which are obtained from refined high purity, ultra-fine synthetic inorganic compounds by a precisely controlled preparation process. Ceramics are inorganic non-metallic materials synthesized from oxides, nitrides, carbides and silicide of various metals with appropriate ingredients after forming and sintering at high temperature. Advanced ceramics are very different from traditional ceramics in terms of composition, properties, manufacturing process and application. Their composition has been developed from the original SiO2 , Al2 O3 , MgO to Si3 N4 , SiC and ZrO2 . Advanced physical and chemical methods were used to prepare ultrafine powders. Sintering has also developed from ordinary atmospheric sintering to advanced sintering methods such as hot-pressure sintering with controlled atmosphere of vacuum and microwave. Advanced ceramics have specific fine microstructure and properties to make them play an important role in the development of modern engineering and high technology. Advanced ceramics in a broad sense include artificial mono-crystalline and amorphous (glass) ceramics and their composites, semiconductors, refractory materials, etc., which belongs to inorganic non-metallic materials. Ceramics are generally divided into two categories: functional and structural ceramics. Bio-ceramics can be classified as functional ceramics. Ceramics related to welding technology are mainly structural ceramics. Advanced ceramics have excellent physical and mechanical properties, such as high strength, high hardness, good resistance to wear, corrosion, high temperature and thermal shock and some unique functions in electricity, magnetism, heat, light and sound. Compared with metal, the coefficient of linear expansion of ceramics is relatively low and is in the range of 10–5 ~ 10–6 K−1 . The melting point (or sublimation, decomposition temperature) is much higher and even some ceramics can work at a high temperature of 2000 ~ 3000 °C and maintain the strength at room temperature while most metals often lose the high strength above 1000 °C. Therefore, ceramics have a broad prospect as a high-temperature structural material for aerospace engines, cutting tools and high-temperature resistant components, etc.

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Three trends in advanced ceramics are following: ➀ Development from single-phase, high-purity to multi-phase composite ceramics, including fiber (or whisker) reinforced ceramic matrix composites, complexphase ceramics reinforced by heterogeneous particle dispersion self-reinforcing materials made of two or more main crystalline phase, gradient functional ceramic and nano-micron ceramic composites. ➁ Development from the micron scale (from powder to microstructure) towards the nano-scale (1 to several hundred nanometers), the transitional structural region between atomic or molecular and conventional micron structures, which results in chemical and physical properties different from those of previous micronscale ceramic, such as super plasticity and changes in electrical and magnetic properties. ➂ Processing of ceramic, such as cutting, shape design and joining (welding). (2) Intermetallic compounds Intermetallics Compounds (IMC) are composed of two or more metal elements in proportion to their long-range ordered crystal structure and metallic properties (metallic luster, electrical and thermal conductivity) that differ from their constituents. It is characterized by the fact that each element has a stoichiometric component and its composition can be varied within a certain range to form a solid solution based on the compound. Metallic elements in IMC are bonded to each other by a mixture of covalent and metallic bonds. Properties of IMC are intermediate between those of ceramics and metals and it is also known as semi-ceramic materials. Plasticity and toughness are lower than those of ordinary metals and higher than those of ceramic. Properties at high-temperature are better than those of ordinary metals and lower than those of ceramic. When two metals form a compound in an integer ratio (or within a certain range of near-integer ratios), the structure is different from that of the two metals that make it up, resulting in a long-range ordered superlattice structure. Intermetallic compounds are divided into two categories: those used as loadbearing structures with good mechanical properties at room and high temperatures, and those used as functional materials with some special physical or chemical properties. Metals lose their strength at high temperatures, while IMC do not have this problem and can be called ‘heroic’ at high temperature. Within a certain temperature range, the strength of intermetallic compounds increases with temperature, which gives them a potential advantage in high-temperature structural applications. When intermetallic compounds were first discovered in the 1930s, they almost have not any ductility at room-temperature. As a result, it was predicted that intermetallic compounds had no practical value in terms of structure. In the mid-1980s, American scientists made a breakthrough in the study of brittleness of intermetallic compounds at room temperature to make its elongation substantially higher and even comparable to the ductility of pure aluminum. This important discovery and the development prospect attracted the attention from all over the world. A work wave of research and development focusing on intermetallic

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compounds has broken through the brittleness of intermetallic compounds such as Ti3 Al, Ni3 Al, TiAl, NiAl, etc., which makes these materials to get a crucial step toward engineering practicality. After the brittleness of intermetallic compounds has been largely improved, a number of issues need to be addressed to make these alloys practical engineering materials, such as further increasing strength at high temperature, improving processability (especially ductility, weldability) and ensuring tissue stability. Alloys or materials based on IMC are a completely new type of material. Conventional metals are based on solid solutions at the beginning or end of the phase diagram, while intermetallic compounds are based on ordered intermetallic compounds in the middle part of the phase diagram. Many intermetallic compounds have paradoxical properties of strength versus temperature. The yield strength of these intermetallic compounds increases with increasing temperature and then decreases after reaching a peak temperature. Intermetallic compounds have unique physicochemical properties, such as unique electricity, magnetism, optics, acoustics, chemical and thermal stability, strength at high temperature. In addition, intermetallic compounds also have good resistance to oxidation and corrosion, superconductivity, semiconductivity and other functional properties. Intermetallic compounds are a class of structural materials at high temperature with great potential for development of advanced materials. There are a wide variety of intermetallic compounds, including all metal-tometal compounds, and they do not follow the traditional laws of chemical valency. Currently, intermetallic compounds used in engineering structures are the three major alloy systems including Ni–Al, Ti–Al and Fe-Al. Ni–Al intermetallic compounds are a class of materials that have been studied further and have many achievements and practical applications. Ni–Al and Ti–Al intermetallic compounds have excellent performance, but its high price make them mainly used in aerospace and other high-tech fields. Intermetallic compounds, known as the largest use of ‘high temperature materials’, is in the aerospace. They are consisted of light metals (such as Ti, Al) with low density, high melting point, high temperature performance, etc. and have an extremely attractive application prospects. (3) Laminated materials Laminated materials (also known as laminated composites) are ‘sandwich’ type structures or multilayer materials (microlaminated materials),which are formed by overlapping two or more materials with different physical and chemical properties at a certain spacing or thickness between layers. The properties of laminated materials depend on the structure and properties of each layer, their volume contents, the mutual solubility and even the intermetallic compounds formed between the two components. Laminated materials are highly valued because they can meet the structural needs for high-performance products. The laminated materials aim to overcome the brittleness of intermetallic compounds by using ductile metals. The interface between layers has an important influence on the internal load transfer, stress distribution, strengthening mechanism

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and fracture process, which makes laminated materials better than the corresponding monolithic materials in terms of performance, with more excellent toughness at high temperature, good resistance to creep, low-temperature fracture strength, microstructural thermo-mechanical stability at high temperature and good prospects for application in the aerospace field. Based on further understanding of the performance characteristics of laminated materials and preparation processes, the analysis of weldability of laminated materials is of great significance to promote their development and application. Ni–Al, Ti–Al and other intermetallic compounds are a class of high-temperature structural materials with great potential for development and broad application prospects in the aerospace because of their excellent properties such as good specific strength, specific stiffness, resistance to oxidation and corrosion. However, the brittleness of intermetallic compounds at room temperature seriously limits its practical applications. The Bauman Technical University in Moscow and GE (funded by the Materials Guidance Division of the United States Air Force Laboratory) company have carried out research and developed laminated composites made of intermetallic compounds and ductile metals, which can overcome the brittleness of intermetallic compounds and offer prospects for the development of aerospace materials. Microlaminated composites are made by adding ductile metal interlayers between brittle intermetallic compound, and their properties depend on the type of the components, volume fraction, layer spacing and thickness ratio. The interlayer interface has an important influence on the internal load transfer, residual stresses, microzone stress and strain distribution, mechanisms of strengthening and fracture. There are three types of strengthening effects of alternating interfaces on microlaminated composites: Orowan strengthening, where the interface acts as an impediment to dislocation motion within the layers; Koehler strengthening, where the resistance to dislocation motion increases due to the formation of forces acting on dislocations due to the difference in modulus between the two sides of the interface; and HallPatch strengthening, where the grain boundaries act as an impediment to dislocation motion. The stress field of laminated composites is an energy dissipative structure that can overcome the fatal weakness of sudden fracture of brittle materials. When the micro-laminated material is impacted or bent, the micro-crack is repeatedly hindered at the layer interface and deflected or blunted, which can effectively relieve the stress concentration at the crack tip and improve the toughness. A well-bonded interface has the effect of preventing crack propagation and relieving stress concentration. (4) Composite materials Composite material is a multi-phase solid material synthesized in a certain forming process, composition proportion or distribution state by two or more substances with different physical and chemical properties. Composite material can be got comprehensively excellent performance that cannot be achieved by a single material by a good reinforcing phase combined with matrix and a proper manufacturing process. The strengths of each component are fully utilized in composite material.

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The composite material maintains the advantages of each component and its relative independence, but it is not a simple superposition of the properties of each component. The development of composite materials can be divided into two stages, early composites and modern composites. The term ‘composite materials’ appeared in the 1940s, when glass fiber reinforced unsaturated polyester resins appeared. A variety of high performance fibers were developed after the 1960s. Various types of materials used as composite matrices (e.g., resin-based, metal-based, ceramicbased, carbon/carbon-based) and reinforcing phases have been used and improved to bring the development of composites to a higher level since 1980s. The stage of development of high-performance modern composites was coming. Manufacturing technology for composite material is essentially the use of the original metal, inorganic non-metallic and polymer materials as components, combining reinforcing phase with matrix together by a certain process to make composite material not only retaining the characteristics of the original material but also presenting some new properties. Composite materials generally have two basic phases: a continuous phase (called matrix); and a dispersed phase (called reinforcing phase). The properties of the composite depend on the properties of these two phases, their proportions, and are also closely related to the characteristics of the interface between the two phases and the geometric features of the reinforcing phase. The dispersed phase is distributed independently throughout the continuous phase, and they are diffusely distributed fillers of fibers, whiskers and particles (denoted by subscripts f, w, and p, respectively). Metal matrix composites include several types of whisker, particle and short fiber reinforced, which are made of elements (e.g., C, B, Si, etc.), oxides (e.g., Al2 O3 , TiO2 , SiO2 , ZrO2 , etc.), carbides (SiC, B4 C, TiC, VC, ZrC, etc.) and nitrides (Si3 N4 , BN, AlN, etc.). Continuous fiber-reinforced metal matrix composites consist of a base metal and reinforcing fibers. The base metal is usually some plastic, ductile metal with good weldability. While the reinforcing phase is a non-metal with high strength, high modulus, high melting point, low density and low linear expansion coefficient, whose weldability is not good. The welding of such materials involves not only the welding of matrix metal, but also the welding of metal with non-metal reinforcing phases and even the welding of different reinforcing phases. (5) Functional materials Materials can be divided into two categories: structural and functional materials. The concept of functional materials was first introduced by J.A. Morton in 1965. Functional materials are materials with specific functions, which play a functional role in objects. Many new functional materials have been mass-produced and applied to promote the further development of modern science and technology. Functional materials are those with excellent electrical, magnetic, optical, thermal, acoustic, mechanical, chemical, biomedical functions. These specially physical, chemical and biological effects can complete the functional interconversion, which

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can be mainly used to manufacture a variety of functional components in various high-tech fields of high-tech materials. The research on functional materials in the world is extremely active with full of opportunities and challenges, and new technologies and patents are emerging all the time. Developed countries are attempting to form a technological monopoly in the field of special functional materials by intellectual property rights to occupy the vast market in China. The strict situation has attracted the great attention of China. Functional materials are not only important basic materials for the development of information, biotechnology, energy, high-tech fields and national defense construction, but also the basis on the transformation and upgrading of basic and traditional industries in China. Functional materials are directly related to the sustainable development of resources, environment and society in China. Functional materials play a unique role in the national economy, social development and national defense construction, involving in information, bioengineering, energy, nano, environmental protection, space and other modern high and new technologies and their industries. Functional materials play an important role not only in promoting and supporting the development of high and new technologies, but also in promoting the transformation and upgrading of related traditional industries to get a high-level development in China. Functional materials have a wide range of applications to form a large-scale and high-tech industry group with broad market prospects and strategic importance. All countries in the world attach great importance to the research, development and application of functional materials, which has become a hot research and development of new materials, and also a hot strategic competition in the development of high technology in the world. Under the certain conditions, structural and functional materials can be transformed into each other. It is difficult to separate structural and functional materials because they share a common scientific basis. Sometimes, a material has both structural and functional properties, for example, an airframe stealth material has three functions: load-bearing, aerodynamics and stealth. At present, some new breakthroughs are coming for international functional materials and their application technologies, such as superconducting, microelectronic, photonic, information, energy conversion and storage, ecological and environmental, biomedical materials, etc. Functional materials are in the midst of rapid development and they are becoming an important means for some developed countries to strengthen their economic and military advantages.

1.2 Applications and Development Prospects of Advanced Materials For modern industry, material is the substance, manufacturing is the way (or means), and application is the purpose. In the application of advanced materials, requirements for the special environment must be considered, such as high or low temperature,

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corrosive medium, etc. Because most structural parts have certain shape fit and accuracy requirements, advanced materials need to have good machinability, such as castability, cold (or hot) formability, weldability and machinability, etc. Unfortunately, welding of advanced materials is very difficult due to their inherent special properties, which sometimes hinders the development and application of advanced materials.

1.2.1 Advanced Ceramics Advanced ceramics have a wide range of applications due to their rich raw materials and high value-added products. However, it is difficult to be processed due to the poor plasticity and toughness, which is not easy to make large or complex-shaped components. The individual use of ceramics is also be restricted. Advanced ceramics are improving with the development of modern electrical, electronic, aviation, atomic energy, metallurgy, machinery, chemical and other industries, as well as computer, space technology, new energy, other science and technology. In practical applications, the joining technology is often used to make ceramic–metal composite components, which can not only give fully respective performance advantages of ceramics and metals, but also reduce production costs. The joining of ceramic with metal has wide application in automotive engine supercharger rotors (which can reduce exhaust emissions), ceramic/steel rockers, ceramic/metal tappets, spark plugs, high-voltage insulators and electronic components (e.g., vacuum tube housings, rectifier housings), etc. The research and development of high efficiency ceramic engines is one of the hot fields of high-tech competition in the world. The use of ceramic engines can increase the operating temperature from 1000 to 1300 °C with an increase of thermal efficiency from 30 to 50%, a reduce of weight by 20% and a saving of fuel by 30–50%. Britain is the first country engaged in the development and applications of structural ceramic. The British government has earmarked tens of millions of pounds to support the research of ceramic gas turbines and reciprocate ceramic engines. And piston ceramic engine has been manufactured. According to the estimation by experts from the United States Ford Motor Company, if ceramic engines are used to all cars in the United States, at least 500 million barrels of oil will be saved each year. For ceramic engines, the United States, Russia, France, Germany and other countries have set research and development plan by investing a huge amount of humanity and funds. The United States government has invested billions of dollars to organize dozens of companies engaged in the research and development of ceramic engines, including General Motors, Ford Motor Company, Knowlton and other large enterprises. They have established some specialized research and development institutions for new ceramic engine. Structural ceramics is known as a new field after microelectronics that can bring huge benefits in Japan. Japanese have spared no expense in competing with the Americans to develop new products even more than the United States. Ceramic

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engine with a power of 213 kW is already in mass production and has equipped millions of cars in Japan. The research and development of ceramic combustion engine in German is also at the forefront of the world. The ‘2000 car’ developed by Mercedes-Benz company is driven by the ceramic gas turbine. In the ‘Eureka Project’ of European Community, three countries, including France, Germany and Sweden, have united to develop ceramic gas turbines since the 1980s and have developed a ceramic turbojet engine of 147 kW with an operating temperature of 1600 °C, which is more than 600 °C higher than ordinary engines.

1.2.2 Intermetallic Compounds In the last two decades, the development and application of intermetallic compounds has been concerned, which is a fundamental change in the field of materials and one of the important tendency for future development. Intermetallic compounds have some properties that other solid solutions do not have due to their special crystal structure. In particular, the strength of solid solution usually decreases with the increase of temperature. But the strength of certain intermetallic compounds increases with temperature within a certain range, which makes it possible to serve as the basis for new high-temperature structural materials. In addition, intermetallic compounds have some properties that are several times or even tens of times greater than those of solid solution. Ni–Al and Ti–Al intermetallic compounds are suitable for aerospace structure and have good application potential, which have been valued by Europe, the United States and other developed countries. Some Ni–Al alloys have been applied or tried for diesel engine, electrical components and aerospace fasteners, etc. Ti–Al alloys can replace nickel-based alloys to make aerospace engine high-pressure turbine stator bearing rings, high-pressure compressor box and engine combustion chamber expansion nozzle, etc. Engine hot-end components is being tried to manufactured by aerospace industry in China, which bring it to a broad application prospects. For example, low-pressure blades of TiAl alloy (Ti-47Al-2Cr-2Nb) were performed on the CF6-80C2 fighter aircraft and 1,000 simulated flight cycles were carried out to prove blades well done by GE Engines Company in the 1990s. Subsequently, TiAl alloy has been used as the grade 5 or 6 low-pressure engine blade for GE-90 engine aiming at replacing the original Rene77 blade to reduce the weight by 80 kg in ‘AITP’ program signed by NASA. At the same time, the application of Ti–Al alloy is also being improved as case, turbine disk, support frame and guide beam, etc. Fe3 Al intermetallic compound can replace many materials including stainless steel, heat resistant steel or high temperature alloy in some occasions due to its good resistance to oxidation and wear. Especially, it is suitable for applications under severe conditions (such as high temperature or corrosive environment) because of its good resistance to sulfation. For example, it can be used for structural parts of thermal

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power plants or carburizing furnace, chemical devices, automobile gas exhaust, petrochemical catalytic cracking device, guide in heating furnace, high temperature grates, etc. In addition, Fe3 Al intermetallic compound can be developed into new electric heating materials due to its excellent oxidation resistance at high temperature and high resistivity. Fe3 Al also can be made into composite structures with WC, TiC, TiB, ZrB and other ceramic materials, which has a broader application prospect.

1.2.3 Laminated Materials The current research of laminated composites is mainly focused on the manufacture process, interfacial properties, strengthening mechanisms, etc., and there is less research on their applications for welding. The University of California used Ag–CuIn filler metal to braze Ti-6Al-4 V/TiAl3 micro-laminated composites in vacuum and the tensile strength of brazed butt joints was only 20 MPa, while the tensile strength of Ti–Al micro-laminated composites was 200 MPa. This required further improvement of process parameters and the search for more reliable welding technology. The overall performance of the laminated material of Ti3 Al base covered by pure nickel layer is good at high temperature, and the main problem of its welding is the lack of room temperature toughness to microcrack at the bond interface. The tendency to microcrack can be reduced by controlling the heat input and using a suitable preheating before welding. The crystalline layer in the weld metal will disappear and the microhardness distribution will be consistent through the weld metal after preheating. However, the rapid cooling of the welding is a non-equilibrium process and the ordering process does not proceed sufficiently, which has an impact on the microstructure properties of the weld metal, which must be considered in the welding of metal/intermetallic compound laminated composites. The microstructure, melting temperature, thermal expansion coefficient, thermal conductivity and other physical and chemical properties in toughness layer and intermetallic compound layer are different due to special laminated structure, which causes the welding of laminated composites more complex and difficult than the individual block material. Interfacial reaction occurs by the effect of the welding thermal cycle on the interface of laminated composite material. And even some potential defects may be transformed into pores or cracks at the interface affected by the welding metallurgical process. Operating temperature in the turbine gas channel of aerospace vehicle engine is increasing (generally above 1100 °C) with the high requirements for thrust-toweight ratio and fuel efficiency, which makes engine blades better resistance to high temperature and good damage fracture toughness. The toughness above 1000 °C of traditional nickel-based alloy declines quickly, which is easy to be oxidized and is difficult to meet the requirements. The use of Ni, Ti, Nb, V and other hightemperature metals and their intermetallic compounds (such as Ni–Al, Ti–Al, Nb-Al, Nb-Ti–Al) as raw materials for the preparation of laminated composites and as toughening elements can overcome the brittleness of intermetallic compounds. Laminated

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1 Overview

composites have more excellent toughness at high temperature, better resistance to creep, fracture toughness at low temperature, resistance to oxidation in the process of thermal cycling, thermo mechanical stability of microstructure at high temperatures and competitive cost. The microlaminated composites made by alternating layers of high temperature metal foils (such as Ti, Ni, V) and Al foils to form intermetallic compounds by rolling or self-propagating synthesis at high temperature may have incomplete Al layers, which limits their application under conditions of high temperature. However, these microlaminated composites can be used as lightweight structural materials and has a promising application in the manufacture of airframe structures due to its good specific strength and stiffness, in which Al can change the brittleness of the intermetallic compound as a toughening element. The laminated materials have good prospects for aerospace engine manufacturing because of their good properties at high temperature and thermo mechanical stability. Although microlaminated composites prepared by vacuum rolling or self-propagating synthesis at high temperature limits their application in high temperature, they can be used as lightweight structural materials for airframes. Welding is an important technology to achieve the connection of multiple aerospace components, but there is still a lack of systematic research and development on welding applications for laminated materials. Key issues should be considered in the welding of laminated composites, such as welding heat input that may cause the expansion of potential defects at the interface and the different thermal expansion coefficients between the composite and base layers that are prone to cracking.

1.2.4 Composite Materials Composite materials were created in the early 1960s in response to the needs for development of aerospace and aviation. Composites are designable by the choice of different matrix and reinforcing phases, combination of materials, the ratio and distribution of reinforcing phases and even requirement for industry. The application advantage of composite materials lies in the combination of different materials to form a variety of new materials with excellent performance, and the integration of structure and function is the development trend of composite materials. Over the past 30 years, the application of composites in fighter aircraft has continued to grow, replacing a significant proportion of conventional structural materials. Composites have shown significant effect of weight reduction. For example, weight reduction effect is obvious due to the thin metal structure subject to small loads (such as the front fuselage) and it is not obvious due to the complex lay-up (such as at the wing root) for structures subject to large loads. However, the weight reduction effect is in the middle for most of the structure in the aircraft. It is generally said that when composite materials account for 20 to 25% of the weight in the structure, there is a substantial increase in the weight reduction effect on the aircraft fuselage.

1.2 Applications and Development Prospects of Advanced Materials

13

The use of composite materials in civil aircraft and helicopters is also gradually increasing. The application of composites also plays a key role in structures including artificial earth satellites, space warfare, transportation systems from heaven and earth, launch vehicle arrow bodies and strategic missile warhead. For example, carbon/carbon composites are used in almost all the end caps of long-range and intercontinental strategic missiles developed by many countries. Carbon/carbon composites, particularly suitable for long-range missiles and ground-returning satellite forward head caps, have the following advantage. ➀ Resistance to high temperature and low density, used for intercontinental missiles, carbon/carbon composites can increase the range by 300 km for weight reduction of each 1 kg; used for spacecraft and space shuttles, carbon/carbon composites can reduce the thrust by 2 kN for weight reduction of each 1 kg and greatly save rocket fuel. ➁ Carbon fiber composites are slow to ablate under the impact of ultra-high temperature and high airflow. They can form a layer of strong and loose ‘sponge body’, which can prevent further ablation and can also play a role in heat insulation after sintering. The first application of large composite structural parts on launch vehicles is the satellite joint bracket of the ‘Long March 2’ in China. It uses a carbon/epoxy composite semi-rigid shell with a ribbed aluminum honeycomb core structure. ‘Common bottom’ is a large aluminum skin fiberglass honeycomb sandwich glued vacuum insulation structural parts, which is the key component of ‘Long March 3’series of launch vehicles using advanced composite molding process to achieve the advanced design and manufacture of large launch vehicles cryogenic propellant tank structure. Their manufacturing has played a key role to improve the launch capability of the rocket. Continuous fiber-reinforced metal matrix composites are limited to a few fields including aerospace and military industry due to their complex manufacturing process and high cost. Discontinuous reinforced metal matrix composites maintain most of the excellent properties of continuous fiber reinforced MCM, whose manufacturing process is simple with low cost. In addition, it is easy for discontinuous reinforced metal matrix composites to be processed second time, resulting in the rapid extremely development in recent years. Although the weldability of these materials is better than that of continuous fiber-reinforced metal matrix composites, it is still very difficult to be welded compared with single metals and alloys. Discontinuous reinforced metal matrix composites are mainly SiCp /Al, SiCw /Al, Al2 O3p /Al, Al2 O3f /Al and B4 Cp /Al, etc. The application of fiber-reinforced metal matrix composites is expanding.

14

1 Overview

1.2.5 Functional Materials The development of functional materials has attracted more interesting in China including many research plans, such as the national scientific and technological ‘863’ or ‘973’ Program, the National Natural Science Foundation and other plans. Special functional materials were called as ‘national defense cutting-edge’ materials in the ‘Tenth Five-Year Plan’ and ‘Eleventh Five-Year Plan’ from the national defense science and technology plan. The implementation of these research plans has made China achieve fruitful results in the field of functional materials. With the support of the ‘863 Program’, new energy materials have been developed such as superconducting materials, flat panel display materials, rare earth functional materials, bio-medical materials and hydrogen storage materials. A number of research results close to or at the international advanced level have been shown in the new fields of functional materials including diamond films and infrared stealth materials. Even some researches have occupied a place in the international arena. Functional materials have also made significant contributions to national defense projects such as the ‘two bombs and one star’ and the ‘four major equipment and four stars’. Functional materials have been developed rapidly to be tens of thousands of varieties in several ten categories in recent years. The application of functional materials has also expanded rapidly, and they are widely used in the fields of electronic information, computers, optoelectronics, aerospace, weaponry, energy, medicine and so on. Although functional materials are not as good as structural materials in terms of output and production value, they have a great influence on the development of various industries, especially in high technology, which sometimes plays a key role. For example, superconducting materials represented by NbTi, Nb3 Sn have been commercialized and have gained applications in several fields such as nuclear magnetic resonance body imaging (NMRI), superconducting magnets and large accelerator magnets. Further applications of low-temperature superconducting materials are limited by the fact that the critical temperature of conventional lowtemperature superconductors is too low and must be used in expensive and complex liquid helium (4.2 K) systems. The development of oxide superconductors at high temperature has broken the temperature barrier and raised the operating temperature from liquid helium (4.2 K) to the liquid nitrogen (77 K). Compared with liquid helium, liquid nitrogen is a very economical refrigerant and has a high heat capacity, which brings great convenience to engineering applications. Oxide superconductors at high temperature are complex multifaceted systems that involve several fields in the research process including condensed matter physics, crystal chemistry, processing technology and microstructure analysis. Some of the latest technologies in the field of materials science research, such as amorphous, nanotechnology, magneto-optical, tunneling microscopy and field ion microscopy, have been used to study these superconductors. Many of these research efforts involve the frontiers of materials science. Important progress has been made in the high temperature superconducting materials including single crystals, thin films, bulk materials, wires and their applications.

Bibliography

15

Shape memory alloy (SMA) is a new type of functional material, which has a special shape memory effect and has a broad application prospect in aerospace, atomic energy, marine development, instrumentation, medical devices and other fields. It is difficult to realize the joining of TiNi shape memory alloy by traditional welding method, and it is also not easy to control the chemical composition, microstructure and phase transition temperature during welding. The function of joint isn’t consistent with the base material to obtain the shape memory effect equivalent to that of the base material. Solid joining methods are promising, transient liquid phase diffusion welding and brazing are also beneficial for welding shape memory alloys. The development of advanced materials has become an important part of strategies for science and technology in some developed countries and regions such as the United States, Europe and Japan. When plans are formulated in national science and technology and industrial development, processing technology of advanced materials is listed as one of the key to be developed on a priority basis in order to maintain their economic and technological leadership. Significant progress has also been made in the development and industrialization of advanced materials in China, which provides strong support for economic and social development. The development of advanced materials has promoted the progress in science and technology and the changes in industrial structure. The research and industrialization application of structural materials with high performance got in some machinery, equipment, large-scale, efficient, high-parameter, multi-functional material basis also makes the rapid development of welding technology in advanced materials, which will promote the continuous progress and forward development of society.

Bibliography 1. Shi Y (2006) Chinese materials engineering dictionary: Materials welding engineering, vol 23. Chemical Industry Press, Beijing 2. Technology Forecasting and National Key Technology Selection Study Group (2004) Technology forecast report in China: information, biology andm aterials. Science and Technology Literature Press, Beijing 3. Zhong Z-L, Ye H-Q (1992) Intermetallic compounds. In: Proceedings of the first national symposium on high temperature structured intermetallic compounds. Machinery Industry Press, Beijing 4. Ren J, Wu A (2000) Joining of advanced materials. Machinery Industry Press, Beijing 5. Yajiang L, Juan W, Yansheng Y et al (2002) Phase constitution near the interface zone of diffusion bonding for Fe3 Al/Q235 dissimilar materials. Scripta Mater 47(12):851–856 6. Shen Z, Qiu ZW (1992) Principles of composite materials and their applications. Science Press, Beijing 7. Wang F, Chen ZL (2002) Application of shape memory alloy materials. Mater Mech Eng 26(3):5– 8

Chapter 2

Welding of Advanced Ceramic Materials

High-tech ceramics are in rapid development and have become key engineering materials. On the whole, ceramics are hard and brittle high melting point materials with low thermal conductivity, good chemical and thermal stability, as well as high compressive strength and unique properties such as insulation and electrical, magnetic, acoustic, optical, thermal and biocompatibility. They can be applied in various fields such as mechanics, electronics, astronautics, medicine and energy, becoming an important part of modern high-technology materials. The welding applications of advanced ceramic materials have also attracted increasing attention.

2.1 Performance Characteristics and Joining Problems of Ceramic Materials Ceramics are inorganic non-metallic materials synthesized from oxides of various metals, nitrides, carbides, and silicides by proper batching, forming, and high temperature sintering. Ceramics have many unique properties, and such materials are generally bonded by covalent, ionic or mixed bonds with strong bonding forces. Thus they have high modulus of elasticity and hardness. Ceramic materials can be divided into two categories according to their application characteristics: functional ceramics and engineering structural ceramics. Functional ceramics are ceramic materials with electric, magnetic, optical, acoustic, thermal and other functions as well as coupling functions. In the view of performance, there are ferroelectric, piezoelectric, photoelectric, acousto-optical, magneto-optical, biological and other functional ceramics. Engineering structural ceramics emphasize on the mechanical properties of the material and are widely used in the engineering field for their high temperature resistance, high strength, super hardness, high insulation, high wear resistance, corrosion resistance and other properties. Common engineering structural ceramics are shown in Table 2.1. © Chemical Industry Press 2023 Y. Li, Joining Technology and Application of Advanced Materials, Advanced and Intelligent Manufacturing in China, https://doi.org/10.1007/978-981-19-9689-4_2

17

18

2 Welding of Advanced Ceramic Materials

Table 2.1 Common engineering structural ceramics Category

Composition material

Oxide ceramics

Al2 O3 , MgO, ZrO2 , SiO2 , UO2 , BeO etc.

Non-oxide ceramics

Carbide

SiC, TiC, B4 C, WC, UC, ZrC etc.

Nitride

Si3 N4 , AlN, BN, TiN, ZrN etc.

Boride

ZrB2 , WB, TiB2 , LaB6 etc.

Silicide

MoSi2 etc.

Fluoride

CaF2 , BaF2 , MgF2 etc.

Sulfide

ZnS, TiS2 , Mx Mo6 S8 (M=Pb, Cu, Cd) etc.

Carbon and Graphite

C

2.1.1 Performance Characteristics of Structural Ceramics 2.1.1.1

Physical and Chemical Properties

The physical properties of ceramic materials are quite different from those of metal materials, mainly in the following aspects: the linear expansion coefficient of ceramics is lower than that of metals, generally in the range of 10–5 ~ 10–6 K−1 ; the melting point (or sublimation, decomposition temperature) of ceramic is much higher than that of metals, some ceramics can work at a high temperatures of 2000 ~ 3000 °C and maintain the strength at room temperature, while most metals will basically lose strength above 1000 °C. Some new special ceramics have electrical conductivity under specific conditions, such as conductive ceramics, semiconductor ceramics, piezoelectric ceramics, etc. There are also some ceramics with special optical properties, such as transparent ceramics, optical fibers, etc., but they are mainly functional ceramics rather than structural ceramics. The microstructure of ceramics is very stable and they have good chemical properties. In its ionic crystal, the metal atoms are surrounded by non-metal (oxygen) atoms and shielded by the non-metal atoms, thus forming an extremely stable chemical structure. Generally, It will no longer interact with the oxygen in the medium and does not oxidize even at high temperatures of 1000 °C. Due to the stable chemical structure, most ceramics have a strong resistance to corrosion by acids, bases, salts, and as well as molten metals.

2.1.1.2

Mechanical Properties

Ceramic materials are mostly crystals composed of ionic bonds (such as Al2 O3 ) or covalent crystals composed of covalent bonds (such as Si3 N4 , SiC), and such crystal structures have obvious directionality. Polycrystalline ceramics have few slip systems, which can hardly produce plastic deformation when subjected to external forces, often undergo brittle fracture, and have poor impact resistance. Due to the ionic

2.1 Performance Characteristics and Joining Problems of Ceramic Materials

19

crystal structure, the hardness and room temperature elastic modulus of ceramics are also high. There are a large number of pores in ceramics, and the density is much lower than that of metals, so their tensile strength is not high, but the compressive strength of ceramics is still relatively high because the pores do not lead to crack expansion under pressure. The ratio of tensile strength to compressive strength of brittle materials cast iron is generally 1/3, while ceramic is about 1/10. Ceramics are very strong ionic/covalent bonded (stronger than metal bonds) and this bonding gives ceramics the associated properties: high hardness, high compressive strength, low thermal conductivity, electrical conductivity and chemical inactivity. This strong bond also exhibits some undesirable properties, such as low elongation. The inherent high hardness of ceramics can be overcome by controlling the microstructure and producing ceramic springs. Composite ceramics have been developed for applications with fracture toughness up to half that of steel. The broader properties of ceramics may have not been recognized. Ceramics are generally thought of as electrical/thermal insulators, while ceramic oxides (based on Y–Ba–Cu–O) have high temperature superconductivity. Diamond, beryllium oxide, and silicon carbide have higher thermal conductivity than aluminum or copper.

2.1.1.3

Several Commonly Used Structural Ceramics

(1) Oxide ceramics Commonly used oxide ceramics include alumina ceramics, beryllium oxide ceramics, and partially stabilized zirconia ceramics. The physical properties of several commonly used oxide ceramics are shown in Table 2.2. (1) Alumina ceramics Alumina ceramics are widely used ceramic materials in engineering. The main components of alumina ceramics are Al2 O3 and SiO2 . The higher the Al2 O3 content, the better the performance, but the process is more complex and the cost is higher. The chemical composition of several oxide ceramics is shown in Table 2.3. Oxide ceramics include not only single component ceramics with a small amount of glass phase or other crystalline phases (such as alumina, beryllium oxide, etc.), but also many multi-component ceramics with natural minerals as raw materials (such as forsterite, etc.). There are more than ten isomers of alumina, and the three main common ones are: α-Al2 O3 , β-Al2 O3 and γ-Al2 O3 . γ-Al2 O3 is a spinel-type cubic structure, which is unstable at high temperatures. It will transform to α-Al2 O3 at 1600 °C. α-Al2 O3 is very stable at high temperatures and has no crystalline transformation until it reaches its melting point of 2050 °C. The main performance characteristics of alumina ceramics is high hardness (87HRA at 760 °C, 82HRA is still maintained at 1200 °C), good wear resistance, corrosion resistance, high temperature resistance, can be used for a long time at the

3.5

9.5

8.5 25–30

Relative dielectric constant/1 MHz

Dielectric strength /kV·mm−1 15–18

>1013

>1013

Specific resistance/l·cm

6.7 7.7

280–350

2000

304

0.218 0.126

25°C 300°C

Thermal conductivity/W·(cm·K)−1

95%Al2 O3 –

– –

250–300 6.6 7.6

25 ~ 300°C 25 ~ 700°C

Compressive strength /MPa

Bending strength /MPa

1200

Elasticity modulus /GPa

Coefficient of linear expansion /10–6 ·K−1

3.2 ~ 3.4 304

Density/g·cm−3

–

75%Al2 O3

Alumina

Melting point (decomposition point) /°C

Materials

Table 2.2 The physical properties of several oxide ceramics 99%Al2 O3

25–30

9.35

>1014

0.314 0.159

6.8 8.0

370–450

2500

382

3.9

2025

15

6.5

>1014

1.592 0.838

6.8 8.4

172

1472

294

2.8

2570

Beryllia (BeO)

–

–

>1014

0.0195 0.0205

14

8.9

>1014

0.419 –

140 ≥10 –

650

850

345

3.56

2800

Magnesia (MgO)

≥10 –

2060

205

3.5

2550

Zirconia (ZrO2)

13

6.0

>1014

0.034 –

10 12

137

579

–

2.8

1885

Forsterite (2MgO·SiO2 )

20 2 Welding of Advanced Ceramic Materials

2.1 Performance Characteristics and Joining Problems of Ceramic Materials

21

Table 2.3 The chemical composition of several oxide ceramics % Materials

75% Alumina ceramics

95% Alumina ceramics

99% Alumina ceramics

Steatite ceramics

Forsterite ceramics

SiO2

15.30

2.50

0.30

55.80

44.50

Al2 O3

75.80

94.70

99.10

3.90

5.10

TiO2

0.25

Trtace

Trtace

Trtace

0.10

Fe2 O3

0.40

0.10

0.14

0.45

0.20

CaO

2.30

2.50

Trtace

0.05

Trtace

MgO

1.85

Trtace

0.25

28.90

49.70

R2 O

0.60

0.20

0.20

0.07

0.20

BaO

3.20

–

–

6.60

–

ZrO2

Trtace

–

–

3.80

Trtace

high temperature of 1600 °C. Alumina ceramics also have good electrical insulation properties, electrical insulation performance in high frequency is particularly outstanding. It can withstand more than 8000v per millimeter thickness. The disadvantage of alumina ceramics is low toughness, poor resistance to thermal vibration, can not withstand rapid changes in temperature. These ceramics are mainly used in the manufacture of cutting tools, molds, bearings, crucibles for melting metals, hightemperature thermocouple sleeves, and some special parts in the chemical industry, such as sealing slip rings, bushings and impellers for chemical pumps. (2) Partially stabilized zirconium oxide (ZrO2 ) ceramics Zirconia ceramics have three crystalline forms: tetragonal structure (t-phase), cubic structure (c-phase) and monoclinic structure (m-phase). With the addition of a suitable amount of stabilizer, the tetragonal structure (t-phase) exists in a metastable state at room temperature, which is called partially stabilized zirconia (PSZ for short). Partially stabilized zirconia ceramics can be applied to structural parts of engines, and its bending strength can be up to 981 MPa at 600 °C. The martensitic transformation of the tetragonal structure (t-phase) to the monoclinic structure (m-phase) under stress is called “stress-induced phase transformation”, which absorbs energy during the phase transformation, relaxes the stress field at the crack tip in the ceramic, increases the resistance to crack propagation, and realizes the toughening of zirconia ceramic. The fracture toughness of partially stabilized zirconia ceramics is much higher than that of other structural ceramics. The commonly used stabilizers in several zirconia ceramics developed so far include MgO, Y2 O3 , CaO, CeO2 , etc. ➀ High strength zirconia ceramics (MG-PSZ) Bending strength reaches 800 MPa and fracture toughness reaches 10 MPa·m1/2 . Vibration resistant MG-PSZ has a bending strength of 600 MPa and fracture toughness of 8 to 15 MPa·m1/2 .

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2 Welding of Advanced Ceramic Materials

➁ Tetragonal polycrystalline zirconia ceramics (Y-TZP) With Y2 O3 as stabilizer, their bending strength can reach 800 MPa, and the maximum can reach 1200Mpa, fracture toughness can reach10 MPa·m1/2 or more. ➂ Tetragonal polycrystalline ZrO2 -Al2 O3 composite ceramics The high elastic modulus of Al2 O3 can refine the grains of polycrystalline zirconia ceramics, improve the hardness, increase the t-phase content of tetragonal structure, and thus improve the strength and toughness of ceramics. The bending strength of ZrO2 -Al2 O3 composite ceramics fabricated by hot press sintering method can reach up to 2400 MPa and fracture toughness up to 17 MPa·m1/2 . (2) Non-oxide ceramics This includes silicon nitride (Si3 N4 ), silicon carbide (SiC), boron nitride (BN) and titanium nitride (TiN). Boron carbide (B4 C) is second only to diamond and cubic boron nitride in terms of hardness among engineering materials and is used for components requiring high wear resistance. Because of their high strength, super hardness, wear resistance and corrosion resistance even at high temperatures, nonoxide ceramics have become key materials in high-tech fields such as machinery manufacturing, metallurgy and aerospace. The physical and mechanical properties of several non-oxide ceramics are shown in Table 2.4. ➀ Silicon nitride ceramics They are hexagonal crystalline systems with Si3 N4 as structural unit, and have extremely strong covalent bonding, with two kinds of crystal structure: α-Si3 N4 and β-Si3 N4 . Silicon nitride ceramics are characterized by high strength. The bending strength of reaction sintered silicon nitride ceramics at room temperature reaches 200 MPa, and the strength can be guaranteed without decay at high temperatures from 1200 to 1350 °C. The bending strength of hot pressed sintered silicon nitride ceramics can be as high as 800 to 1000 MPa at room temperature, and the bending strength can reach 1500 MPa after adding certain additives. The hardness of silicon nitride ceramics is very high, second only to diamond, cubic boron nitride and boron carbide. Engines made of silicon nitride ceramics can work at higher temperatures, allowing the engine’s fuel to be fully burned. So that the thermal efficiency can be improved, and the energy consumption and environmental pollution can be reduced. ➁ Silicon carbide ceramics With high thermal conductivity, high corrosion resistance and high hardness, silicon carbide is a covalently bonded compound with high bonding energy and has the structural type of diamond. The common silicon carbide crystalline types are cubic structure β-SiC, which is stable below 2100 °C, and hexagonal structure α-SiC, which is stable above 2100 °C. Under the pressure of 101.33 MPa, silicon carbide decomposes at about 2830 °C. Silicon carbide ceramics are characterized by high strength at high temperatures, with bending strength remaining at a high level of 500 to 600 MPa at 1400 °C. Silicon carbide ceramics have excellent resistance to wear, corrosion and creep. Due to the high temperature strength of silicon carbide ceramics, they can be used in the manufacture of nozzles for rocket tail pipe, throat nozzles for pouring molten metal,

>1013

9.4 ~ 9.5

Specific resistance /l·cm

Dielectric constant

Thermal 0.30 conductivity/W·cm−1 ·K−1

3

Coefficient of linear expansion/× 10–6 ·K−1

9.4 ~ 9.5

>1013

0.14

2.7

160 ~ 180

20 ~ 100

320

65

Elastic modulus /GPa

Bending strength /MPa

2.2 ~ 2.6

45

10 ~ 103

0.81

4.6 ~ 4.8

78 ~ 90

450

93

3.2

45

10 ~ 103

0.43

4

45

405

90–92

3.09

3.4 ~ 5.3

>1014

–

7.5

–

–

2 (Mohs)

2.27

3.4 ~ 5.3

>1014

–

–

–

–

4.8 (Mohs)

–

8.8

>1014

0.7 ~ 2.7

4.5 ~ 5.7

40 ~ 50

279

1400 (HV)

3.32

–

>1012

–

–

70 ~ 80

290

92 ~ 93

3.18

80 ~ 85

3 ~ 3.2

–

>1012

–

–

97 ~ 116

31.5

95

3.29

–

Pressureless Hot sintering pressed sintering

91 ~ 93

–

Sialon

Density /g·cm−3

Cubic

Aluminium nitride (AlN)

Hardness /HRA

Hexagonal

Boron nitride (BN)

1900 1900 2600 2600 3000 3000 2450 – (sublimation) (sublimation) (decomposition) (decomposition) (decomposition) (decomposition) (decomposition)

Pressureless sintering

Hot pressed sintering

Hot pressed sintering

Reactive sintering

Silicon carbide (SiC)

Silicon nitride (Si3 N4 )

Melting point (decomposition point) /°C

Properties

Table 2.4 Physical and mechanical properties of several non-oxide ceramics

2.1 Performance Characteristics and Joining Problems of Ceramic Materials 23

24

2 Welding of Advanced Ceramic Materials

thermowells, heating furnace tubes and gas turbine blades, bearings, etc. They can also be used in heat exchangers, refractory materials, etc. ➂ Sialon ceramics Ceramics composed of Si3 N4 and Al2 O3 are called Theron ceramics. Their forming and sintering properties are better than those of pure Si3 N4 ceramics, and the physical properties are similar to those of β-Si3 N4 and the chemical properties are close to those of Al2 O3 . These ceramics can be formed by hot extrusion, moulding, casting and other techniques, and can be sintered at 1600 °C in an atmospheric pressure-free atmosphere to achieve the properties of hot-pressed silicon nitride ceramics. It has the highest strength among ceramic materials sintered at atmospheric pressure. In recent years, Sailong ceramics have been developed rapidly. (3) Ceramic composites One way to improve the performance of ceramic materials is to make ceramic matrix composites. The toughness and thermal shock resistance of oxide ceramics can be improved by adding other compounds or metal elements to form composite Al2 O3 ceramics.The mechanical properties of several alumina composite ceramics and hotpressed alumina ceramics are shown in Table 2.5. The bending strength of composite ceramics can be significantly improved because the dispersed second phase prevents the growth of Al2 O3 grains and also hinders microcrack propagation. The bending strength of Al2 O3 complex-phase ceramics containing 5% (by volume) SiC can reach more than 1000 MPa, and the fracture toughness is increased to 4.7 MPa·m1/2 . Ceramics can be used as reinforcing agents for composite systems (e.g. GRP) and metal matrix composites (e.g. alumina-reinforced Al/Al2 O3 ), i.e. ceramic fibers, whiskers or particles are mixed into the ceramic matrix materials. This allows the matrix and the incorporated material to maintain their inherent properties, while the Table 2.5 Mechanical properties of hot pressed Al2 O3 ceramics and their complex phase ceramics Main performance

Hot pressed sintering Al2 O3

Hot pressed sintering Al2 O3 + metal

Hot pressed sintering Al2 O3 + TiC

Hot pressed sintering Al2 O3 + ZrO2

Hot pressed sintering Al2 O3 + SiC(w)

Density /g·cm3 3.4 ~ 3.99

5.0

4.6

4.5

3.75

Melting point /°C

2050

–

–

–

–

Bending strength /MPa

280 ~ 420

900

800

850

900

Hardness /HRA

91

91

94

93

94.5

Thermal conductivity /W (cm·K)−1

0.04 ~ 0.045

0.33

0.17

0.21

0.33

Average grain size /μm

3.0

3.0

1.5

1.5

3.0

2.1 Performance Characteristics and Joining Problems of Ceramic Materials

25

combined properties of the ceramic composite far exceed those of the single material itself. Ceramic composites are divided into two main categories: fiber-reinforced and whisker or particle-reinforced composites. ➀ Fiber-reinforced ceramic composites Fibers are continuous or nearly continuous filaments which can enhance toughness and resistance to high temperatures while maintaining or improving strength. Materials that can be made into fibers include Al2 O3 , SiC, Si3 N4 , etc. However, it is difficult to process the ceramic matrix after adding fibers, and many ceramic composites reinforced by fibers fail to achieve the purpose of improving performance because of uneven fiber distribution, decreased fiber properties after processing (welding), or insufficient matrix compactness. ➁ Whisker or particle-reinforced ceramic composites Whiskers are short singlecrystal fibers, either in the form of rods or needles, with an aspect ratio of about 100 and a diameter of less than 3 μm. Al2 O3 ceramic composites reinforced with SiC whiskers have attracted extensive attention. Adding SiC whiskers into a single Al2 O3 ceramic or a multi- component matrix results in a material with much higher strength and fracture toughness, as well as excellent thermal shock resistance, wear resistance and oxidation resistance. Al2 O3 series ceramic composites toughened with ZrO2 are made with diffusely distributed partially stabilized ZrO2 particles to improve the strength and toughness of Al2 O3 ceramic matrices. Ceramics are widely used in engineering and technology because of their good dielectric properties, heat resistance, vacuum denseness, and corrosion resistance. It has long term thermal stability, resistance to corrosion of various media, strict electrical insulation properties and magnetic insulation properties, and has a broad application prospect.

2.1.1.4

Preparation Method of Composite Ceramics

A variety of methods can be used to prepare composite ceramics. The preparation process of composite ceramics is as follows: ingredients design → powder mixing → pressing and forming → sintering. Taking Al2 O3 –TiC composite ceramics as an example, during the sintering process, it is difficult to sinter because Al2 O3 and TiC will react and the reaction gas is generated. Sintering additives, surface treatments or hot-pressing sintering (HP) or hot-isostatic-pressing sintering (HIP) processes are generally required. Sintering is the key process to obtain the desired microstructure and endow the material various properties. Al2 O3 –TiC composite ceramics can be classified according to the sintering method. (1) Pressureless sintering (pressureless sintering, PS). Pressureless (atmospheric pressure) sintering refers to the sintering process in which the sintered blanks are sintered without applied pressure and only under atmospheric

26

2 Welding of Advanced Ceramic Materials

Table 2.6 Effect of sintering aids on the properties of Al2 O3 –TiC-based composite ceramics Sintering aids

Methods

Relative density /%

Bending strength /MPa

Vickers hardness /GPa

Fracture toughness / MPa·m1/2

TiH2

PS

94.9

386 ~ 574

18.1 ~ 19.3

4.2 ~ 4.6

MgO

PS (Powder Embedded Sintering)

96.7

504 ~ 746

–

3.7 ~ 4.7

CaO

PS (Powder Embedded Sintering)

>97.0

–

–

–

Y2 O3

PS (Powder Embedded Sintering)

97.0

–

–

4.6

pressure. Since Al2 O3 and TiC will react to produce gas at high temperature, it is difficult to densify (relative density < 94%) with conventional sintering methods. In order to promote sintering, various sintering aids such as TiH2 , MgO, CaO, Y2 O3 , Cr2 O3 are often added to the Al2 O3 –TiC system, and methods such as rapid temperature rise and powder burial are adopted to make the relative density of the sintered blanks reach 98%. This sintering method can form a liquid phase which is conducive to densification during the sintering process, prevent abnormal grain growth and make the material microstructure uniform. Pressureless sintering allows continuous operation, has low production cost, and unlimited product shape and size. The effect of sintering additives on the properties of Al2 O3 –TiC composite ceramics is shown in Table 2.6. (2) Hot-pressing sintering (hot-pressing sintering, HP) Hot pressing sintering is to apply a large enough pressure to promote sintering while heating and sintering. Due to the simultaneous heating and pressure, it is conducive to the contact, diffusion and flow of powder particles and other mass transfer processes, which can reduce the sintering temperature, shorten the sintering time and inhibit the growth of grains; Without adding sintering aids, it is easy to obtain sintered bodies with theoretical density and porosity close to zero.Since the driving force of hot-pressure sintering on the powder is 20 to 100 times greater than that of atmospheric pressure sintering, the sintering temperature of materials that can be sintered by atmospheric pressure sintering can be reduced by 100 to 150°C if hot-pressure sintering is used. However, the shape and size of the material is restricted during hot-press sintering. It cannot be mass-produced, and the cost is also higher. The effect of hot pressing process on the properties of Al2 O3 –TiC composite ceramics is shown in Table 2.7.

1700

1750

1750

A12 O3 –30TiC

A12 O3 –25TiC

1650

A12 O3 –30TiC

A12 O3 –25TiC

1600

1620

A12 O3 –29TiC

A12 O3 –30TiC

25

30

–

40

25

20

25

30

1700

1750

A12 O3

Pressure/MPa

Hot pressing process

Temperature/ºC

A12 O3 –30TiC

Composition(weight percentage) /%

20

–

30

30

60

60

30

60

Time/min

Table 2.7 Properties of hot-pressed Al2 O3 –TiC-based composite ceramics

100

–

99.5

99.9

99.0

–

99.5

99.5

Relative density/%

450

762

704 ~ 866

637 ~ 805

516

500

401 ~ 471

Bending strength/MPa

21.0

–

19.8 ~ 21.6

–

20.9

20.9

–

18.1 ~ 19.3

Vickers hardness/GPa

5.7

5.7

4.0 ~ 4.6

4.0 ~ 4.6

5.2

5.2

5.1

3.0 ~ 3.4

Fracture toughness/MPa·m1/2

2.1 Performance Characteristics and Joining Problems of Ceramic Materials 27

28

2 Welding of Advanced Ceramic Materials

(3) Self-propagating high temperature synthesis (SHS) It is a technique for synthesizing materials by using the self-heating and selfconduction of the chemical reaction heat between reactants. The preparation of composite ceramics by self propagating high temperature synthesis method takes graphite, TiO2 powder and Al powder as raw materials. They are mixed according to the ratio of reaction formula 3TiO2 + 4Al + 3C = 3TiC + 2Al2 O3 . A large amount of heat is released when the combustion powder reacts, so as to maintain the progress of the reaction until the spread is completed.This method is simple, energy and time saving, and it can synthesize non-equilibrium phases and intermediate products that are difficult to synthesize by conventional processes. The reaction is not easy to control, and the products are loose and porous, but if pressurized at the same time, dense ceramics can be synthesized in one step, which could be an important process for the synthesis of composite materials in the future. (4) Spark plasma sintering, SPS Spark plasma sintering is a new sintering technology, which has the characteristics of fast heat rate, short sintering time, uniform grain size, conducive to controlling the structure of the sintered body,and high density of the obtained material. When sintering, Al2 O3 and TiC powder are mixed in a certain ratio and then placed in a container. Under high pressure, the powder is heated directly through the mold by using DC pulse current, so that the powder is instantly at a high temperature and the preparation process can be completed in just a few minutes. Due to the short preparation time, the production cost can be reduced. (5) Other preparation methods Other sintering methods for the preparation of composite ceramics include hotisostatic-pressing sintering (HIP), gas pressure sintering (GPS), multi-step sintering (MS), etc. Hot isostatic sintering involves placing powder compacts into a high pressure vessel so that the powder is subjected to isotropically balanced gas pressure during the heating process, so as to densify the material. Al2 O3 –TiC composite ceramics prepared using this technique have high densities and excellent properties, but their equipment is more complex and expensive. The gas pressure sintering method can inhibit the volatilization and decomposition of the reactants, and the sintering temperature is higher. The content of TiC cannot be too high. When it exceeds 30%, the densities and properties of Al2 O3 –TiC composite ceramics are lower. This method can be used to prepare components with complex shapes, and the properties of the prepared materials are slightly lower than those of hot pressing method and hot isostatic pressing method. Multi-step sintering, i.e., pressureless sintering or self-propagation of high-temperature synthetic powders followed by hot isostatic pressing or hot pressing sintering, etc. This method uses less burning aid, and the properties of the prepared material can be close to those prepared by hot pressing, but the process time is longer. The cost of above methods are more higher. The preparation method has a great influence on the mechanical properties of the composite ceramics. For example, the mechanical properties of Al2 O3 –TiC

2.1 Performance Characteristics and Joining Problems of Ceramic Materials

29

composite ceramics prepared by different methods at the room temperature are shown in Table 2.8. It can be seen that the material made by the pressureless sintering method has poorer properties; the material strength of hot isostatic sintering is better than that of other preparation methods; the material sintered by the selfpropagating high-temperature synthesis method is not as strong as the hot-pressure sintered and gas pressure sintered materials. The Al2 O3 –TiC ceramic prepared by the self-propagating high-temperature synthesis method + hot-pressure sintering (or hot isostatic sintering) ceramics have better overall performance, and the material strength, hardness and toughness are better than those made by single preparation methods. Al2 O3 –TiC composite ceramics are often used as high speed cutting tools or high temperature heat generators, so their thermal stability performance is critical. Studies on the high temperature oxidation resistance of Al2 O3 –TiC ceramics showed that trace oxidation of Al2 O3 –TiC composite ceramics occurred at 400 °C. When the temperature T < 900 °C, the oxidation process generates brittle TiO2 by phase interface reaction, and the oxidation increment versus time is given by the equation: 1−(1−α)2/3 = Kt. When the temperature is between 900–1100 °C, the oxidation mechanism changes to parabolic, and the oxidation increment versus time is given by the equation: α 2 = Kt. Table 2.8 Properties of Al2 O3 –TiC based composite ceramics prepared by different processes Composition(weight percentage)/% A12 O3 –26TiC

Preparation method

Density /g·cm−3

Bending strength /MPa

Vickers hardness /GPa

Breaking tenacity /MPa·m1/2

PS

–

386 ~ 574

18.1 ~ 19.3

4.2 ~ 4.6

SHS + HIP

–

511 ~ 765

20.9 ~ 23.1

4.2 ~ 4.6

A12 O3 –28TiC

HP

4.16

516

20.9

5.2

A12 O3 –30TiC

GPS

–

563 ~ 639

18.3 ~ 19.1

3.6 ~ 3.8

HIP

4.30

780

20.0

–

SHS + PS

4.23

307 ~ 467

20.7 ~ 21.3

4.1 ~ 4.7

SHS + PS + HIP

4.25

435 ~ 661

21.9 ~ 23.7

4.1 ~ 5.7

SHS + HP

4.23

416 ~ 832

22.2 ~ 24.6

3.6 ~ 3.8

A12 O3 –37TiC

SHS

–

587

18.4

4.5

A12 O3 –40TiC

SHS + HP

4.47

658 ~ 824

21.4 ~ 23.8

5.1 ~ 5.9

A12 O3 –47TiC

SHS + HP

4.54

571 ~ 941

20.8 ~ 22.0

5.5 ~ 5.9

SHS + HP

4.27

537 ~ 597

22.8 ~ 23.0

5.5 ~ 6.1

Note PS is pressureless sintering; HP is hot-pressure sintering; SHS is self-propagating hightemperature synthesis; GPS is pneumatic sintering; HIP is hot isostatic sintering

30

2 Welding of Advanced Ceramic Materials

2.1.2 Basic Requirements for Ceramic-to-Metal Joining Engineering ceramic materials have a broad application prospect in aerospace, machinery, metallurgy, chemical industry, electronics, etc. due to their excellent properties of high strength, corrosion resistance, low thermal conductivity and high wear resistance. However, the inherent hardness and brittleness of ceramic materials make them difficult to be processed and to make complex-shaped components, which are greatly restricted in engineering applications. One of the ways to promote the practical use of ceramics is to join it with high plastic toughness of metal materials to make composite components, so as to give full play to the performance advantages of the two materials and make up for their respective shortcomings. Therefore, welding is one of the key technologies to promote the application of ceramics. (1) Form of ceramic joining Ceramic materials have poor machinability, low plasticity and impact toughness, weak thermal shock resistance, and make it difficult to manufacture large size and complex shape parts. Therefore, ceramics are usually used in composite structures with metallic materials. When ceramics and metallic materials are successfully joined, ceramics will provide additional functionality to the part and improve its application performance. Therefore, a reliable connection between ceramics and metallic materials is the key to promote the application of ceramic materials. Welding is an important form of processing in which ceramics are used in production. For example, welding of ceramics to metals occupies a very important place in the nuclear industry and in the production of electric vacuum devices. The joining of ceramic materials can take several forms as follows. ➀ Joining of ceramics to metallic materials. ➁ Joining of ceramics to non-metallic materials (e.g. glass, graphite, etc.). ➂ Joining of ceramics to semiconductor materials. (2) Requirements for joint performance. More applications are the welding of ceramic and metal materials. The application of this welding structure in electrical appliances, electronic devices, nuclear energy industry, aerospace and other fields are gradually expanding, and the requirements for ceramic and metal joint performance are also increasingly high. The general requirements for the performance of ceramic–metal joints are as follows. ➀ The ceramic-to-metal welded joint must have high strength, which is a basic performance requirement for welded structural parts. ➁ The welded joint must be gas-tight with vacuum. ➂ The joints shall have minimum residual stress and shall be heat resistant, corrosion resistant and thermally stable during use. ➃ The welding process should be as simple as possible, with a stable process and low production costs.

2.1 Performance Characteristics and Joining Problems of Ceramic Materials

31

2.1.3 Problems with Ceramic-to-Metal Joining Due to the essential differences between the atomic structures of ceramic materials and metals, and the special physicochemical properties of ceramic materials themselves, there are a number of problems exist in both the joining of ceramics to metals and the joining of ceramics themselves. When a ceramic is connected to a metal, an interface needs to be made between the connecting materials in order to achieve a reliable coalesce between the two. This interface material should meet the following requirements. (i)

The interface material has a similar coefficient of linear expansion to the material being welded. (ii) A reasonable type of bonding, i.e. ionic/covalent bonding. (iii) Lattices mismatch between ceramic and intermetallic. The main problems that arise in welding ceramics to metallic materials are as follows. (1) Thermal expansion and thermal stress in ceramic-to-metal welding The linear expansion coefficient of ceramics is relatively small, which is quite different from that of metal. When ceramics and metals were jointed by heating, the thermal expansion and contraction causes a large residual stress in the joint area, weakening the mechanical properties of the joint; cracks can also occur when the thermal stress is large, leading to fracture and damage of the connected ceramic joint. One way to control stress is to reduce the temperature gradient in and around the weld area as much as possible during welding, and control the heating and cooling rates. Reducing the cooling speed is conducive to stress relaxation and stress reduction. Another way to reduce stress is to use a metallic interlayer, using a plastic material or a metallic material with linear expansion coefficient close to the coefficient of linear expansion of the ceramic. (2) Ceramics and metals are difficult to wet Ceramic materials are poorly wettable or not wettable at all. When brazing or diffusion bonding is used to join ceramic and metal materials, it is difficult to select a suitable brazing material to bond to the substrate because the molten metal is difficult to wet on the ceramic surface. In order to achieve a brazed joint between a ceramic and a metal, one of the most essential conditions is to make the brazing material wettable to the ceramic surface, or to improve the wettability to the ceramic, and finally to achieve a brazed joint. For example, good wettability can be obtained by using the active metal Ti to react at the interface to form a compound of Ti. In the ceramic joining process, the ceramic surface can also be metallized (physically or chemically overlaid with a layer of metal) before the ceramic-to-ceramic or ceramic-to-metal joint is made. This method actually turns the ceramic-to-ceramic or ceramic-to-metal connection into a metal-to-metal connection, but the bond strength of joints in this method is not high and it is mainly used for sealed welds.

32

2 Welding of Advanced Ceramic Materials

(3) Easily formed brittle compounds As the physical and chemical properties of ceramics and metals are very different, in addition to bonding conversion at the interface of the joint, various chemical reactions are prone to occur, generating various carbides, nitrides, silicides, oxides and multiple compounds at the bonding interface. These compounds are hard and brittle, which are the main cause of cracks and brittle fracture of the joint. When determining the phases of interfacially brittle compounds, it is necessary to prepare a variety of samples for calibration, due to large errors in the quantitative analysis of some light elements (C, N, B, etc.). The phase structure determination of multiple compounds is generally determined by comparing the X-ray diffraction method with the standard diffraction pattern but some compounds do not have standard patterns, which makes it difficult to determine the phase. (4) Ceramic–metal bonding interface Ceramic and metal joints have differences in atomic structure energy levels between the interfaces. Ceramics are bonded to the metal by a transition layer (diffusion or reaction layer). The interfacial reaction between the two materials has a great influence on the formation, microstructure and properties of the joint. The interfacial reaction and microstructure of the joint is an important topic in the study of ceramic– metal welding. Ceramic materials contain mainly ionic or covalent bonds, exhibit very stable electronic coordination and are difficult to be wetted by metal brazing materials with metallic bonds, so it is difficult to create a fusion joint between the metal and the ceramic by the usual fusion welding methods. When brazing ceramic with metal brazing filler material, either the surface of the ceramic is metallized first, the surface of the ceramic to be brazed is modified, or active elements are added to the brazing material so that chemical reaction occurs between the brazing material and the ceramic, and through the reaction the surface of the ceramic decomposes to form a new phase, generating a chemical adsorption mechanism, so as to form a solid ceramic–metal bonded interface.

2.1.4 Joining Methods of Ceramics and Metal The joining methods between ceramics and metals include mechanical connection, bonding and welding. The main welding methods commonly used are brazing, diffusion bonding, electron beam welding, laser welding, etc., as shown in Table 2.9. It is very difficult to weld ceramics and metals directly, and can hardly be achieved by general fusion welding methods, or even by direct welding. Therefore, special technological measures must be taken for the welding of ceramics and metals, so that the metal can wet the ceramic or react chemically with them. The wetting of the metal to the ceramic and the chemical reaction between the metal and the ceramic,

2.1 Performance Characteristics and Joining Problems of Ceramic Materials

33

Table 2.9 Joining methods of ceramics to metal Joining methods of ceramics to metal

Brazing

Ceramic surface metallization method

Powder metallurgy bonding Other metallization methods

Mo–Mn method Mo–Fe method Steam coating metallization Sputtering metallization Ion coating

Active metallization Ti–Ag–Cu method Ti–Ni method, Ti–Cu method, Ti–Ag method Oxide solder method Fluoride solder method Diffusion bonding Other methods

Direct diffusion bonding Indirect diffusion bonding(Add middle layer in diffusion bonding) Electron beam welding (EBW) Laser welding (LW) Ultrasonic pressure welding (UPW)

as well as the difference between the thermal expansion and contraction of the two during the connection and the resulting thermal stress, even causing cracking, are the main problems when joining ceramics to metals. (1) Ceramic-to-metal brazed joining The most applied ceramic–metal joining methods is the brazing method, which is generally divided into indirect brazing and direct brazing. The classification, principles and applicable materials of ceramic-to-metal brazing methods are shown in Table 2.10. Indirect brazing (also known as the two-step method) involves metallizing the ceramic surface firstly and then brazing with a common brazing filler material. The most commonly used method of ceramic surface metallization is the Mo-Mn method. In addition, physical vapor deposition (PVD), chemical vapor deposition (CVD), thermal spraying, and ion implantation are also used. The indirect brazing process is complex and its application is somewhat limited. The direct brazing method (also known as the one-step method), also known as reactive metallization brazing. It involves the addition of reactive elements, such as the transition metals Ti, Zr, Hf, Nb, Ta, etc., to the brazing filler material, which decompose the ceramic surface through chemical reaction to form a reaction layer. The reaction layer mainly consists of metal and ceramic compounds, and most of these products exhibit the same structure as the metal and can therefore be wetted by the molten metal. The direct brazing method can make the manufacturing process

Add some metals (such as Ag, Cu, Ni, Ceramic–metal joining etc.) to the oxidizing metals (Ti, Zr, Nb, Ta, etc.) to form a low melting point alloy as solder (the molten metal of this solder has small surface tension, viscosity and good wettability), add it to the gap between the connected ceramics and metals, and heating and brazing in an inert atmosphere furnace such as vacuum or Ar

The composite oxide glassy solder Ceramic–metal joining with a melting point lower than that of the connected ceramics and metals is used, mixed into paste with organic adhesive, embedded in the joint, and then heated and fused in hydrogen

Active metallization method

Ceramic fusion method

Applied materials

Principle

Mo or Mo Mn powder (particle size 3 Ceramic–metal joining ~ 5 μm) is mixed with organic solvent to form a paste as solder, which is coated on the ceramic surface and heated in steam atmosphere for brazing

Classification

Mo-Mn method

Table 2.10 Classification, principles and applicable materials of ceramic–metal brazing methods Illustration

(continued)

Al2 O3 -CaO-MgO-SiO2 solder is used for the connection between ceramics and heat-resistant metals, and the heating temperature is above 1200 °C. Al2o3-mno-sio2 solder is used to connect ceramics with ferrous alloys and heat-resistant metals, and the heating temperature is above 1400 °C

It is suitable for occasions with large output, and the shape of the workpiece can be arbitrary. When Al2O3 is connected with metal, Ti-Cu, Ti-Ni, Ti–Ni–Cu, Ti-Ag–Cu, Ti-Au-Cu and other solders can be used; Where high temperature strength is required, Ta, Cr, Mo, Nb and other solders can be added to Ti-V system and Ti-Zr system, and the brazing temperature is 1300 ~ 1650 °C

It is used for the joining between Al2 O3 and other oxide ceramics and metals, such as the sealing of the connection between ceramics and metals in various electronic tubes and electrical machinery

34 2 Welding of Advanced Ceramic Materials

Amorphous binary alloy (Ti-Cu Si3 N4, SiC and other ceramic - ceramic 、Ti-Ni or Zr-Cu、Zr–Ni) foil with joining, Si3 N4, SiC to metal joining thickness of about 40 ~ 50 μm and width of about 10 μm is used as brazing filler metal, placed in the joint surface, and then heated and brazed in vacuum or Ar atmosphere furnace

The surface friction function and stirring effect of ultrasonic vibration are used, and Sn–Pb alloy solder (usually added with Zn, Sb, etc.) is used for dip brazing

Amorphous alloy method

Ultrasonic soldering method

Glass, Al2 O3 ceramic joining, etc

Copper oxide (CuO) powder (particle The joining between oxide ceramics size 2 ~ 5 μm) is used as intermediate themselves (Al2O3, MgO, ZrO2) and material, heated in vacuum or oxide ceramics- metal oxidizing atmosphere, and brazed by the good wettability of molten copper on Al2O3 ceramic surface and oxide reaction

Applied materials

Principle

Classification

CuO method

Table 2.10 (continued)

(continued)

When Al2 O3 with purity (mass fraction) of 96% is brazed with Sn Pb filler metal and Zn, the joint strength can be greatly improved. Al2 O3 with purity (mass fraction) of 99.6% is difficult to be brazed by this method

It is a variant of active metallization method. Using Cu–Ti alloy foil as solder to connect non oxide ceramics such as Si3 N4 -Si3 N4 or SiC–SiC, higher joint strength can be obtained

The usual brazing condition is: the vacuum degree is 6.67 × 10−5 Pa; Heating for 20 min at about 600 °C in a vacuum furnace

Illustration

2.1 Performance Characteristics and Joining Problems of Ceramic Materials 35

Use the hydroxide heat-resistant glass Glass, Al2 O3 ceramic joining, etc as the intermediate layer in the joint, heat it under Ar or N2 atmosphere and irradiate it with laser to activate it for brazing

Applied materials

Principle

Classification

Laser activation brazing method

Table 2.10 (continued) –

Illustration

36 2 Welding of Advanced Ceramic Materials

2.1 Performance Characteristics and Joining Problems of Ceramic Materials

37

of ceramic structural parts simple and has become one of the research hotspots in recent years. The key to direct brazing of ceramics is to use an active brazing material. On the premise that the solder can wet the ceramics, it is also necessary to consider whether the difference in the coefficient of linear expansion between the ceramic and the metal during high temperature brazing will cause cracking.The insertion of an intermediate buffer layer between the ceramic and the metal is effective in reducing stress and improving the strength of the joint. The limitations of direct brazing are the lower high temperature strength of the joint and the spreading of the brazing material in large-area brazing. (2) Solid-state diffusion bonding Solid-state diffusion bonding is generally divided into two forms, direct and indirect, mainly by vacuum diffusion bonding, but also by hot isostatic method of diffusion bonding. The classification, principles and applicable materials of ceramic-to-metal solid-phase joining methods are shown in Table 2.11. Solid-state diffusion bonding is a common method of ceramic–metal joining, which refers to the process of forming a reliable connection as a whole by making the connected surfaces contact each other under a certain temperature and pressure, expanding the physical contact of the connected surfaces by local plastic deformation of the contact surfaces, or by the transient liquid phase generated on the connected surfaces, and then diffusing occur between the atoms of the bonding layer. Its distinctive features are stable joint quality, high bonding strength, good high temperature performance of the joint and corrosion resistance. In solid-phase diffusion bonding, the bonding temperature, pressure, time and surface condition of the weld are the main factors affecting the quality of diffusion bonding. The bonding of the interface in the solid-phase diffusion connection is achieved by plastic deformation, diffusion and creep mechanisms, and its bonding temperature is high, and the ceramic–metal solid-phase diffusion bonding temperature is usually 90% of the melting point of the metal. Due to the mismatch between the linear expansion coefficient and elastic modulus of ceramics and metals, it is easy to generate large stresses near the interface, which makes it difficult to realize direct solid-phase diffusion bonding. In order to relieve the residual stresses in the ceramic–metal joint and to control the interfacial reaction, inhibit or change the interfacial reaction products to improve the joint performance, diffusion bonding with an intermediate layer is often used. (3) Fusion welding of ceramics to metals The high melting point and ceramic pyrolysis make it difficult to join ceramics and metals using the general fusion welding method. Although the fusion welding method is fast and efficient, and can form a joint with stable performance at high temperature, in order to reduce the welding stress and prevent the generation of cracks, the auxiliary heat source must be used for preheating and slow cooling, and the process parameters are difficult to control, and the equipment investment is expensive. The main methods of fusion welding of ceramics and metals include electron beam welding, laser welding, and arc welding, etc. Because ceramic materials are

Principle

The joint surface between ceramic and metal is covered with metal foil, heated to a temperature lower than the melting point of metal (1065 °C for Cu) in a furnace with slightly oxidizing atmosphere (oxygen, phosphorus, sulfur, etc.), and the joining is realized by the eutectic effect after the reaction between gas and metal

The bonding surface is processed into approximate mesh, and the workpiece is put into the vacuum chamber after assembly (vacuum degree 133 × 10−3 Pa), apply static pressure in all directions at appropriate temperature (pressure 50 ~ 250 MPa), and the joint will be formed in a short time (in order to promote interface bonding, sometimes metal powder or TiN and other ceramic powder are placed on the interface as a middle layer)

Classification

Gas-metal eutectic method

Hot isostatic pressing (HIP)

Illustration

Ceramic-ceramic joining, ceramic–metal joining. It is especially suitable for the connection between Al2 O3 , Zr2 O, SiC and metal

(continued)

Due to the simultaneous compression in all directions, the interface is closely connected as the plastic deformation in the connection area is small and the strength of the joint is higher. Ceramic powder cover the metal surface and can form a thick and dense surface layer

The bonding between ceramics and Cu, – Fe, Ni, Co, Ag, Cr, etc. It is especially suitable for the joining between Al2 O3 and Cu

Applied materials

Table 2.11 Classification, principles and applicable materials of ceramic–metal solid-phase joining methods

38 2 Welding of Advanced Ceramic Materials

It is a method of direct bonding by means of reaction between ceramic and metal after contact. It can also be divided into no pressure mode and pressure method

A diffusion brazing method in which Ceramic-metal joining an intermediate layer (solder) is sandwiched in the gap of the joint, then heated and pressurized in a vacuum furnace, and connected through the diffusion of interface atoms

Reaction bonding method

Diffusion bonding method

Illustration

Joining between oxide ceramics and precious metals (Pt, Pd, Au, etc.) and transition metals (such as Ni), ceramic-metal joining

It is used for the joining between nickel base heat-resistant alloy and Si3 N4 in diesel engine exhaust valve

No pressure method: heat in the atmosphere (Ar or vacuum) to 90% of the melting point of the metal, and apply the pressure that makes the bonding surface physically contacted to connect; Pressure method: apply external pressure while heating in hydrogen atmosphere (the temperature is 90% of the melting point of the metal) to deform the metal and form a joint

The joining between glass and metal, If external pressure is applied at the Al2 O3 and Cu, Fe, Ti, Al, etc. It is also same time, the joining can be realized under low voltage and temperature suitable for the joining between ceramics and semiconductors

When the joint area is heated to high temperature, DC voltage is applied to polarize the bonding interface, and direct connection is carried out through metal diffusion to ceramics. Usually, 0.1 ~ 1.0 kV DC voltage is added to the connection area for 40 ~ 50 min at the temperature of 500 ~ 600 °C

Applied materials

Principle

Classification

Diffusion bonding add voltage method

Table 2.11 (continued)

2.1 Performance Characteristics and Joining Problems of Ceramic Materials 39

40

2 Welding of Advanced Ceramic Materials

extremely brittle and have very low plasticity and toughness, making their fusion welding very limited. The methods, principles and applicable materials of ceramicto-metal fusion welding are shown in Table 2.12. Brazing and diffusion bonding method of ceramic–metal connection are relatively mature and widely used; Electron beam welding and laser welding are also expanding their applications. In addition, the connection of ceramics and metal can also be achieved by ultrasonic pressure welding, friction pressure welding and other methods. Table 2.12 Ceramic–metal fusion welding methods, principles and applicable materials Classification

Principle

Applied materials

Illustration

Laser welding

It is a method of fusion welding by irradiating ceramic joint with high energy density laser beam. The laser adopts pulse oscillation mode with high output power. The workpiece needs to be preheated before welding to prevent cracks caused by thermal shock caused by centralized laser heating

Joining between oxide ceramics (Al2 O3 , mullite, etc.), Si3 N4 , SiC and ceramics

The preheating temperature of Al2 O3 is 1030 °C. Without the intermediate layer, the joint strength which close to the ceramic strength can be achivevd. Unfocused laser beam can be used during preheating. In order to increase the penetration, the welding speed should be slow, but too slow will cause the coarser grain

Electron beam welding

It is a method of fusion welding by using high energy density electron beam to irradiate the joint area

The same as laser welding. In addition, Al2 O3 and Ta, graphite and W can also be Joined by this method

The same as laser welding. Welding must also be done in a vacuum chamber

Arc welding

Heat the joint area with gas flame. When the temperature rises to the conductivity of the ceramic, apply voltage to the joint through the special electrode in the gas flame torch, so that the arc discharge between the joint surfaces and generate high heat for fusion welding

Joining between certain ceramic-ceramic, ceramic-certain metals (such as ZrB2 and Mo, Nb, Ta, ZrB2 , SiC and or Ta)

Carbide ceramics and boride ceramics with conductivity can be directly welded. During welding, it is necessary to control the rising rate of current and the maximum current value

2.2 Weldability Analysis of Ceramic Materials

41

2.2 Weldability Analysis of Ceramic Materials There are essential differences between ceramics and metals, coupled with the special physical and chemical properties of ceramics themselves, so there are many problems in the welding of ceramics and metals. The linear expansion coefficient of ceramics is relatively small, which is quite different from that of metal. Residual stresses will be generated in the welded joint, and higher stresses will lead to cracks at the joint and even cause fracture. The main problems in the welding of ceramics and metals include stresses and cracks, interfacial reactions, low bond strength, etc.

2.2.1 Welding Stress and Cracks The chemical composition and thermophysical properties of ceramics and metals are quite different, especially in terms of the linear expansion coefficient (shown in Fig. 2.1). For example, the of linear expansion coefficient of SiC and Si3 N4 are only 4 × 10–6 K−1 and 3 × 10–6 K−1 , respectively, while the linear expansion coefficient of aluminum and iron are as high as 23.6 × 10–6 K−1 and 11.7 × 10–6 K−1 , respectively. In addition, the elastic modulus of ceramics is also very high. In the heating and cooling process of welding, ceramic and metal produces a large difference in expansion and contraction, resulting in large thermal stresses in the vicinity of the joint. Due to the extremely uneven distribution of thermal stress, stress concentration occurs at the joint interface, resulting in cracks in the joint area. When concentrated heating, especially when using a high-energy dense beam heat source for fusion welding, the ceramic side near the welded joint produces a high stress zone, the ceramic itself is a hard and brittle material, it is easy to produce cracks in the welding process or after welding. Fig. 2.1 Linear expansion coefficient of ceramics and metals

42

2 Welding of Advanced Ceramic Materials

Welding of ceramics to metals is generally carried out at high temperatures and the difference between the welding temperature and room temperature is also an important factor in increasing the residual stress in the joint area. In order to reduce the stress concentration in ceramic–metal welded joints, it is effective to add a plastic material or a metal with linear expansion coefficient close to that of the ceramic as an intermediate layer between the ceramic and the metal. For example, Cu foil with a thickness of 20 mm is added as a transition layer between the ceramic and Fe–Ni-Co alloy, and a diffusion welded joint with tensile strength of 72 MPa can be obtained at heating temperature of 1050 °C, holding time of 10 min, and pressure of 15 MPa. The interlayer used in diffusion welding can reduce the diffusion temperature, the pressure and the holding time to facilitate interfacial diffusion and remove impurity elements and also to reduce the residual stresses generated in the joint area.The effect of interlayer thickness on the reduction of residual stresses in diffusion welding of Al2 O3 ceramics with 0Cr13 ferritic stainless steel is shown in Fig. 2.2. Materials with low modulus of elasticity and yield strength and good plasticity are mostly selected for the interlayer, and the stress in the ceramic/metal joint is reduced by plastic deformation of the interlayer metal or alloy. The use of a metal with a lower modulus of elasticity and yield strength as an interlayer is a way to transfer the stresses in the ceramic to the interlayer. The use of two different metals as a composite interlayer is also an effective way to reduce stresses in ceramic/metal welds. Generally Ni is used as a plastic metal and W as a low linear expansion coefficient material. The main metals commonly used as interlayer in diffusion bonding of ceramics and metals are Cu, Ni, Nb, Ti, W, Mo, copper-nickel alloys, steel, etc. The requirement for these metals is that the linear expansion coefficient is similar to that of ceramics, and there is no isotropic transformation occurs during the manufacture and operation of the component to avoid the sudden changes in the linear expansion coefficient, which can destroy the matching of ceramics and metals and lead to the failure of 1200

Residual stress /MPa

1000 800

M0

600

Cermet

400

Ti

200

Nb 0 0.0

0.5

1.0

1.5

2.0

2.5

3.0

Interlayer thickness /μm Fig. 2.2 Effect of interlayer thickness on residual stresses in Al2 O3 /stainless steel joint (Heating temperature 1300 °C, holding time 30 min, pressure 100 MPa)

2.2 Weldability Analysis of Ceramic Materials

43

the welded structure. The interlayer can be made directly from metal foil, or the metal powder can be pre- placed on the ceramic surface by vacuum evaporation, ion sputtering, chemical vapor deposition (CVD), spraying, electroplating, etc., and then welded to the metal. As the thickness of the intermediate layer increases, the residual stresses are reduced. This effect of Nb is most pronounced as the linear expansion coefficient of Nb is closed to alumina ceramics. However, the effect of the interlayer is sometimes complicated, and if there is a chemical reaction at the interface, the effect of the interlayer will depend on the type and thickness of the reactant. Improper selection of the interlayer can even lead to deterioration of the joint performance. For example, the formation of brittle phases due to chemical reactions or increased stresses due to mismatched linear expansion coefficient can lead to cracks in the joint area. When brazing ceramics and metals, some brazing materials with good plasticity and low yield strength, such as pure Ag, Au or Ag–Cu brazing filler metals, can be used in order to maximize the release of stress in the brazed joint; Sometimes low-melting point active brazing materials are also used, for example, Ag52-Cu20In25-Ti3 and In85-Ti15 indium-based brazing materials are used to vacuum braze AlN and Cu. Indium-based brazing materials have good wettability to AlN ceramics. Brazed joints with good microstructure and properties could be formed by controlling brazing temperature and time, as shown in Fig. 2.3. In order to avoid welding cracks in ceramic–metal joints, in addition to the add of interlayer or a reasonable choice of brazing material, the following process measures can be taken.

Brazing time

Brazing temperature

Bearing capacity

Bearing capacity

➀ Select the welded ceramic and metal reasonably so that the difference between the linear expansion coefficients of the two is as small as possible without affecting the service performance of the joint. ➁ The temperature gradient at and near the welding position should be reduced as much as possible, the heating rate should be controlled, and the cooling rate should be reduced to facilitate stress relaxation and reduce the welding stress.

Brazing temperature

Brazing time

Fig. 2.3 Effect of brazing temperature and time on the load carrying capacity of the joint

44

2 Welding of Advanced Ceramic Materials

➂ Take measures such as notching, protrusion and end thinning to reasonably design the ceramic-to-metal joint structure. Residual stresses are generated at the joint during diffusion bonding of ceramic to steel. The stresses arise due to the mismatch in thermal expansion between ceramic and steel and the large difference in elastic modulus. In addition, the strain hardening coefficient, yield stress, and interlayer thickness can also have an effect on the formation and distribution of stresses. When the stresses reach a certain strength, cracks may occur in different areas of the joint. When the thermal expansion coefficient of ceramics is lower than that of steel (α c < α m ), the stress and crack distribution in the ceramic-steel diffusion joint is shown in Fig. 2.4a. When the thermal expansion coefficient of ceramics is higher than that of steel (α c > α m ), the stress and crack distribution in the ceramic-steel diffusion joint is shown in Fig. 2.4b. In both cases, the cracks arise in the maximum tensile stress region on the ceramic side, because the ceramic side tends to weaken under tensile stress. For example, the presence of microcracks was also observed near the interface of the Al2 O3 –TiC/18–8 steel diffusion bonded joints specimen. The cracks present in the Al2 O3 –TiC ceramic side near the interface, forming longitudinal cracks roughly parallel to the interface, as shown in Fig. 2.5. The formation of longitudinal cracks within the Al2 O3 –TiC ceramic near the Al2 O3 –TiC/18–8 steel diffusion interface is due to the thermal expansion coefficient of the Al2 O3 –TiC ceramic (7.6 × 10–6 K−1 ) is lower than that of the 18–8 stainless steel (16.7 × 10–6 K−1 ) and the reaction products in the transition zone of the interface, forming longitudinal cracks parallel to the interface near the Al2 O3 –TiC ceramic interface. Microcracks were also observed within the reaction zone of the interlaye of the Al2 O3 –TiC/18–8 steel diffusion bonded joint specimens, as shown in Fig. 2.5. This type of crack is located within the reaction zone of the intermediate layer and a transverse crack perpendicular to the interface as shown in Fig. 2.5b.

ceramic

ceramic

tension tension

pressure

tension

pressure

steel

steel

(a) αc

αm

(b) αc

αm

Fig. 2.4 Illustration of thermal stresses and cracks in joints due to mismatch of thermal expansion coefficients

2.2 Weldability Analysis of Ceramic Materials

45

Al2O3-TiC

Al2O3-TiC

A A

B B

(a) Longitudinal crack

(b) Transverse crack

Fig. 2.5 Crack morphology on the Al2 O3 –TiC ceramic side near the diffusion bonding interface

Transverse crack is a common defect in ceramic-to-metal welded (diffusion, brazed) joints because the thermal expansion coefficient of the diffusion reaction layer is higher than that of the ceramic substrate in most cases. Through careful observation of the transverse cracks in the reaction zone of the interlayer, it is found that the transverse cracks in the interlayer reaction zone start in the interlayer reaction zone and expand to the interface between the reaction zone of the intermediate layer and the diffusion reaction zone on the steel side or cross the intersection of the two into the precipitation phase of the steel-side reaction zone and terminate their expansion. It does not no longer continue to expand across the precipitation phase boundary. For Al2 O3 –TiC/steel diffusion bonded joints, both longitudinal and transverse cracks are related to the interfacial reaction of the elements during diffusion bonding and residual stresses in the joint. The above longitudinal and transverse cracks were found in Al2 O3 –TiC/18–8 steel diffusion welding tests. The two main factors of crack formation are metallurgical and mechanical factors, for Al2 O3 –TiC/18–8 steel diffusion welded heads, Ti and Fe are completely miscible in the liquid state and limited dissolution in the solid state according to the Ti– Fe phase diagram, and Ti and Fe are prone to form TiFe and TiFe2 intermetallic compounds. During diffusion welding, the composite interlayer melts to form Cu– Ti liquid phase, and the molten Cu–Ti liquid phase diffuses into the 18–8 steel. At the same time, the elements (Fe, Cr, Ni) in the 18–8 steel also dissolve and diffuse into the Cu–Ti liquid phase, so that the front zone at liquid/solid interface will be enriched with Ti, Cu, Fe, Cr, and Ni. Ti is an active element, which is easy to react with Cu, Fe, Cr, and Ni. Fex Tiy , Fe (Ti), TiFe and TiFe2 intermetallic compounds are hard and brittle compounds, which have the characteristics of high hardness and low plasticity, so that the plasticity of the joint is reduced and cracks are easy to appear. This is the metallurgical factor for the formation of cracks in the reaction zone of the interlayer. Although Ti-Cu and Ti-Ni compounds may also be generated at the interface of Al2 O3 –TiC/18–8 steel, these compounds are not very brittle and have

46

2 Welding of Advanced Ceramic Materials

certain plasticity. Due to the presence of the brittle and hard Ti–Fe compound layer, the Al2 O3 –TiC/18–8 steel diffusion welded joints are more prone to tearing at the Ti-Fe compound layer. Compounds of Ti–Fe, Ti–Cu, Ti–Ni, etc. may also be formed in the interfacial transition zone of Al2 O3 –TiC/Q235 steel diffusion welded joints, but cracks in Al2 O3 –TiC/Q235 steel joints fracture when the cracks start in the CuTi compound layer, indicating that the Ti–Fe compound layer generated in the interfacial transition zone of Al2 O3 –TiC/Q235 steel is thin, and the high plasticity residual Cu present at the interfacial transition zone, which makes the joint plastic to a certain extent. It is not sufficient to cause damage when the Ti–Fe compound layer is thin. When the CuTi compound layer is thick, the plasticity of Cu in the interfacial transition zone is not sufficient to resist the brittle damage caused by the thicker CuTi layer. However, the shear strength of Al2 O3 –TiC/18–8 steel joints is lower than that of Al2 O3 –TiC/Q235 steel joints due to the superior plasticity of Cu–Ti compounds over Ti-Fe compounds. The thermophysical properties between Al2 O3 –TiC ceramics and steel are different, especially the linear expansion coefficient. These differences may cause residual thermal stresses, which are a mechanical factor of crack formation.

2.2.2 Interfacial Reactions and Interface Formation Processes (1) Interfacial reaction products The joining between ceramic and metal is bonded by means of a transition layer (diffusion or reaction layer). The interfacial reaction between the ceramic/transition layer/metal material has a strong influence on the formation and performance of the joint. The physical phase structure of the interfacial reaction products is the key to influence the ceramic–metal bonding. In the ceramic and metal diffusion bonding, the ceramic and metal react at the interface to form compounds. The structure of the formed phase depends on the type of ceramic and metal (including the interlayer), but also with the welding conditions (such as heating temperature, surface condition, intermediate alloy and thickness, etc.). The interface reaction of SiC ceramic and metal generally produces the metal carbide, silicide or ternary compounds, and sometimes also produces tetrads and other multi compounds or amorphous phases, the reaction equation is as following: Me + SiC → MeC + MeSi Me + SiC → MeSix C y For example, SiC reacts with the Zr at the interface could form ZrC, Zr2 Si and the ternary compound Zr5 Si3 Cx . The possible interfacial reaction products in SiC ceramics with metal joints are shown in Table 2.13.

2.2 Weldability Analysis of Ceramic Materials

47

Table 2.13 Interfacial reaction products of SiC ceramics with metal connected joints Joint assembly

Temperature/K Time/min Pressure Atmosphere/mPa Reaction producs /MPa

SiC/Ni

1223

90

0

Ar

Ni2 Si + C, Ni5 Si + C, Ni3 Si

SiC/Fe–16Cr

1223

960

0

Ar

(Ni, Cr)2 Si + C, (Ni, Cr)5 Si2 + C, (Cr3Ni5Si1.8)C

SiC/Fe–17Cr

1223

960

0

Ar

(Fe, Cr)7 C3 , (Fe, Cr)4 SiC, α +C

SiC/Fe–26Ni

1223

240

0

Ar

(Fe, Ni)2 Si + C, (Fe, Ni)5 Si2 + C, α + C

SiC/Ti–25Al–10Nb 973

6000

0

–

(Ti, Nb)C, (Ti, Nb)3 (Si, Al), (Ti, Nb)5 (Si, Al)3 , (Ti, Nb)5 (Si, Al)3 C

SiC/Zr/SiC

1573

60

7.3

1.33

Zr5 Si3 Cx , Zr2 Si, ZrCx

SiC/Mo

1973

60

20

20 000

Mo5 Si3 C, Mo5 Si3 , Mo2 C

SiC/Al–Mg/SiC

834

120

50

4 000

Mg2 Si, MgO, Al2 MgO4 , Al8 Mg5

SiC/Ti/SiC

1673

60

7.3

1.33

Ti3 SiC2 , Ti5 Si3 Cx , TiC, TiSi2 , Ti5 Si3

SiC/Ta/SiC

1773

480

7.3

1.33

TaC, Ta5 Si3 Cx , Ta2 C

SiC/Nb/SiC

1790

120

7.3

1.33

NbC, Nb2 C, Nb5 Si3 Cx , NbSi2

SiC/Cr/SiC

1573

30

7.3

1.33

Cr5 Si3 Cx , Cr3 SiCx , Cr7 C3 , Cr23 C6

SiC/V/SiC

1573

120

7.3

1.33

V5 Si3 Cx , V5 Si3 , V3 Si, V2 C

SiC/Al/SiC

873

120

50

4 000

Al–Si–C–O Amorphous phase

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2 Welding of Advanced Ceramic Materials

The interfacial reactions between Si3 N4 ceramics and metals generally produce nitrides, silicides or ternary compounds of the metal. For example, the interfacial reactions between Si3 N4 and Ni-20Cr alloys produce Cr2 N, CrN and Ni5 Si2 , but does not generate compounds with Fe, Ni and Fe–Ni alloys. The possible interfacial reaction products in Si3 N4 ceramic–metal joints are shown in Table 2.14. The interfacial reaction products obtained in Si3 N4 ceramic and Ti, Mo, and Nb interfacial reactions are different when N2 and Ar are used as protective atmospheres, respectively, even if the same heating temperature and time are used. The interfacial reactions between Al2 O3 ceramics and metals generally produce oxides, aluminides or ternary compounds of that metal, e.g. the reactions between Al2 O3 and Ti produce TiO and TiAlx . The possible products of interfacial reactions in Al2 O3 ceramic–metal joints are shown in Table 2.15. The reactions between ZrO2 and metals generally produce oxides and zirconides of that metal, e.g. The reactions between ZrO2 and Ni produce NiO1-x , Ni5 Zr and Ni7 Zr2 . (2) Formation of diffusion interface In the diffusion bonding process of ceramics and metals with a composite interlayer, the diffusion ability of the interlayer elements in the two base materials differs due to the great differences in microstructure, composition, physical and chemical properties and mechanical properties of ceramics and metals, resulting in different degrees of reaction between the interlayer and the base materials on both sides, thus producing asymmetry in the formation process of the diffusion bonding interface. Taking the diffusion bonding of Al2 O3 –TiC composite ceramics and W18Cr4V HSS as an example, there are obvious asymmetries in the interfacial structure and element distribution. To elucidate the Al2 O3 –TiC/W18Cr4V diffusion bonding process, Fig. 2.6 illustrates the asymmetry of the formation process of the interface between Al2 O3 –TiC ceramic and W18Cr4V steel in diffusion bonding.The Al2 O3 –TiC/W18Cr4V diffusion bonding process is divided into four stages. Stage 1: Ti–Cu–Ti interlayer melting stage. Figure 2.6a shows the Ti–Cu–Ti composite interlayer is placed between Al2 O3 –TiC ceramic and W18Cr4V steel before diffusion bonding. After the diffusion bonding process starts, pressure is gradually applied to the upper surface of the specimen, and the softer Cu in the interlayer undergoes plastic deformation, accelerating the interface contact and providing a channel for atomic diffusion and interfacial reactions. With the increase of heating temperature, the solid phase diffusion occurs between Al2 O3 –TiC/W18Cr4V interface The element diffusion distance is short due to the small diffusion coefficient of the element in the solid state. According to the phase diagram of the Cu–Ti binary alloy (Fig. 2.7), at the Cu/Ti interface, the CuTi phase is generated firstly instead of Cu3 Ti2 . When the temperature increases to 985 °C, a liquid phase region with a large concentration gradient starts to appear at the local contact site of the Cu/Ti interface (Fig. 2.6b), and then the liquid phase spreads throughout the interface and extends to both sides of Cu and Ti (Fig. 2.6c). Since the diffusion coefficient of Cu (DCu = 3 × 10 m−92 /s) is larger

120

1323

1523

1273

1073–1473

1473

Si3 N4 /V/Mo

Si3 N4 /AISI316

Si3 N4 /Ni–Cr

Si3 N4 /Ni–Nb–Fe–36Ni/MA6000 60

95

1440

90

60

1673

60 60

1473

1473

60

1473

Si3 N4 /Nb

60

1473

Si3 N4 /Cr

Si3 N4 /Mo

120

1073

Si3 N4 /Ti

60

Si3 N4 /Ni–20Cr /Si3 N4

240

1200

1473

Si3 N4 /Incoloy909

Time/min

Temperature/K

Joint assembly

Table 2.14 Interfacial reaction products of Si3 N4 ceramic-to-metal joint

100

0

7

20

7.3

0

—

0

0

0

0

50

200

Pressure/MPa

—

Ar

1

5

Ar

N2

—

Ar

N2

Ar

N2

0.14

Ar

Atmosphere/mPa

NbN, Ni8 Nb6 , Ni6 Nb7 , Ni3 Nb

Ni2 Si, Ni3 Si2 , Cr3 Si, Cr5 Si3 , (Cr, Si)3 Ni2 Si

α-Fe, γ-Fe

V3 Si, V5 Si3

Nb5 Si3 , NbSi2 , Nb2 N, Nb4.62 N2.14

Nb5 N, Nb4 N3 , Nb4.62 N2.14

CrN, Cr2 N, Cr3 Si

Mo3 Si, Mo5 Si3 , MoSi2

Mo3 Si, Mo5 Si3

TiN + Ti2 N + Ti5 Si3

TiN + Ti2 N

CrN, Cr2 N, Ni5 Si2

TiN, Ni16 Nb6 Si7

Reaction producs

2.2 Weldability Analysis of Ceramic Materials 49

Time/min

1273

1373

Ni/ZrO2 /Zr

ZrO2 /Ni–Cr-(O)/ZrO2

30

1373

—

Al2 O3 /Ta-33Ti

Al2 O3 /Ni

1440

1313

Al2 O3 /Cu/Al2 O3

30

180

60

—

30

1143

1273

Al2 O3 /Ti/1Cr18Ni9Ti

30

Al2 O3 /Cu/AISI1015

Temperature/K

803

Joint assembly

Al2 O3 /Cu/Al

10

2

—

3

5

3

15

6

Pressure/MPa

Table 2.15 Interfacial reaction products of Al2 O3 ceramic-to-metal joint Atmosphere/mPa

100

1

—

0.13

0.13

O2

1.33

1.33

Reaction producs

NiO1−x Cr2 O3−y ZrO2−z , 0 < x,y,z < 1

Ni5 Zr, Ni7 Zr2

NiO, Al2 O3 , NiO·Al2 O3

TiAl, Ti3 Al, Ta3 Al

Cu2 O, CuAlO2

Cu2 O, CuAlO2 , CuAl2 O4

TiO, TiAlx

Al + CuAlO2 , Cu + CuAl2 O4

50 2 Welding of Advanced Ceramic Materials

2.2 Weldability Analysis of Ceramic Materials

51

Fig. 2.6 Schematic representation of Al2 O3 –TiC/W18Cr4V diffusion connection interface formation a Initial state; b Local liquefaction of Cu–Ti interface; c Cu–Ti liquid phase covered the whole interface; d Cu is completely melted; e Ti is completely melted; f Cu–Ti diffuses mutually and Ti reacts with the base metal to form a reaction layer; g Cu–Ti liquid phase reaction forms compounds, and the reaction layer increases; h Homogenization of solid phase composition

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2 Welding of Advanced Ceramic Materials

Fig. 2.7 Phase diagram of Cu–Ti binary alloy

than that of Ti (DTi = 5.5 × 10–14 m2 /s), Cu diffuses faster than Ti and all of the Cu melts first (Fig. 2.6d) and then all of Ti also melts (Fig. 2.6e). The melted Ti and Cu form a Cu–Ti liquid phase with a concentration gradient filling the entire interface between Al2 O3 –TiC and W18Cr4V. Due to the pressure applied on the specimen surface, part of the liquid phase is extruded from the interface under the pressure and the Cu–Ti liquid phase zone narrows. The melting of the Ti–Cu–Ti interlayer is very rapid due to the presence of liquid phase diffusion and concentration gradients, and the completion time of interlayer melting is very short compared to the entire joining time (instantaneous liquid phase). At this stage, the diffusion of Ti to the base metals on both sides is limited. The centerline of the liquid phase region remains the original interlayer centerline after the interlayer melting is completed (Fig. 2.6e). Phase 2: Homogenization of the liquid phase composition. The concentration distribution of the freshly melted Cu–Ti liquid phase is not uniform, so there is further interdiffusion between Cu and Ti. Ti is an active element and the Cu–Ti liquid phase filler metal is wettable to the Al2 O3 –TiC/W18Cr4V steel interface. The applied pressure promotes the expansion of the Cu–Ti liquid alloy. During this process, Ti in the Cu–Ti liquid phase filler metal diffuses and reacts to both sides of the Al2 O3 –TiC/W18Cr4V interface (Fig. 2.6f), and elements in the base material also diffuse to the Cu–Ti liquid phase, homogenizing the composition in the liquid phase region. Since there are tiny voids between the grains of Al2 O3 –TiC ceramics, it is favorable for Ti to diffuse in Al2 O3 –TiC ceramics. The C atoms in W18Cr4V steel are very small and diffuse rapidly, which are easy to diffuse into the

2.2 Weldability Analysis of Ceramic Materials

53

Cu–Ti liquid phase and react with Ti at the liquid/solid interface to form TiC, which hinders the diffusion of Ti into W18Cr4V. Therefore, the diffusion distance of Ti to Al2O3 tic is greater than that to W18Cr4V side. At the end of this stage, the liquid phase centerline shifts to the Al2 O3 –TiC side. Stage 3: Liquid solidification process. With the diffusion of Ti atoms at the liquid–solid interface, at the interface between Al2 O3 –TiC and the liquid phase, Ti reacts with Al and O atoms in Al2 O3 –TiC ceramics, generating Ti–Al and Ti–O compound reaction layers. At the interface between the liquid phase and W18Cr4V, Ti reacts with Fe and C in W18Cr4V steel to generate TiC, FeTi reaction layers. The solute atoms in the liquid phase zone gradually decrease, and when the concentration of solute atoms is less than the solidus concentration, the liquid phase starts to solidify (liquid–solid interface advances into the liquid phase), the interfacial reaction layer continues to grow, and the Cu–Ti liquid phase gradually decreases, and finally the liquid phase zone disappears completely, as shown in Fig. 2.6g. Since the diffusion rate of Ti to the Al2 O3 –TiC side is greater than that to the W18Cr4V side, the thickness of the reaction layer on the Al2 O3 –TiC side is greater than that on the W18Cr4V side at the end of liquid-phase solidification, and the centerline of the interface is shifted from the position of the original centerline of the interlayer. Phase 4: Homogenization of the solid phase composition. After complete solidification in the liquid phase zone, there is still a large concentration gradient of elements in the Al2 O3 –TiC/W18Cr4V interfacial transition zone as the diffusion joining process proceeds. Through the holding temperature stage, the interfacial elements diffuse with each other and the composition in each reaction layer is further homogenized, forming an interfacial layer with relatively homogeneous composition, as shown in Fig. 2.6h. The homogenization of the solid phase composition takes a long time, and the interface generally does not reach complete homogenization. Therefore, the microstructure morphology and element distribution show asymmetry in the transition zone of Al2 O3 –TiC/W18Cr4V interface. (3) Mechanism of interface reactions in diffusion bonding. X-ray diffraction (XRD) analysis of the shear fracture of Al2 O3 TiC/W18Cr4V diffusion joint shows the presence of various reaction products such as TiC, TiO, Ti3 Al, Cu, CuTi, CuTi2 , FeTi, Fe3 W3 C in the Al2 O3 –TiC/W18Cr4V interfacial transition zone. These reaction products are located in different reaction layers in the Al2 O3 –TiC/W18Cr4V interfacial transition zone, see Fig. 2.8. The interfacial reactions occurring in each reaction layer in the interfacial transition zone from the Al2 O3 –TiC ceramic side to the W18Cr4V steel side are analyzed as follows. (i) Al2 O3 –TiC/Ti interface (reaction layer A). The Al2 O3 phase and the TiC phase in Al2 O3 –TiC composite ceramics will react more vigorously only at temperatures greater than 1650 °C. The diffusion bonding

54

2 Welding of Advanced Ceramic Materials

Fig. 2.8 Organization of the transition zone at the interface of the Al2 O3 –TiC/W18Cr4V diffusion connection

Al2O3-TiC A B C

D

W18Cr4V

temperature in the experiment was 1160 °C, which is much lower than 1650 °C. TiC is an ion-bonded compound of NaCl structure with Gibbs free energy ΔG° (TiC) = −190.97 + 0.016 T, which is little affected by temperature change. Ti, a transition metal element, is very reactive and is used as an active element in the joining of ceramics and metals, reacting with ceramics to form reaction layers. At the Al2 O3 –TiC/Ti interface, the reaction is mainly between Ti in the Ti-Cu–Ti interlayer and the Al2 O3 ceramic. During diffusion bonding of Al2 O3 –TiC/W18Cr4V, Ti reacts with Al2 O3 as following: 3Ti + Al2 O3 = 3TiO + 2Al

(2.1)

TiO and Al atoms are generated. According to the Ti–Al binary phase diagram, at the diffusion bonding temperature, the following reactions between Ti and Al may occur. Ti + 3Al = TiAl3

(2.2)

Ti + Al = TiAl

(2.3)

3Ti + Al = Ti3 Al

(2.4)

Since only the Ti3 Al phase is finally produced, the following reactions remain. TiAl3 + 2Ti 3 = TiAl

(2.5)

TiAl + 2Ti = Ti3 Al

(2.6)

2.2 Weldability Analysis of Ceramic Materials

55

At the beginning of the diffusion reaction, Ti, Al diffuse each other. Because of the fast diffusion rate of Al, TiAl3 is formed first at the interface of Ti, Al, followed by TiAl at the interface of TiAl3 and Ti, and finally TiAl and Ti react to form Ti3 Al. Ti is a strong carbide forming element, so the free Ti in the interlayer reacts with the C in the Al2 O3 –TiC ceramic to form TiC:. Ti + C = TiC

(2.7)

Coexistence with TiC in Al2 O3 –TiC aggregates at the Al2 O3 –TiC/Ti interface. From the above analysis, it is clear that the reaction layer A mainly produces TiO, Ti3 Al and TiC phases. (ii) Within the Ti-Cu–Ti interlayer (reaction layer B). During the diffusion bonding of Al2 O3 –TiC ceramics and W18Cr4V steel with the Ti-Cu–Ti interlayer, the reaction in reaction layer B is mainly between Ti and Cu. Due to the small solubility of Ti in Cu, Ti exists mainly in the form of intermetallic compounds. According to the phase diagram of the Cu–Ti binary alloy, the Cu– Ti liquid phase starts to form at the Cu/Ti interface when the heating temperature reaches 985 °C. Within the Cu–Ti liquid phase region, Ti and Cu diffuse rapidly and are able to diffuse adequately. This system has the lowest formation free energy of CuTi and it is the easiest to be produced. The reaction products are also related to the relative concentration of Cu–Ti, and Cu and Ti can also produce CuTi in addition to CuTi2 . Excess Cu in the Cu–Ti liquid phase is squeezed out of the interface under pressure due to the pressure applied in the diffusion bonding process. Due to the rapid diffusion of C atoms, C in Al2 O3 –TiC ceramics and W18Cr4V steel quickly diffuses inside the Cu–Ti liquid phase and reacts with Ti to form TiC, which is diffusely distributed in the Cu–Ti liquid phase and exists in the Cu–Ti solid solution as TiC particles after solidification, enhancing the properties of the interfacial transition zone. The main phases in the reaction layer B are CuTi, CuTi2 and TiC. (iii) Ti side of Ti/W18Cr4V interface (reaction layer C). After the formation of Cu–Ti instantaneous liquid phase in the Ti-Cu–Ti interlayer, the C atoms in the W18Cr4V steel will rapidly diffuse to the Ti/W18Cr4V interface. Since Ti is a strong carbide forming element, Ti and C form TiC phase at the Ti/W18Cr4V interface. With the extension of holding time, TiC accumulates at the Ti/W18Cr4V interface and generates a continuous TiC layer. Fe and Ti have little mutual solubility and exist mainly as Fe-Ti intermetallic compounds.Ti in the Cu–Ti liquid phase diffuses into the W18Cr4V steel, while the W18Cr4V steel dissolves and diffuses into the Cu–Ti liquid phase.Ti and Fe react as follows. 2Fe + Ti = Fe2 Ti

(2.8)

56

2 Welding of Advanced Ceramic Materials

Fe + Ti = FeTi

(2.9)

FeTi, Fe2 Ti are formed and as the reaction proceeds, Fe2 Ti is converted to FeTi. At the Ti/W18Cr4V interface Ti reacts preferentially with C to form TiC, which hinders the diffusion of Ti into the W18Cr4V steel, so FeTi is present only in a very small area at the Ti/W18Cr4V interface.The reactive layer C on the Ti side at the Ti/W18Cr4V interface is mainly the TiC phase and a small amount of FeTi phase. (iv) Ti/W18Cr4V interface near the W18Cr4V steel side (reaction layer D) The high number of carbides in W18Cr4V HSS has a great influence on the properties of the steel. During diffusion bonding, C in W18Cr4V HSS diffuses to the Ti/W18Cr4V interface and reacts with Ti to form TiC, forming a decarburized layer on the W18Cr4V side with reduced C concentration, and this region mainly contains Fe, W and a small amount of C, generating Fe3 W3 C, making the carbide particles in W18Cr4V steel fine and the unreacted Fe in the α-Fe form is preserved. So the reacting layer D is mainly carbides such as Fe3 W3 C and α-Fe. The interfacial structure in Al2 O3 –TiC/W18Cr4V joint from Al2 O3 –TiC side to W18Cr4V side is in the following order: Al2 O3 –TiC/TiC + Ti3 Al + TiO/CuTi + CuTi2 + TiC/TiC + FeTi/Fe3 W3 C + α-Fe/W18Cr4V, as shown in Fig. 2.9. The formation of the phase structure in the interfacial transition zone is closely related to the diffusion bonding parameters. The boundaries of the reaction layers in the interfacial transition zone are not obvious and sometimes cross together. As seen in Fig. 2.7, Ti appears in almost all the interfacial reaction products, indicating that Ti is involved in all processes of the interfacial reaction. In the Al2 O3 –TiC/W18Cr4V diffusion bonding process, Ti is the main control element of the interfacial reaction.

Fig. 2.9 Phase structure of the transition zone at the Al2 O3 –TiC/W18Cr4V interface

2.2 Weldability Analysis of Ceramic Materials

57

2.2.3 Bond Strength at the Diffusion Interface The properties of diffusion welded joints vary greatly with different diffusion conditions and different interfacial reaction products. With the increase of heating temperature, the interface diffusion reaction is sufficient, which improves the strength of the joint. When diffusion bonding of alumina with steel using an aluminium foil of 0.5 mm thickness as an interlayer, the effect of heating temperature on the tensile strength of the joint is shown in Fig. 2.10. Excessive temperatures may cause changes in the properties of the ceramics or the appearance of a brittle phase that reduces the performance of the joint. In addition, the tensile strength of ceramic-to-metal diffusion welded joints is related to the melting point of the metal. In alumina-to-metal diffusion welded joints, as the melting point of the metal increases, the tensile strength of the joint increases. The tensile strength (σ b ) of ceramic-to-metal diffusion welded joints versus holding time (t) is given by. σb = B0 t 1/2

(2.10)

where B0 is a constant. However, there is an optimum value for the holding time at a certain heating temperature. The effect of holding time on the tensile strength of the joint in Al2 O3 /Al diffusion bonding joints is shown in Fig. 2.11a. When Nb was used as an interlayer during diffusion bonding of SiC and stainless steel, the NbSi2 phase with lower strength and quite different linear expansion coefficient with SiC appeared when the time was too long, and the shear strength of the joint was reduced as shown in Fig. 2.11b. If V 70

Tensile strength /MPa

60 50

calculated value

40

measured value

30 20 10 0 400

450

500

550

600

650

700

Bonding temperature / Fig. 2.10 Effect of heating temperature on the strength of alumina/steel diffusion welded joints

58

2 Welding of Advanced Ceramic Materials 250

70

200

Shear strength /MPa

Tensile strength /MPa

60 50 40 30 20 10 0

0

20

40

60

80

100

120

150

100

50

0

140

0

20

Holding time /min

40

60

80

100

120

Holding time /min

Fig. 2.11 Effect of holding time on joint strength

were used as an intermediate layer to diffuse join AlN with too long holding time, the shear strength of the joint was also reduced due to the appearance of the brittle phase of V5 Al8. The purpose of applying pressure in diffusion welding is to produce microscopic plastic deformation at the contact surface, reduce the surface roughness and damage the surface oxide film, increase the surface contact area, and provide conditions for atomic diffusion. In order to prevent large deformation of the ceramic and metal structure, the pressure applied in diffusion bonding is generally small (99.99

960.5

960.5

Au–Ni

Au 82.5, Ni 17.5

950

950

Cu–Ge

Ge 12, Ni 0.25, Cu Bal

850

965

Ag–Cu–Pd

Ag 65, Cu 20, Pd 852 15

898

Au–Cu

Au 80, Cu 20

889

889

Ag–Cu

Ag 50, Cu 50

779

850

Ag–Cu–Pd

Ag 58, Cu 32, Pd 824 10

852

Au–Ag–Cu

Au 60, Ag 20, Cu 835 20

845

Ag–Cu

Ag 72, Cu 28

779

779

Ag–Cu–In

Ag 63, Cu 27, In 10

685

710

(i)

The brazing material does not contain chemical elements with high saturated vapour pressure, such as Zn, Cd, Mg, etc., to avoid contamination of the electronics by these chemical elements during the brazing process or to cause dielectric leakage. (ii) The oxygen content of the brazing material must not exceed 0.001% to avoid the generation of water vapour when brazing in an inert atmosphere. (iii) The brazed joint should have good relaxation to minimize the thermal stresses caused by the difference in linear expansion coefficient between the ceramic and the metal. In order to minimize the welding stress when choosing the brazing filler metals for the connection between ceramics and metals, sometimes brazing materials with good plasticity and low yield strength, such as pure Ag, Au or Ag–Cu eutectic brazing materials, have to be used. The vitrification method uses capillary action to achieve the joining, this method does not add metal to brazing material but inorganic brazing material (vitreous body), such as oxide and fluoride brazing material. The glass phase formed by the melting of the oxide brazing material is able to penetrate into the ceramic, infiltrate the metal surface and finally form a joint. The oxide braze formulation in typical vitrification method is shown in Table 2.18. The glass body is not ductile enough to withstand the shrinkage of the ceramic after curing, and can only rely on formulating the composition so that its linear expansion coefficient is as close as possible to that of the ceramic. The practical application of this method is also quite strict.

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2 Welding of Advanced Ceramic Materials

Table 2.18 Typical glassy oxide braze formulations Series

Composition /%

Melting point linear expansion coefficient /°C /10–6 ·K−1

Al-Y-Si

Al2 O3 15, Y2 O3 65, SiO2 20

–

Al-Ca-Mg-Ba

Al2 O3 49, CaO 3, MgO 1550 11, BaO 4 1410 Al2 O3 45, CaO 36.4, MgO 4.7, BaO 13.9

– 8.8

Al-Ca-Ba-B

Al2 O3 46, CaO 36, BaO 16, B2 O3 2

9.4 ~ 9.8

Al-Ca-Ba-Sr

Al2 O3 44 ~ 50, CaO 35 ~ 1500 (1310) 40, BaO 12 ~ 16, SrO 1.5 ~ 1500 5 Al2 O3 40, CaO 33, BaO 15, SrO 10

7.7 ~ 9.1 9.5

Al-Ca-Ta-Y

Al2 O3 45, CaO 49, Ta2 O3 3, Y2 O3 3

(1380)

7.5 ~ 8.5

Al-Ca-Mg-Ba-Y

Al2 O3 40 ~ 50, CaO 30 ~ 40, MgO 10 ~ 20 BaO 3 ~ 8, Y2 O3 0.5 ~ 5

1480 ~ 1560

6.7 ~ 7.6

Zn-B-Si-Al-Li

ZnO 29 ~ 57, B2 O3 19 ~ 56, SiO2 4 ~ 26, Li2 O 3 ~ 5, Al2 O3 0 ~ 6

(1000)

4.9

Si-Ba-Al-Li-Co-P

SiO2 55 ~ 65, BaO 25 ~ 32, (950 ~ 1100) Al2 O3 0 ~ 5, Li2 O 6 ~ 11, CaO 0.5 ~ 1, P2 O5 1.5 ~ 3.5

(1320)

Si-Al-K-Na-Ba-Sr-Ca SiO2 43 ~ 68, Al2 O3 3 ~ 6, (1000) K2 O 8 ~ 9, Na2 O 5 ~ 6, BaO 2 ~ 4, SrO 5 ~ 7, CaO 2 ~ 4, with a small amount of Li2 O、MgO、TiO2 、B2 O3

7.6 ~ 8.2

10.4

8.5 ~ 9.3

Note The data in parentheses are reference temperatures

Solder with different melting points and linear expansion coefficients can be obtained by adjusting the solder formula, so that it can be used for the joining between different ceramics and metals. This vitreous intermediate material is actually the bonding phase between Si3 N4 ceramic grains (e.g. Al2 O3 , Y2 O3 , MgO etc.) and the impurity SiO2 , which is present at sintering. The bonding takes place at temperatures above 1530 °C (corresponding to the eutectic point of Y-Si-Al-O-N), without pressure and usually protected by nitrogen.

2.3 Brazing of Ceramic to Metal Joints Powder Making ingredient paste Biscuit Porce Clean -ting -lain Metal Clean parts Solder Clean

69 Paste prepare

Assem -ble

Metallization

Polish

Brazing

Inspect

Ni plating

Sintering

Finished product

Fig. 2.17 Process flow for Mo-Mn method ceramic metallization brazed joining

(3) Ceramic metallization brazing process The process flow of ceramic metallization brazing connection is shown in Fig. 2.17. The key points of ceramic metallization brazing process are: Taking Mo-Mn metallization as an example, the process flow of ceramic metallization brazed joints by the Mo-Mn metallization method is shown in Fig. 2.17. The key points of the ceramic metallization brazing process are as follows: The preparation and coating process for metallized pastes is as follows. ➀ Cleaning of parts Ceramic parts can be cleaned with detergent in an ultrasonic cleaner, then cleaned with deionized water and dried. Metal parts, on the other hand, should be cleaned by alkaline washing and acid washing to remove oil and oxidation film on the metal surface, and then cleaned with flowing water and dried. The cleaned parts should immediately go to the next process and should not be touched with bare hands in the middle process. ➁ Paste preparation The powder of various raw materials is weighed in proportion, and add appropriate amounts of nitrocellulose solution, butyl acetate, diethyl oxalate, etc. This is an important process of ceramic metallization, the paste is mostly composed of pure metal powder with an appropriate amount of metal oxides. The particle size of powder is between 1 to 5 μm, mixed into a paste with organic binder, and then evenly coated on the surface of the ceramic to be metallized. The thickness of the coating is about 30 to 60 μm. ➂ Ceramic metallization The coated ceramic parts are placed in a hydrogen furnace and hold at a temperature of 1300–1500 °C for 0.5–1 h. ➃ Nickel plating The metallized layer is mostly Mo-Mn layer, which is difficult to infiltrate with brazing material, and must be plated with another layer of nickel with 4–5 μm thickness. ➄ Assembling the joint Assemble the treated metal and ceramic parts together, and prepare the brazing material at the joints.

70

2 Welding of Advanced Ceramic Materials

Fig. 2.18 Brazing of ceramic probe component 1—Bronze; 2—Ceramic; 3—Stainless steel

➅ Brazing Brazing is carried out in an inert atmosphere or in a vacuum furnace, and the brazing temperature is determined by the brazing material. Neither the heating nor the cooling rate should be too fast during brazing to prevent the ceramic parts from cracking. ➆ Inspection For some special requirements of ceramic sealing parts, such as vacuum devices or electrical parts, tests such as leakage, thermal shock, thermal baking and insulation strength should be carried out. An example of the application of the ceramic metallization method of brazing is shown below. Figure 2.18 shows a probe component used in an oil testing instrument, the material is purple copper and stainless steel. The components are isolated from each other with Al2 O3 ceramic, which acts as an insulator and requires a leak-free seal after brazing. The brazing process uses the Mo-Mn method to metalize the inside of the hole at one end of the Al2 O3 ceramic tube and the outer surface of the tube to be welded, and then a nickel layer with a thickness of 35 μm is plated on the outside of the metalized layer. BAg72Cu was used as the brazing material, a bright and dense joint was obtained by holding the joint for 5 min under the vacuum of 1.33 × 10–2 Pa and a brazing temperature of 850 °C.

2.3.3 Active Metallization Brazing of Ceramics to Metals Transition group metals (e.g. Ti, Zr, Nb, etc.) are highly chemically active and these elements have a large affinity for oxides, silicates, etc. They can form reactive layers on the ceramic surface through chemical reactions. After the addition of such active metals to the brazing materials such as Au, Ag, Cu, Ni, etc., the so-called active brazing solders are formed. The active brazing solders is very easy to react chemically with the ceramic in the liquid state to form a ceramic-to-metal joint.

2.3 Brazing of Ceramic to Metal Joints

71

The reaction layer mainly consists of metal and ceramic composites (exhibiting the same microstructure as the metal and can be wetted by the molten metal), so as to achieve the purpose of joining with the metal. The active metal has strong chemical activity and the protection of the active elements during brazing is important. Once these elements have been oxidized they can no longer react with the ceramic. Therefore active metalization brazing method is generally carried out in a vacuum or inert protective atmosphere above 10–2 Pa, and the brazing joint is completed in one pass. (1) Active brazing material Active solder usually uses Ti as the active element and can be applied to braze oxide ceramics, non-oxide ceramics and various inorganic dielectric materials. They are developing rapidly because they are brazed directly to ceramics with the active metal and the process is simple. A comparison of several commonly used active metalization brazing methods is shown in Table 2.19. The high-temperature active brazing materials used for direct brazing of ceramics and metals are shown in Table 2.20. The binary brazing materials are mainly Ti-Cu and Ti-Ni, which have a low vapor pressure of less than 1.33 × 10–3 Pa at 700 °C and can be used in the range of 1200 to 1800 °C. Ternary brazing materials are Ti-CuBe or Ti-V-Cr, of which 49Ti-49Cu-2Be has similar corrosion resistance to stainless steel and lower vapor pressure, and is used in leak-proof and oxidation-proof vacuum sealing joints. Ti-Zr-Ta brazing materials without Cr can also be brazed directly to MgO and Al2 O3 ceramics, and the joints obtained with this solder are able to work at temperatures higher than 1000 °C. The Ag–Cu–Ti system brazing material developed in China is capable of directly brazing ceramics with oxygen-free copper, and the shear strength of the joint can reach 70 MPa. (2) Active brazing joining process Taking active metal Ti–Ag–Cu method as an example, the process flow of active brazing connection of ceramics to metals is shown in Fig. 2.19. Figure 2.19 process flow for active brazing of ceramic to metal connections. Key points of active metallization brazing process. ➀ Parts cleaning Ceramic parts can be cleaned in the ultrasonic cleaning machine, and metal parts are cleaned by alkaline and acid washing to remove oil and oxidation film on the metal surface. The cleaned parts will immediately go to the next process. ➁ Paste-making The purity of titanium powder used for paste-making should be above 99.7%, and the particle size should be within the range of 270–360 mesh. When making the paste, take cellulose solution which is half weight of the titanium powder, adding a small amount of diethyl oxalate dilution, and then mixed them into a paste.

3~5

Ti71.5ni28.5 foil with thickness of 10 ~ 20 m was used as solder for brazing

990 ± 10

Ti-Ni

Metal material

High alumina, magnesium olivine ceramics

Ti

High alumina, sapphire, Cu,Ti,Nb transparent alumina, magnesium olivine, glass ceramics, mica, graphite and non-oxide ceramics

Brazing temperature/°C Holding time/min Ceramic material 3~5

Solder prepare method

Ag–Cu–Ti Ti powder with 850 ~ 880 thickness of 20 ~ 40 m was precoated on the ceramic surface, and then Ag69Cu26Ti5 solder with thickness of 0.2 mm was used for brazing

Solders

Table 2.19 Comparison of several commonly used active metallization brazing methods

(continued)

High brazing temperature, low vapor pressure, good wettability for ceramics, especially suitable for matching brazing Ti and magnesium olivine ceramics. The disadvantage is that the brazing temperature range is narrow, and the surface cleaning of parts is strict

It has good wettability to ceramics, good air tightness of joints, and is widely used. It is commonly used for matching brazing of large parts and brazing of soft metals and high-strength ceramics. The disadvantage is that the solder contains a large amount of Ag, the vapor pressure is high, and it is easy to deposit on the ceramic surface, which reduces the insulation performance

Characteristics and Application

72 2 Welding of Advanced Ceramic Materials

High alumina magnesium olivine and non-oxide ceramics

900 ~ 1000

Ti 25% ~ 30%, Cu in balance, Ti(Cu) foil or powder with above composition were used as solder and brazing

Cu–Ti

2~5

Brazing temperature/°C Holding time/min Ceramic material

Solder prepare method

Solders

Table 2.19 (continued) Characteristics and Application

Cu,Ti,Ta,Nb,Ni-Cu High brazing temperature, low steam pressure, good wettability of ceramics, alloy is brittle and hard, suitable for matching brazing or high strength ceramic brazing

Metal material

2.3 Brazing of Ceramic to Metal Joints 73

74

2 Welding of Advanced Ceramic Materials

Table 2.20 High temperature reactive brazing materials for direct brazing of ceramics to metals Solders

Melting point/°C

Brazing temperature/°C

Application and joint performance

92Ti–8Cu

790

820 ~ 900

Ceramic–metal joining

75Ti–25Cu

870

900 ~ 950

Ceramic–metal joining

72Ti–28Ni

942

1140

Ceramic–ceramic, ceramic–graphite, ceramic–metal joining

68Ti–28Ag–4Be

–

1040

Ceramic–metal joining

54Ti–25Cr–21 V

–

1550 ~ 1650

Ceramic–ceramic, ceramic–graphite, ceramic–metal joining

50Ti–50Cu

960

980 ~ 1050

Ceramic–metal joining

50Ti–50Cu (at.%)

1210 ~ 1310

1300 ~ 1500

Ceramic–sapphire, ceramic–lithium joining

49Ti–49Cu–2Be

–

980

Ceramic–metal joining

48Ti–48Zr–4Be

–

1050

Ceramic–metal joining

47.5Ti–47.5Zr–5Ta

–

1650 ~ 2100

Ceramic–tantalum joining

7Ti–93 (BAg72Cu)

779

820 ~ 850

Ceramic–titanium joining

5Ti–68Cu–26Ag

779

820 ~ 850

Ceramic–titanium joining

100Ge

937

1180

Self bonding silicon carbide–metal (σb = 400 MPa)

85Nb–15Ni

–

1500 ~ 1675

Ceramic–niobium (σb = 145 MPa)

75Zr–19Nb–6Be

–

1050

Ceramic–metal joining

56Zr–28 V–16Ti

–

1250

Ceramic–metal joining

83Ni–17Fe

–

1500 ~ 1675

Ceramic–tantalum (σb = 140 MPa)

66Ag–27Cu–7Ti

779

820 ~ 850

Ceramic–titanium joining

Powder ingredient Porce -lain

Biscuit -ting

Metal parts Solder

Making paste Clean Clean

Paste prepare Assem -ble

Clean

Fig. 2.19 Active brazing joining process of ceramic and metal

Brazing

Inspect

Finished product

2.3 Brazing of Ceramic to Metal Joints

75

➂ Application of the paste The active brazing paste is applied uniformly to the brazing surface of the ceramic by means of a brush or other spraying method. The coating should be uniform and the thickness is generally around 25 to 40 μm. ➃ Assembly The paste on the ceramic surface is dried and assembled with the metal parts and BAg72Cu brazing material. ➄ Brazing Brazing is performed in vacuum or inert atmosphere. When the vacuum reaches 5 × 10–3 Pa, gradually increase the temperature to 779°C to make the brazing material melt, then increase the temperature to 820–840 °C, holding for 3–5 min (too high temperature or too long holding time will make the active elements react strongly with the ceramic parts, causing the brazing seam to be loosely structure and form air leakage) and then cool down. In the heating or cooling process, pay attention to the heating and cooling speed to avoid cracking of the ceramic due to heating and cooling too fast. ➅ Inspection Baking resistance inspection and air tightness inspection should be carried out for brazed parts. For vacuum devices or electrical parts, tests such as air leakage, thermal shock, thermal baking and electrical insulation strength should be carried out.

2.3.4 Examples of Ceramic-to-Metal Brazing (1) Several application examples Ceramic-to-metal joining structures are widely used in the electronics industry, and applications are being developed in mechanical, metallurgical, and energy fields. Some examples of applications are shown in Fig. 2.20. (1) Automotive engine supercharger rotor To improve the performance of automotive engines and to save fuel, ceramic–metal composite parts have received much attention. Si3 N4 ceramics have good prospects for manufacturing automotive engine supercharger rotors due to their low density, good high temperature strength and resistance to wear without lubrication. Such ceramic-steel composite rotors are about 40% lighter in mass than conventional allmetal rotors and resistant to temperatures up to 1000 °C. These properties improve the acceleration performance and combustion efficiency of turbines and reduce exhaust emissions. Such composite rotors have also found application in heavy-duty diesel engines. The structure of this automotive engine supercharger rotor is shown in Fig. 2.21a, which is composite structure of Si3 N4 ceramic turbine and metal shaft, connected

76

2 Welding of Advanced Ceramic Materials

Kovar cylinder

Alumina ceramic sheet Kovar Transition tube Oxygen-free copper needle

Ceramic shaft

Solder Metal shaft

Fig. 2.20 Example of application of ceramic to metal brazed structures a Vacuum switch tube shell b sleeve sealing transition needle sealing mandrel c Composite structure of inner and outer sleeve seal and transition needle seal d Joining structure of ceramic turbine shaft and metal shaft

as a whole by adding an interlayer of active brazing material and a sleeve. The key to forming this ceramic and metal composite structure is as the following two points: ➀ The use of a multilayer buffer layer consisting of Ni-W alloy and Ni of 2 to 4 mm thickness, which reduces the maximum stress in the ceramic from 1250 to 210 MPa when directly connected. ➁ The active brazing material was chosen to wet the surface of Si3 N4 ceramics well for brazing without metallization. The vacuum degree for brazing is 3 × 10–2 Pa. (2) Ceramic and metal rocker An automotive company has introduced a ceramic and metal composite rocker. This rocker is partially made of Si3 N4 ceramic, which reduces wear by 5 to 10 times compared to all-metal parts, thus extending the maintenance period. This rocker is made by joining Si3 N4 ceramic inserts to a steel substrate through an interlayer. Si3 N4 ceramic inserts are pre-coated with a titanium layer firstly and

2.3 Brazing of Ceramic to Metal Joints Solder

77 Raised the amount ceramic chip Solder

Shaft sleeve

Steel thimble thickness

Metal Tungsten alloy Steel Multilayer buffer

Fig. 2.21 Example of a ceramic and metal composite structure a Ceramic and metal composite supercharger rotor b Si3 N4 and Steel Composite Ceramic Tappet

then brazed to the steel substrate with BAg72Cu brazing material at 850 °C in inert atmosphere. Due to the low service temperature (mainly wear resistance), a Cu sheet with a thickness of 0.5 mm is sufficient for the interlayer to meet the process requirements. (3) Ceramic and metal tappet The tappet and cam are an important friction pair in the engine valve train, and the contact surface of the tappet is subjected to intense friction during service. The composite ceramic tappet made of Si3 N4 ceramic has superior wear resistance compared to the cold-excited cast iron and hard dense cast iron tappets commonly used today. The structure of Si3 N4 ceramic composite tappet with steel is shown schematically in Fig. 2.21b. Si3 N4 ceramic is jointed with steel sleeve by brazing technique. This Si3 N4 ceramic and steel composite tappet can be used in heavy-duty diesel engines and has good application prospects. (2) Notes for brazed joint design. ➀ Reasonable selection of brazing materials matching. Select ceramics with similar linear expansion coefficient to match each other with metals, such as Ti with Mg olivine ceramics and Ni with 95% Al O23 ceramics, the linear expansion coefficient is basically the same in the range of room temperature to 800 °C. Use the plasticity of the metal to reduce the brazing stress, such as brazing 95% Al2 O3 ceramic using oxygen-free copper. Although the linear expansion coefficient of the metal and ceramic is very different, but due to the full use of the plasticity and ductility of soft metals good joints can still be obtained.

78

2 Welding of Advanced Ceramic Materials

Fig. 2.22 Typical form of flexural brazed joint

The selection of high strength and high thermal conductivity ceramics, such as BeO and AlN, can reduce the thermal stress at the brazed joint and improve the bond strength of brazing joint. ➁ Stress reduction by elastic deformation of metal parts. A “flexural brazed joint structure” is designed to relieve stresses by using the elastic deformation of thin walls of non-brazed part of metal parts. A typical flexural brazing joint is shown in Fig. 2.22. ➂ Avoid stress concentration Ceramic parts should be designed to avoid sharp corners or disparity in thickness, and try to use rounded transition. Change the shape of the end of the metal parts when casing and sealing, so that the metal end of the seal is thinned, which can increase plasticity and reduce stress concentration. Control the heating temperature of brazed parts to prevent the production of weld beading. The linear expansion coefficient of brazing material is generally large, and if brazing material builds up, it can cause local stress concentration and lead to ceramic cracking. ➃ Pay attention to the influence of brazing materials. Try to use brazing materials with low strength and good plasticity such as Ag–Cu eutectic, pure Ag, Cu, Au, etc. to maximize the stress release. The braze layer is as thin as possible under the premise of ensuring the seal. Select the appropriate

2.4 Diffusion Bonding of Ceramics to Metals

79

length of the solder foot. The length of the solder foot has a great influence on the strength of the joint when sealing, generally 0.3 ~ 0.6 mm is appropriate.

2.4 Diffusion Bonding of Ceramics to Metals 2.4.1 Characteristics of Ceramic-to-Metal Diffusion Bonding Diffusion bonding is a common method for ceramic/metal joints. It is a process in which the surfaces to be joined are in contact with each other at a certain temperature and pressure, and the physical contact of the surfaces to be joined is expanded by causing local plastic deformation of the contact surfaces, or by the transient liquid phase generated on the surfaces to be joined, and then combined with mutual diffusion between the atoms of the interface to form an overall reliable joint. This joining method is characterized by stable joint quality, high bonding strength, good high temperature performance and corrosion resistance of the joint. (1) Direct diffusion bonding This method requires a very flat and clean surface of the jointed part, which is subjected to high temperature and pressure to achieve atomic contact and diffusion migration of atoms at the joining interface. (2) Indirect diffusion bonding This method is the most common diffusion joining method used in ceramic welding. The joining is completed at a certain temperature and pressure by adding a metal intermediate layer with good plasticity between the parts to be joined. Indirect diffusion bonding can reduce the joining temperature, avoid the coarse microstructure of the connected parts, reduce the problems caused by the mismatch of thermophysical properties when joining different materials. It is an effective method of connecting ceramics to metals. Indirect diffusion bonding is divided into the following two kinds: ➀ The ceramic, metal and intermediate layer all remain in a solid state and non-melting. Only the contact area between the ceramic and the metal gradually expands by heating and pressurizing. Surface diffusion and volume diffusion of certain components occur, eliminating interface pores, causing the interface to move and eventually forming a reliable joint. ➁ The intermediate layer melts instantaneously, and a trace liquid phase appears instantaneously in the joint area during diffusion joining, also known as transient liquid phase diffusion welding (TLP). This method combines the advantages of brazing and solid phase diffusion bonding, using the formation of a low melting point eutectic between the base material to be welded and the interlayer at a certain temperature. Isothermal solidification and accelerated diffusion processes occur through the diffusion of solute atoms to form a diffusion bonded joint with uniform microstructure.

80

2 Welding of Advanced Ceramic Materials

Transient liquid-phase diffusion bonding can be applied to ceramic-to-ceramic or ceramic-to-metal joints, and the effects of the instantaneous liquid-phase diffusion joint formation process, interlayer design, effects of temperature and pressure on the joint performance, and the bonding mechanism can be investigated in depth. Trace liquid phases help to improve the interfacial contact state, can reduce the bonding temperature and allow the use of lower diffusion pressures. There are two main methods to obtain trace liquid phases. a. Use of eutectic reactions The liquid phase diffusion bonding (called eutectic reaction diffusion bonding) is carried out by taking advantage of the possible formation of low melting point eutectic between certain dissimilar materials. This method requires that once the liquid phase is formed it should be immediately cooled down to solidify so as not to continue to generate excess liquid phase. Thus, the temperature and holding time should be strictly controlled. When applying the principle of eutectic reaction diffusion bonding to the addition of an interlayer diffusion joining, the total amount of liquid phase can be controlled by the thickness of the interlayer, and this method is called instantaneous liquid phase diffusion bonding (or transition liquid phase diffusion bonding). b. Addition of special brazing materials Using materials that are close to the base material composition but contain a small amount of elements that can lower the melting point and can diffuse rapidly in the base material(such as B, Si, Be, etc.) as brazing solder. This brazing material is added as an interlayer in the form of foil or coating. Compared with conventional brazing, the thickness of this braze layer is thinner and the braze solidification is finished in an isothermal state, whereas in conventional brazing the braze is solidified during the cooling process. Diffusion bonding has a wide range of applications and reliable quality control in the welding of ceramics to metals. The main diffusion bonding processes for ceramic materials includes: ➀ Direct diffusion bonding of homogeneous ceramic materials. ➁ Diffusion bonding of the same ceramic material with another thin layer of material. ➂ Direct diffusion bonding of dissimilar ceramic materials. ➃ Diffusion bonding of heterogeneous ceramic materials with a third thin layer of material. When welding ceramics to metals, diffusion bonding with a filler interlayer and eutectic reaction diffusion bonding are often used. The main advantages of diffusion bonding of ceramic materials are: high joint strength, easy dimensional control, and suitability for joining dissimilar materials. The main disadvantages are the high diffusion temperature, the long bonding time and the process under vacuum, the large equipment investment in one-time, and the restricted size and shape of the workpiece.

2.4 Diffusion Bonding of Ceramics to Metals

81

Diffusion bonding of ceramics to metals can be carried out both in vacuum and in an inert atmosphere. Mutual chemical interactions are more likely to occur when there is an active film on the metal surface. Therefore, filling the welding vacuum chamber with a reductive and active medium (which maintains a thin active film on the metal surface) results in a stronger bond and higher strength of the diffusion bonded joint. Qualified joint strengths are obtained by diffusion welding temperature reaches 900 °C between alumina ceramics and oxygen-free copper. Higher strength value can only be obtained at welding temperatures of 1030 to 1050 °C, because copper has great plasticity and is prone to deform under pressure, which increases the actual contact surface. Factors affecting the strength of diffusion welded joints are the heating temperature, holding time, pressure, ambient medium, surface condition of the joining surface and the matching of chemical reactions and physical properties (such as the linear expansion coefficient, etc.) between the joined materials.

2.4.2 Process Parameters for Diffusion Bonding In solid-phase diffusion bonding, the bonding temperature, pressure, holding time and surface condition of the weld are the main factors affecting the quality of diffusion bonding. The bonding of the interface in the solid-phase diffusion bonding is achieved by plastic deformation, diffusion and creep mechanisms, and its bonding temperature is high, and the ceramic/metal diffusion bonding temperature is usually 0.8 ~ 0.9 times of the metal melting point. Due to the mismatch between the linear expansion coefficient and elastic modulus of ceramics and metals, it is easy to generate large stresses near the interface, which makes it difficult to achieve direct diffusion joints. In order to relieve the residual stress in the ceramic–metal joint and to control the interfacial reaction, inhibit or change the interfacial reaction products to improve the joint performance, diffusion bonding with an interlayer is often used. (1) Heating temperature The heating temperature has the most significant effect on the diffusion process. When joining metal and ceramic, the temperature sometimes reaches more than 90% of the melting point of metal. In solid-phase diffusion bonding, the chemical reaction caused by the mutual diffusion of elements can form a sufficient interfacial bond. The thickness of the reacting layer (X) can be estimated by the following equation: X = K 0 t n exp (−Q/RT )

(2.11)

where K 0 is a constant; t is the bonding time(s); n is the time index; Q is the diffusion activation energy (J/mol), depending on the diffusion mechanism; T is the thermodynamic temperature (K); and R is the gas constant, R = 8.314 J/(K·mol).

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2 Welding of Advanced Ceramic Materials

The same trend is observed for the effect of heating temperature on the strength of the joint, and the effect of temperature on the tensile strength of the joint (σ b ), obtained from tensile tests, can be expressed by the following equation: ( ) σb = B0 exp −Q app /RT

(2.12)

where B0 is a constant; Qapp is the apparent activation energy, which can be the sum of various activation energies. The increase in heating temperature increases the strength of the joint, but the increase in temperature may cause a change in the properties of the ceramic or the appearance of brittle phase that may reduce the performance of the joint. The tensile strength of ceramic-to-metal diffusion welded joints is related to the melting point of the metal. In alumina-to-metal diffusion welded joints, the tensile strength of the joint increases as the melting point of the metal increases. For example, the interfacial structure and bending strength of diffusion joints with Si3 N4 ceramics joined with aluminum as an interlayer vary considerably at different heating temperatures. The effect of heating temperature on the bending strength of Si3 N4 /Al/Si3 N4 diffusion joints is shown in Fig. 2.23. It can be seen that at low temperature joints, due to the residual Al interlayer at the joint interface, the bending strength of diffusion joints decreases sharply with increasing temperature, mainly because the properties of Al affect the joint strength. The bending strength of the joints treated at 1970 K increased with increasing heating temperature (Fig. 2.23b) due to the formation of AlN ceramics from the residual Al at high temperature. AlN ceramics is stronger than Al, and the denser AlN and AlSi aggregation band increased the joint strength. (2) Holding time The relationship between the thickness of the reaction layer and the holding time in SiC/Nb diffusion bonded joint is shown in Fig. 2.24.

Fig. 2.23 Si3 N4 /Al/Si3 N4 Diffusion joint organization and flexural strength a changes in interface structure b influence of temperature on bending strength

2.4 Diffusion Bonding of Ceramics to Metals

83

Reaction layer thickness /μm

100

80

60

40

20

0

0

50

100

150

200

1/2

Holding time t /s Fig. 2.24 SiC/Nb diffusion weld head reaction layer thickness versus holding time t (from top to bottom: 1–1773 k, 2–1673 k, 3–1573 k, 4–1473 k)

Fig. 2.25 Thickness of SiC/Ti reaction layer as a function of heating temperature and time

Reaction layer thickness /m

When other conditions are the same, the thickness of the diffusion weld reaction layer increases as the heating temperature and joining time increase, as shown in Fig. 2.25. The same trend is observed for the effect of holding time on the strength of diffusion welded joints. The relationship between tensile strength (σ b ) and holding time (t) is: σ b = B0 t 1/2 , where B0 is a constant.

Joining time

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2 Welding of Advanced Ceramic Materials

Table 2.21 Interfacial reaction products in different types of ceramic-to-metal joints Joint assembly Interfacial reaction product Joint assembly

Interfacial reaction product

Al2 O3 /Cu

CuAlO2 , CuAl2 O4

Si3 N4 /Al

AlN

Al2 O3 /Ti

NiO·Al2 O3 , NiO·SiAl2 O3

Si3 N4 /Ni

Ni3 Si, Ni(Si)

SiC/Nb

Nb5 Si3 , NbSi2 , Nb2 C, Nb5 Si3 Cx , NbC

Si3 N4 /Fe–Cr alloy

Fe3 Si, Fe4 N, Cr2 N, CrN, Fe x N

SiC-Ni

Ni2 Si

AlN/V

V(Al), V2 N, V5 Al8 , V3 Al

SiC/Ti

Ti5 Si3 , Ti3 SiC2 , TiC

ZrO2 /N-Cr-(O)/ZrO2 NiO1-x Cr2 O3-y ZrO2-z , 0 < x,y,z < 1

(3) Pressure In order to prevent deformation of the component, the pressure applied for ceramic-tometal diffusion bonding is generally less than 100 MPa.When solid-phase diffusion bonding ceramic to metal, the ceramic-to-metal interface reacts to form compounds, and the type of formed compound is related to the joining conditions (e.g., temperature, surface condition, type and content of impurities, etc.). The possible products of interfacial reactions in different types of ceramic–metal joints are shown in Table 2.21. Joint performance varies considerably with different diffusion conditions and different interfacial reaction products. In general, the joint strength of vacuum diffusion bonding is higher than that of joints bonded in argon and air. The use of an intermediate layer in diffusion welding of ceramics and metals not only reduces the residual stresses generated in the joint, but also lowers the heating temperature, reduces the pressure and shortens the holding time, promotes diffusion and removes impurity elements. The selection of the intermediate layer is critical, and improper selection can cause deterioration of the joint performance. For example, the bending strength of the joint may be reduced due to the formation of brittle reactants as a result of intense chemical reactions, or the residual stresses may be increased due to a mismatch in the linear expansion coefficient, or the corrosion resistance of the joint may be reduced, even leading to cracks and fractures. The interlayer can be incorporated in different forms, usually as a powder, in foil form or by metallization. The process parameters for diffusion welding of various ceramic material combinations and their properties are shown in Table 2.22. Al2 O3 , SiC, Si3 N4 and WC ceramics were researched and developed earlier and developed more maturely. While AlN, ZrO2 ceramics have been developed relatively late. The hardness and strength of ceramics are high and not easily deformed, so the diffusion bonding of ceramics and metal requires a flat and clean surface to be connected, and the diffusion bonding must be under high pressure (pressure 0.1 to 15 MPa), high temperature (usually 0.8 to 0.9 of the melting point of the metal T m ), and much longer welding time than other welding methods. Ceramic materials commonly used in diffusion bonding of ceramics to metals are alumina ceramics and

120 240

1600

1550

600

1025 ~ 1050

1050

800

1375

1450

1450

Al2 O3 + Nb

Al2 O3 + Pt

Al2 O3 + Al

Al2 O3 + Cu

94%Al2 O3 + Cu

Al2 O3 + Cu4 Ti

Al2 O3 + Fe

Al2 O3 + mild steel

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1025

1000

Al2 O3 /Cu/Al2 O3

1375

900

727 ~ 877

1100

Si3 N4 + Invar

Si3 N4 + Nimonic80A 6 ~ 60

7

120

2

60

1250

Al2 O3 /Ag/Al2 O3

30

240

1350

Al2 O3 /Fe/Al2 O3

Al2 O3 /Ni/Al2 O3

15

1650

Al2 O3 /Pt/Al2 O3

15

1100

Al2 O3 + Cr

30

625

Al2 O3 + high alloy steel

1.7 ~ 6

20

50 ~ 60

155

1.7 ~ 5

1.7 ~ 20

60

20

1350

Al2 O3 + Ni

Hold time /min

Heat temperature /°C

Bonding materials

0 ~ 50

0 ~ 0.15

6

50

15 ~ 20

50

6

50

0.8

120

50

0.2 mm Ni

25μ m V

–

–

–

–

Al-Si/Al/Al-Si

–

–

–

Fe–Ni-Cr

Al-Si

Al

None

None

ZrO2

10 ~ 20μ m Al

Cu,Ni,Kovar

Interlayer and thickness

Vacuum

Vacuum

Vacuum

Vacuum

Vacuum

Vacuum

Vacuum

Vacuum

Vacuum

Vacuum

Vacuum

Vacuum

Vacuum

0.1 MPa nitrogen

1 MPa nitrogen

Vacuum

Air

–

Environmental atmosphere

57 (S) (continued)

120 (S)

0 ~ 40

125

187室温, >100(800°C)

87 (S)

200 (B)

95 (S)

20 (S)

100 (S)

>90 (A)

50b (A)

208b (A)

220(A)室温, 135(A)1000°C

380(A)室温, 230(A)1000°C

175 (B)

320 ~ 490(B)

–

Strength /MPa

86 2 Welding of Advanced Ceramic Materials

60 10

10 ~ 15

10

10

6

Pressure /MPa

Ag 25μ m

0.1mmCu

0.1mmNi

–

Interlayer and thickness

Vacuum

Vacuum

Vacuum

Vacuum

Environmental atmosphere

–

240 (A)

150 (A)

97 (T)

Strength /MPa

Note The letters in brackets after the strength value represent various performance test methods, A represents four-point bending test, B represents three-point bending test, T represents tensile test, S represents shear test; the superscript b represents the maximum value

250 ~ 450

900

BeO + Cu

60

1100

ZrO2 + ZrO2

120

1000

ZrO2 /Cu/ZrO2

Hold time /min

Heat temperature /°C

Bonding materials

Table 2.22 (continued)

2.4 Diffusion Bonding of Ceramics to Metals

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2 Welding of Advanced Ceramic Materials

zirconia ceramics. Metals welded to such ceramics are usually copper (oxygen-free copper), titanium (TA1), titanium-tantalum alloy (Ti-5Ta), etc. For example, alumina ceramics have high hardness and low plasticity, and will retain this property during diffusion bonding. Even if a glassy phase is present within the alumina ceramic (mostly distributed around the corundum grains), the ceramic will not show creep until it is heated above 1100 to 1300 °C. The actual contact between the ceramic and most metals during diffusion welding is formed during the local plastic deformation of the metal. Table 2.23 lists the matched combinations of Al2 O3 ceramics to different metals, diffusion welding conditions, and joint strength. When there is difficulty in joining the ceramic and metal directly by diffusion bonding, the method of adding an interlayer can be used, and the plastic deformation of the metal interlayer can reduce the processing accuracy on the ceramic surface. For example, adding Cu foil with a thickness of 20 μm as a transition layer between the ceramic and Fe–Ni–Co alloy, a diffusion bonded joint with a tensile strength of 72 MPa can be obtained under the process conditions of heating temperature1050°C, holding time 10 min and pressure15 MPa. Table 2.23 Diffusion welding conditions and joint strength of various Al2 O3 ceramics with different metals Ceramic–metal combination 95% alumina ceramics (contain MnO)

72% alumina ceramics

99.7% alumina ceramics

94% alumina ceramics

Atmosphere

Heating temperature /°C

Bending strength /MPa

Fe–Ni-Co

H2 (vacuum)

1200

100 (120)

Stainless steel

H2 (vacuum)

1200

100 (200)

Ti

Vacuum

1100

140

Ti-Mo

Vacuum

1100

100

Fe–Ni-Co

H2

1200

100

Stainless steel

H2 (vacuum)

1200

115

Ti

Vacuum

1100

125

Ni

Vacuum

1200

130

Stainless steel

Vacuum

1250 ~ 1300

180 ~ 200

Ni

Vacuum

1250 ~ 1300

150 ~ 180

Ti

Vacuum

1250 ~ 1300

160

Fe–Ni-Co

Vacuum

1250 ~ 1300

110 ~ 130

Fe–Ni alloy

Vacuum

1250 ~ 1300

50 ~ 80

Nb

Vacuum

1250 ~ 1300

70

Ni–Cr

H2 (vacuum)

1250 ~ 1300

100

Pd

H2 (vacuum)

1250 ~ 1300

160

3# steel

H2 (vacuum)

1250 ~ 1300

50

Stainless steel

H2

1250 ~ 1300

30

Note 1. The vacuum degree is 10–2 ~ 10–3 Pa 2. Holding Time is 15 to 20 min

2.4 Diffusion Bonding of Ceramics to Metals

89

The inter transition layer can be achieved by using metal foil directly, or by vacuum evaporation, ion sputtering, chemical vapor deposition (CVD), spraying, electroplating, etc. Diffusion bonding can also be achieved using the sintered metallization method, reactive metallization method, metal powder or brazing material as described earlier. The diffusion bonding process is not only used for welding metals to ceramics, but also for welding microcrystalline glass, semiconductor ceramics, quartz, graphite, etc. to metals. The process parameters for diffusion bonding of inorganic non-metallic materials to metals are shown in Table 2.24. Table 2.25 lists the process parameters for diffusion bonding of oxygen-free copper to Al2 O3 ceramics in H2 atmosphere. The strength of ceramic-to-metal diffusion bonding joints is not only related to the properties of the material itself, but also the joining process plays a decisive role in the mechanical properties of ceramic/metal diffusion bonded joints. The process parameters of the diffusion bonding directly affect the physical phase structure and strength properties of the bonding interface. Another set of process parameters and joint strengths of ceramic-to-metal diffusion joints are shown in Table 2.26.

2.4.3 Characteristics of the Al2 O3 Composite Ceramic/metal Diffusion Interface (1) Interface bonding features. At heating temperature of 1130 °C, joining time of 45 min and joining pressure of 20 MPa, the interface between the Al2 O3 –TiC composite ceramic and the W18Cr4V steel diffusion joint was tightly bonded without defects such as poor bonding and microscopic voids. The specimens of Al2 O3 –TiC composite ceramic and W18Cr4V steel diffusion joint were cut by wire cutting and prepared as metallographic specimens for analysis. Scanning electron microscopy observation of the microstructure near the Al2 O3 –TiC/W18Cr4V diffusion interface (Fig. 2.26) shows that there are white massive structure and black particles diffusely distributed on the intermediate reaction layer of the Al2 O3 –TiC/W18Cr4V diffusion interface. The energy spectrum analysis (Table 2.27) of the gray matrix structure ➀, white massive structure ➁, black particles ➂ and white dot phase ➃ in the figure shows that the gray matrix ➀ is mainly composed of Cu and a small amount of Ti. The white massive tissue ➁ is mainly composed of Cu and Ti, while the black particles ➂ are mainly Ti and the white dot ➃ contain W. It is determined that the gray matrix is Cu–Ti solid solution, the white massive structure is CuTi, the black particles are TiC, and the white dots are WC. The Cu and Ti in the reaction layer come from the dissolution diffusion during the joining of Ti-Cu–Ti intermediate layer. The W in the white dot phase is the result of the diffusion of W

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Cu foil

–

–

94%Al2 O3 ceramic + Ni、Mo、kovar alloy

95%Al2 O3 ceramic + Cu

95%Al2 O3 ceramic + 4J42

Al foil

–

420

–

Glass ceramics + Cu

94%Al2 O3 ceramic + Cu

850 ~ 900

Evaporated Cu 5 ~ 10μ m

Silica glass + Cu

–

840

–

Alumina silicate glass + Nb

Glass ceramics + Al

590

Cu foil 0.05 mm

Boron-silicon glass + kovar alloy

1150 ~ 1250

1000 ~ 1020

1050

1050

620

950

Bonding temperature/°C

Materials combination Transition layer

15 ~ 18

20 ~ 22

18

10 ~ 12

8

5

5~8

10

50 ~ 100

5

Pressure/MPa

8 ~ 10

20 ~ 25

15

50 ~ 60

60

45

15 ~ 20

30

15

20

Holding time /min

Table 2.24 Process parameters for diffusion bonding of inorganic non-metallic materials to metals

10–1

–

–

–

– (continued)

In H2 , φ135mm Porcelain insulator

In H2

In H2, bending strength 230 MPa

–

– 10–2

10–2

Tensile Strength 29 MPa, Resistance to thermal shock at 700°C

10–1 ~ 5 × 10–2

Tensile Strength 139 MPa, 600°C thermal shock 16 times

Tensile Strength 18 MPa, Resistance to Cs, 650°C, 800 h

(2 ~ 5) × 10–2

10–2 ~ 10–3

Tensile Strength 10 MPa

Notes

5 × 10–2

Vacuum degree/Pa

90 2 Welding of Advanced Ceramic Materials

1000 ~ 1100

500

Ag foil 25 μm

–

CVD deposited Ni

Deposited Ni

Ni foil

Plated Au、(Ni)

Plated Ag 6 ~ 400 8μ m,insert Ag foil 10 ~ 30μ m 1100 ~ 1150

–

–

Al foil 0.1 mm

sapphire + (Fe–Ni alloy)

BeO ceramic + Cu

ZnS Optical ceramic + Cu、 kovar

(ZnO-TiO) ceramic + Ti

(Al2 O3 -SiC-Si) ceramic + (Ni–Cr)

ZrO2 ceramic + Pt

Silicon crystal + Cu

Silicon crystal + Mo

Silicon crystal + W

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Al–Mg glass 1 ~ 10μ m 550 ~ 750

Mn(Ni) + Zn Ferrite head

1000 ~ 1050

Cu foil 0.6 mm

(yttrium-gadolinium) garnet ferrite + Cu

370

1150 ~ 1300

650

750

850

250 ~ 450

Bonding temperature/°C

Materials combination Transition layer

Table 2.24 (continued)

10 ~ 50

16 ~ 20

23

17

5 ~ 300

20

2~3

15

15

8 ~ 10

10 ~ 15

2

Pressure/MPa

15 ~ 90

15 ~ 20

60

30

50 ~ 60

60

5 ~ 20

15

15

40

10

10

Holding time /min

Tensile Strength 68 MPa

(continued)

The electromagnetic properties of ferrite are not affected after welding

– 10–1 10–1

– 10–1

Heat cycle 5 times between 300 ~ -196°C

–

10–1

–

10–1

–

(Ni–Cr) alloy contains Ni 80%, Cr 20%

10–2

10–2

–

In Ar

–

The alloy contains 46% Ni

Notes

10–2

–

–

5 × 10–2

Vacuum degree/Pa

2.4 Diffusion Bonding of Ceramics to Metals 91

–

Cr, Ni powder

Graphite + Mo、Nb 1650 ~ 1750

1250 ~ 1300

7

1100

Graphite + stainless steel

1

850

Ni foil 1 μm

1

1~2

3

850

Chemical plating Ni 10 ~ 30μ m

Graphite + Ti

Pressure/MPa

Bonding temperature/°C

Materials combination Transition layer

Table 2.24 (continued)

5

5

45

35

35

Holding time /min

–

5 × 10–4

Inert gas, Cr powder 80%, Ni powder 20%

–

– –

10–1

–

Notes

10–1

10–1

Vacuum degree/Pa

92 2 Welding of Advanced Ceramic Materials

1000

1000

1000

1000

7 + 0.4

7 + 0.5

7 + 0.5

Al2 O3 ceramics + oxygen-free copper

Al2 O3 ceramics + copper

Al2 O3 ceramics + copper

Bonding temperature/°C

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20

20

20

Holding time/min

Process parameters

7 + 0.4

Thickness/mm

Al2 O3 ceramics + oxygen-free copper

Ceramic and metal

19.6

21.56

21.56

19.6

Pressure/MPa

10

10

15

10

Heating rate/°C·min−1

10

3

10

3

Cooling rate/°C·min−1

Table 2.25 Process parameters for diffusion bonding of Al2 O3 ceramics with oxygen-free copper in H2 atmosphere

60

70

70

60 ~ 70

Total heating time/min

120

120

120

120

Total cooling time/min

2.4 Diffusion Bonding of Ceramics to Metals 93

Si3 N4 assembly

SiC assembly

1473 873

F6 8×8 —

25

600

SiC/Cr/SiC

SiC/Al-Si/Kovar

—

Si3 N4 /Ni

1273

1373

F10

—

2000 + 1000

Si3 N4 /AISI316

Si3 N4 /Fe-36Ni + Ni/MA6000 F10

1328 1323

F10 F10

25

250

Si3 N4 /V/Mo

Si3 N4 /Invar/AISI316 1473

1423

15 × 15

125

3.5 × 2.5

1473

15 × 15

125

Si3 N4 /Ni-20Cr/Si3 N4

1723 1623

F6

100

100

SiC/Co-50Ti/SiC

SiC/Fe-50Ti/SiC

1020

— F6

—

—

SiC/Ni/SiC

SiC/Cu/SiC

1200

1773

F6

20

SiC/Ti/SiC

1790

12

25

SiC/Nb/SiC

SiC/V/SiC

1373–1673

1773

F6 F6

Temperature/K

Section dimension /mm

F6

20

Thickness of interlayer /μm

SiC/Ta/SiC

Materials combination

Table 2.26 Process parameters and joint strength for ceramic-to-metal diffusion bonding

60

120

180

90

90

60

60

45

30

—

—

30

30

60

30–180

600

480

Time/min

5

100

7

7

20

100

100

20

20

20

15

4.9

7.3

7.3

30

7.3

7.3

Pressure/MPa

6.65

—

1

2

5

Ar

0.14

1.33

1.33

—

—

30

1.33

1.33

1.33

1.33

1.33

Atmosphere /mPa

(continued)

32 (tensile)

75 (bending)

37 (shear)

95 (shear)

118 (shear)

300 (bending)

100 (bending)

133 (shear)

60 (shear)

80 (—)

90 (—)

113 (bending)

89 (shear)

250 (shear)

130 (shear)

187 (shear)

72 (shear)

Joint strength /MPa

94 2 Welding of Advanced Ceramic Materials

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ZrO2 assembly

Al2 O3 assembly

2200

700

—

—

—

Al2 O3 /Al-Si/mild steel

Al2 O3 /Ti-5Ta/Al2 O3

Al2 O3 /Ag

Al2 O3 /SUS321

Al2 O3 /AA7075

126

100

Al2 O3 /Cu/AISI1015

ZrO2 /Ni–Cr-(O)/ZrO2

200

Al2 O3 /Cu/Al

125

—

ZrO2 /Ni–Cr/ZrO2

200

Al2 O3 /1Cr18Ni9Ti

50 + 100 + 1000

Al2 O3 /Ti/1Cr18Ni9Ti

Si3 N4 /Cu–Ti-B + Mo + Ni/40Cr 1173 1143 1273 773 1273 873 1423 1173 1300 633 1373 1373

F10 F10 20 × 20 F10 F32 F16 F8 F13 F10 F15 F15

1423

F14

15 × 15

10 + 60 + 60 + 60 + 10

Si3 N4 / Ni + Ni–Cr + Ni + Ni–Cr + Ni/ Si3 N4

1423

15 × 15

Temperature/K

Section dimension /mm

200

Thickness of interlayer /μm

Si3 N4 /Ni–Cr/Si3 N4

Materials combination

Table 2.26 (continued)

120

120

600

10

0

20

30

30

20

60

30

40

60

60

Time/min

10

10

6

25

3

0.2

5

3

6

7

15

30

22

22

Pressure/MPa

100

100

665

1.33

Air

0.13

30

O2

1.33

1.33

1.33

6

Ar

Ar

Atmosphere /mPa

(continued)

620 (bending)

574 (bending)

60 (shear)

60 (tensile)

70 (tensile)

56 (tensile)

23 (tensile)

100 (bending)

108 (tensile) 55 (shear)

18 (tensile)

32 (tensile)

180 (shear)

391 (bending)

160 (bending)

Joint strength /MPa

2.4 Diffusion Bonding of Ceramics to Metals 95

ZrO2 /AISI316/ZrO2

Materials combination

Table 2.26 (continued)

100

Thickness of interlayer /μm

Temperature/K

1473

Section dimension /mm F15 60

Time/min

10

Pressure/MPa

100

Atmosphere /mPa 720 (bending)

Joint strength /MPa

96 2 Welding of Advanced Ceramic Materials

2.4 Diffusion Bonding of Ceramics to Metals (a)

Al2O3-TiC

97 (b)

Interface transition zone

W18Cr4V

10µm

10µm

Fig. 2.26 Tissue characteristics of Al2 O3 –TiC/W18Cr4V diffusion joints (SEM) a Diffusion bonding interface b Interfacial transition zone

Table 2.27 Energy spectrum analysis of different morphological tissues within the reaction layer (mass fraction) % Test position

Al

O

Ti

W

Cr

Fe

Cu

➀

3.16

4.63

14.56

1.11

3.88

2.04

70.62

➁

1.24

0.66

34.66

3.27

1.87

2.74

55.56

➂

1.55

0.12

87.81

5.08

2.75

1.08

1.61

➃

0.68

0.07

2.15

92.22

2.12

1.55

1.21

elements in W18Cr4V HSS, and these diffused W and C in W18Cr4V form WC during the joining process, which is diffusely distributed in the reaction layer. (2) Division of the interface transition zone When Al2 O3 –TiC is diffusion bonded with W18Cr4V, diffusion occurs between Ti and Cu due to the concentration gradient at the Ti-Cu–Ti interlayer interface, and when the heating temperature is higher than the Cu–Ti eutectic temperature, the Cu– Ti liquid phase diffuses and reacts into the Al2 O3 –TiC ceramic and W18Cr4V steel on both sides. The elements in the parent material also diffuse into the interlayer, forming a diffusion reaction layer (or called interface transition zone) with different microstructure near the interface between Al2 O3 –TiC ceramic and W18Cr4V. The backscattered electron image and elemental line scan results near the Al2 O3 – TiC/W18Cr4V diffusion interface are shown in Fig. 2.27. As seen in Fig. 2.27a, there are obvious interfacial transition zones between Al2 O3 –TiC ceramics and W18Cr4V steel, which can be divided into four reaction layers according to their locations, namely Al2 O3 –TiC/Ti interface reaction layer A, Cu– Ti solid solution layer B, reaction layer C at Ti side of Ti/W18Cr4V interface and reaction layer D at W18Cr4V steel side. As seen from the line scan of elements shown in Fig. 2.27b, layer A contains Ti, Al, and O, mainly from Al2 O3 –TiC ceramics and Ti in the interlayer; layer B

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Fig. 2.27 Backscattered electron images and elemental line scans near the Al2 O3 –TiC/W18Cr4V diffusion interface a Electron backscattered phase b Element line scanning

contains Cu and a small amount of Ti from the Ti-Cu–Ti interlayer; layer C contains Ti, mainly from the Ti-Cu–Ti interlayer; layer D is Fe and Cr from W18Cr4V steel. The elemental distribution in each layer is consistent with that of the initial state of joining. When the heating temperature is 1100 °C and the bonding time is 30 min, the element diffusion is not sufficient and the diffusion distance is short. As the heating temperature was increased and the holding time extended, the element diffusion was further intensified and the interfacial reaction was more adequate. As the diffusion

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joining process parameters varied, the microstructure of each reaction layer in the interfacial transition zone will also be changed. There are Al, Ti, Cu, Fe, W, Cr, V and other elements in Al2 O3 –TiC/W18Cr4V interface transition zone. In the diffusion bonding process, the element diffusion and mutual reaction make the microstructure of the interface transition zone is complex, forming several characteristic zones A, B, C, D. The microstructure of the reaction layer A near the Al2 O3 –TiC ceramic side is the dark gray matrix mixed with a large number of TiC black particles, and the TiC particles are aggregated at the Al2 O3 –TiC/reaction layer A and B interfaces, as shown in Fig. 2.28a and b. Ti in the interlayer reacts with Al2 O3 , and the TiC particles that do not participate in the reaction aggregated near the interface. Reaction layer B has matrix with light gray color, and there are much smaller black and white particles within the light gray matrix than reaction layer A. The boundaries of reaction layer A and reaction layer B are not obvious and they cross each other. Reaction layer C has shows a black band, as shown in Fig. 2.28c; some white dot particles are present in reaction layer D (Fig. 2.28d), which may be the result of segregation in the composition of the microregion.

(a)

(b)

Al2O3-TiC

Reaction layer B Reaction layerA

5µm

5µm (d)

(c)

Reaction layer D

Reaction layer C

W18Cr4V

1µm

1µm

Fig. 2.28 Microstructure of the transition zone at the Al2 O3 –TiC/W18Cr4V diffusion interface a Reaction layer A b Reaction layer B c Reaction layer C d Reaction layer D and W18Cr4V

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(b)

Interface transition zone

W18Cr4V

Al2O3-TiC

Interface transition zone

10µm

W18Cr4V

20µm

Fig. 2.29 Microstructure of the transition zone at the Al2 O3 –TiC/W18Cr4V interface at different holding times and pressures a 1130 °C × 30 min, P = 10 MPa b 1130 °C × 60 min, P = 15 MPa

The diffusion bonding temperature determines the extent of diffusion and interfacial reactions of elements near the interface. The holding time t is the main factor determining the diffusion uniformity of elements near the diffusion joint interface. The role of the joining pressure p is to cause microscopic plastic deformation at the contact interface and promote close contact at the joining surface. The microstructure of the Al2 O3 –TiC/W18Cr4V interface transition zone for different holding times and joining pressures at a heating temperature of 1130 °C is shown in Fig. 2.29. As seen in Fig. 2.29, the width of the Al2 O3 –TiC/W18Cr4V interfacial transition zone is only about 25 μm at the holding time of 30 min and joining pressure of 10 MPa. The microstructure is not uniform, with a few microscopic voids at the interface between the interfacial transition zone and W18Cr4V, and the interface is not tightly bound. When the holding time was 60 min and the joining pressure was 15 MPa, the microstructure morphology of the interfacial transition zone was basically the same, with some white massive structure and black particles distributed on the gray matrix. The effect of pressure on the microstructure of the ceramic/metal diffusion interface is manifested by promoting the close contact between the interfaces and providing the necessary conditions for the diffusion reaction between the interlayer and the base material on both sides. During the ceramic/metal diffusion bonding process, the heating temperature, holding time and joining pressure interact with each other to influence the microstructure and properties of the transition zone at the ceramic/metal interface. (3) Microhardness at the interface transition zone The microhardness of the transition zone at the ceramic/metal interface reveals the changes in the microstructure of the region. The microhardness of the Al2 O3 – TiC/W18Cr4V interface transition zone and the parent material on both sides nearby were measured by microhardness tester with test load of 100 g and loading time of 10 s. The microhardness distributions near the Al2 O3 –TiC/W18Cr4V interface zone at different heating temperatures and holding times are shown in Figs. 2.30 and 2.31.

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Fig. 2.30 Microhardness near the Al2 O3 –TiC/W18Cr4V interface (1110 °C × 45 min) a Test position b Microhardness

Fig. 2.31 Microhardness distribution in the transition zone of the Al2 O3 –TiC/W18Cr4V interface (1130 °C × 60 min) a Test position b Microhardness

As seen in Fig. 2.30, the microhardness of the interfacial transition zone from the Al2 O3 –TiC side to the W18Cr4V side at a heating temperature of 1110 °C and a holding time of 45 min is about 350 HM, and that of the W18Cr4V HSS is about 470 HM. The microhardness of the Al2 O3 –TiC ceramics is much higher than that of the W18Cr4V steel, which also further indicates that the microstructure and properties of Al2 O3 –TiC and W18Cr4V are very different. The microhardness of the interfacial transition zone is lower than that of the parent material on both sides, which is mainly due to the low heating temperature and short holding time, the diffusion of Cu and Ti in the Ti-Cu–Ti interlayer is not sufficient, and only a small amount of Ti diffuses into Cu. It can be seen from the figure that the interfacial transition zone is much narrow and the microhardness test point is located in the middle of the interfacial transition zone, where the Cu layer is located, so the microhardness is lower.

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The microhardness distribution of the Al2 O3 –TiC/W18Cr4V interfacial transition zone under the conditions of heating temperature 1130 °C × 60 min is shown in Fig. 2.31. The microhardness gradually decreases from the Al2 O3 –TiC side to the W18Cr4V side. The microhardness of the interface transition zone near the Al2 O3 – TiC side is about 1200 HM, which is higher than that of the interface transition zone near the W18Cr4V side (800 HM). The microhardness of the interfacial transition zone at the process parameter of 1130 °C × 60 min is higher than that of the interfacial transition zone at the process parameter of 1110 °C × 45 min. This is due to the fact that increasing the heating temperature and prolonging the holding time allowed the Ti in the Ti-Cu–Ti interlayer to diffuse into the Cu, which improved the hardness of the interfacial transition zone. Ti is an active element and reacts with the elements from Al2 O3 –TiC and W18Cr4V to form compounds that also improve the microhardness of the interfacial transition zone. As seen in Figs. 2.30 and 2.31, the microhardness of the Al2 O3 –TiC/W18Cr4V interfacial transition zone is lower than that of the Al2 O3 –TiC ceramics, indicating that no high hardness brittle phase with higher hardness than that of the Al2 O3 –TiC ceramics is generated during the Al2 O3 –TiC/W18Cr4V diffusion bonding. (4) Phase structure of the interface transition zone. When diffusion joining of Al2 O3 –TiC ceramics and W18Cr4V steel with a Ti-Cu– Ti interlayer, a large elemental concentration gradient exists between the interlayer and the parent material on both sides. At high temperature of diffusion joining, Ti and Cu in the interlayer undergo mutual diffusion and chemical reaction, and the activity of Ti makes Ti react with Al, O, C in Al2 O3 –TiC ceramic and with Fe, W, Cr and C in W18Cr4V steel to form new compounds. The reactions between various elements of Al2 O3 –TiC ceramics and W18Cr4V steel may also occur in Al2 O3 –TiC and W18Cr4V interface transition zone, which will produce a variety of generated phases. Specimens were cut from the Al2 O3 –TiC/W18Cr4V diffusion joint using a wire cutting machine and analyzed by X-ray diffractometer (XRD, type D/MAX-RC) for the phase composition of the interface transition zone. Prior to the test, the joint specimens were divided into two parts from the Al2 O3 –TiC/W18Cr4V diffusion interface by applying shear force from the Al2 O3 –TiC side and the W18Cr4V side, see in Fig. 2.32a. The specimen size was 10 mm × 10 mm × 7 mm, and the analytical surface for the X-ray diffraction test is shown in Fig. 2.32b. The X-ray diffraction test was performed using a Cu-Kα target with an operating voltage of 60 kV, an operating current of 40 mA, and a scanning speed of 8°/min. The X-ray diffraction patterns of both sides of the Al2 O3 –TiC/W18Cr4V diffusion interface are shown in Fig. 2.33. Comparison X-ray diffraction analysis (XRD) data of the Al2 O3 –TiC and W18Cr4V diffusion interface with the standard powder diffraction card published by the Joint Committee on Powder Diffraction Standards (JCPDS). Result shows that on the Al2 O3 –TiC ceramic side of the diffusion joint, four main phases such as Al2 O3 , TiC, TiO, and Ti3 Al are present. On the W18Cr4V side, the phases are more complex and include Al2 O3 , TiC, Cu, CuTi, CuTi2 , Fe3 W3 C, FeTi, etc.

2.4 Diffusion Bonding of Ceramics to Metals

a

103

b

Fig. 2.32 Al2 O3 –TiC/W18Cr4V specimen for X-ray analysis and analysis position a Al2 O3 TiC/W18Cr4V sample b XRD analysis surface

During the diffusion bonding of Al2 O3 –TiC composite ceramics with W18Cr4V steel, no mutual reaction occurs between the Al2 O3 matrix and the TiC reinforcing phase in Al2 O3 –TiC composite ceramics at the bonding temperature of 1130 °C. At the Al2 O3 –TiC/Ti interface, Ti reacts with Al2 O3 to form Ti3 Al and TiO because Ti is an active element and the thickness of the Ti foil is small. The brittleness of the Ti3 Al phase is great, and the interface on the Al2 O3 –TiC ceramic side containing more Ti3 Al phase, which is the weaker part of the diffusion joint in performance. The Al2 O3 phase and TiC phase measured by X-ray diffraction test on the W18Cr4V side were derived from Al2 O3 –TiC ceramic, which remained on the W18Cr4V surface after Al2 O3 –TiC/W18Cr4V diffusion joint shear fracture. It indicates that the shear specimen fractured at the diffusion interface near the Al2 O3 –TiC ceramic side. The Ti-Cu–Ti interlayer generated during the diffusion joining process Cu–Ti solid solution or Cu–Ti compounds such as CuTi, CuTi2 etc. The unreacted part of Cu remains in the form of monomers. At the W18Cr4V/Ti interface, Ti is a carbide-forming element and is highly susceptible to forming TiC with C in the steel, which will prevent the diffusion of Ti into Fe. Since the solubility of Ti in Fe is very small, the diffusion of Ti into Fe will form FeTi or Fe2 Ti intermetallic compounds in addition to solid solution. W18Cr4V HSS contains Fe, W, Cr, V, and C. At the diffusion bonding temperature of 1130 °C, reactions between these elements may also occur to form new compounds. XRD analysis revealed the Fe3 W3 C phase.

2.4.4 Diffusion Bonding of SiC/Ti/SiC Ceramics Reliable joining of SiC ceramics can be achieved under diffusion bonding conditions using Ti as an interlayer. The SiC/Ti/SiC interfacial reaction was studied in the range of joining temperatures from 1373 to 1773 K and holding times from 5 to 600 min.

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Fig. 2.33 X-ray diffractogram of the Al2 O3 –TiC/W18Cr4V diffusion interface a Al2 O3 -TiC ceramic side b W18Cr4V steel side

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The highest shear strength was obtained at the optimum joining parameters of 1773 K × 60 min. (1) SiC/Ti/SiC interfacial reaction Figure 2.34 shows a schematic of the reaction process at the SiC/Ti/SiC diffusion interface. In the early stage of the reaction, SiC reacts with Ti to form TiC and Ti5 Si3 Cx , and TiC has limited growth on the Ti side due to the fast diffusion of C, while Ti5 Si3 Cx is formed on the SiC side. As the Si and C on the SiC side diffuse through the Ti5 Si3 Cx layer to the middle, the Ti in the middle region also diffuses into the Ti5 Si3 Cx . Due to the slow diffusion of Si, it is difficult to reach elemental equilibrium in Ti5 Si3 Cx , so the TiC phase precipitates as block in Ti5 Si3 Cx . The interfacial structure at this point is shown in Fig. 2.34b, showing a laminar arrangement of SiC/Ti5 Si3 Cx + TiC/TiC + Ti/Ti. The laminated Ti5 Si3 Cx phase is generated at the interface of SiC/Ti5 Si3 Cx + TiC when the joining time is extended to 0.9 ks, and the interface structure becomes SiC/Ti5 Si3 Cx /Ti5 Si3 Cx + TiC/TiC + Ti/Ti, as shown in Fig. 2.34c. The sequence of phase formation in this reaction system is independent of the thickness of the Ti interlayer, but the appearance time of each reacting phase decreases with the increase of the bonding temperature. Regarding the appearance of Ti5 Si3 Cx singlephase layer, the analysis revealed that it is mainly caused by the different diffusion rate of each element, and the diffusion rate of Ti element into the ceramic is slower than the diffusion rate of Si and C elements into the Ti metal. The Ti content at the interface near the ceramic side so low that TiC cannot be formed. In the middle stage of the reaction, the reaction cannot be balanced due to the aggregation of Si and C elements at the SiC/Ti5 Si3 Cx interface, and the hexagonal crystalline Ti3 SiC2 phase is formed at the interface again. As shown in Fig. 2.34d

Fig. 2.34 Variation of SiC/Ti/SiC interface structure with diffusion connection time

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and e, and the arrangement of interfacial layer changes to SiC/Ti3 SiC2 /Ti5 Si3 Cx /Ti5 Si3 Cx + TiC/TiC/Ti. To investigate the equilibrium process of SiC and Ti, the joining time is further extended and the reaction enters the later stage where all the Ti participates in the reaction and disappears away at the interface. Due to the diffusion of Si and C on both sides, the mixed phase of Ti5 Si3 Cx + TiC were also all disappears, at this time the order of the interfacial layer changes to SiC/Ti3 SiC2 /Ti5 Si3 Cx + Ti3 SiC2 /Ti3 SiC2 /SiC. After the bonding time reached more than 36 ks, the TiC single phase all participated in the reaction and the fine Ti3 SiC2 phase was observed in the Ti5 Si3 Cx phase, while the Ti3 SiC2 layer and the Ti3 SiC2 /Ti5 SiC3x interfaces formed rhombohedral crystals of TiSi2 compounds as shown in Fig. 2.34g. Further extending the joining time to 108ks, the interfacial microstructure is shown in Fig. 2.34h, the Ti5 Si3 Cx phase also disappears, and the joint interface becomes mixed microstructure consisting of Ti3 SiC2 and TiSi2 , which basically achieves the phase equilibrium in the Ti-Si–C ternary phase diagram. (2) Conditions for the formation of interfacial reaction phases. The relationship between the reaction products at the SiC/Ti interface as a function of joining temperature and time is shown in Fig. 2.35, and the symbols in the figure are experimental data points. The graph gives the conditions for the formation of each reactant (bonding temperature and time). This figure serves to predict the type of compound produced at the interface based on the joining conditions and also to determine the joining conditions based on the type of compound desired. The thickness of the Ti interlayer used for the test was 50 μm. The first line from the low temperature side is the generation curve of single phase Ti5 Si C3x . In the region below this curve, the interfacial reaction products are TiC

Joining temperature

Fig. 2.35 Variation of reaction products with temperature and time

Joining time

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and Ti5 Si3 Cx , forming a mixture of bulk TiC and TiC + Ti5 Si3 Cx microstructure. If the temperature and time required for this line were reached, lamellar Ti5 Si3 Cx was formed. With the increase of temperature or the extension of the joining time, the Ti3 SiC2 phase appears at the interface. At this time the diffusion path of SiC and Ti interfacial reaction is fully formed and the interfacial structure presentes as SiC/Ti3 SiC2 /Ti5 Si3 Cx /Ti5 Si3 Cx + TiC/TiC + Ti/Ti. Further increase of the joining temperature or extension of bonding time, the more stable silicide TiSi2 appears at the interface. (3) Mechanical properties of diffusion joints. The results of shear tests on SiC/Ti/SiC diffusion bonded joints showed that the shear strength of the diffusion bonded joints was about 44 MPa at joining temperature of 1100 °C, and the joint shear strength rose to 153 MPa at joining temperature of 1200 °C. When the joining temperature was further increased to 1500 °C, the joint shear strength reached a maximum of 250 MPa. From the site of fracture occurrence, it is known that fracture occurs at the interface of SiC/Ti5 Si3 Cx + TiC at the temperature interval below 1200 °C. For the joints above 1400 °C, fracture occurs on the SiC ceramic parent material near the bonding layer and develops within the SiC in the direction of the bonding surface. The analysis of the fracture microstructure shows that the fracture surface is flat at 1100 °C. The fracture surface is more concave and convex at 1200 °C, and there are more massive reactive phase of Ti5 Si3 Cx + TiC on the SiC fracture surface. The hardness of TiC is the highest among all Ti compounds, and the difference between the linear expansion coefficients of TiC and SiC is the smallest. Besides, there is also a good correspondence in crystallography between them, so it can be inferred that the strength of SiC/TiC interface is higher. The joint has the maximum shear strength at the joining temperature of 1500 °C. The interface is directly connected between SiC and Ti3 SiC2 , and there is also a good crystalline correspondence between them, although the brittle phase TiSi2 is also present, it is diffusely distributed in Ti3 SiC2 , so the joint exhibits high bond strength. The diffusion welded joints with the best bonding parameters (1500 °C × 3.6 ks) were selected to determine the high-temperature shear strength of SiC/Ti/SiC diffusion joints. The test results showed that the high-temperature shear strength of the joint could be maintained up to about 800 °C. The shear strength was slightly higher than that at room temperature, showing good high-temperature resistance. The broken position at high-temperature was the same as at room temperature, which also occurred on the SiC ceramic base material near the diffusion interface.

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2.5 Electron Beam Welding of Ceramics to Metals 2.5.1 Characteristics of Electron Beam Welding of Ceramics and Metals Since the 1960s, foreign countries have begun to apply electron beam welding to the welding process of metals and ceramics. This method has expanded the scope of application of engineering materials, but also to improve the gas tightness and mechanical properties of ceramic welded parts, meeting the needs of many aspects. Electron beam welding is a welding method in which a high energy density electron beam is used to bombard the welded part to make it locally heated and melted. Electron beam welding of ceramics and metals is a very effective method. Since the welding is done under vacuum conditions, it can prevent the pollution of oxygen and nitrogen in the air, which is beneficial to the welding of ceramics and metals, and the gas tightness after welding is good. Electron beam can form a very small diameter by focusing, which can be as small as 0.1 ~ 1.0 mm. Its power density can be increased to 106 ~ 108 W/cm2 degree. Thus the electron beam penetration effect is very strong, the heating area is very small and also the width of weld bead is small.The melt depth is very large, and the ratio of melt width to melt depth can reach (1:10) ~ (1:50). In this way, not only the welding heat affected zone is small, but also the stress is very small. This can ensure the accuracy of the structure after welding for ceramic finishing parts as the last process. The disadvantages of this method are the complexity of the equipment, the strict requirements for the welding process and the high production costs. In vacuum electron beam welding of ceramics and metals, there are various forms of joints for the welded parts, and the more suitable joint form is flat welding as the best. It can also be used in lap welding or sleeve welding, the assembly gap between the workpiece should be controlled at 0.02 ~ 0.05 mm and should not be too large, otherwise it may produce defects such as lack of penetration. The vacuum electron beam welding machine for ceramics and metals is composed of four parts: electron optical system (including electron gun, magnetic focusing and deflecting system), vacuum system (including vacuum chamber, diffusion pump, mechanical pump), and transmission mechanism, power supply and control system. The main component of the electron beam welding machine is the electron optical system, which is the key to obtain a high energy density electron beam, and can ensure the process stability of electron beam welding after being equipped with a stable and easily adjustable power supply system. The acceleration voltage of electron beam welding gun is divided into high voltage type (110 kV or more), medium voltage type (40–60 kV) and low voltage type (15–30 kV). For the welding of ceramics and metals, the most suitable type is high vacuum low voltage type electron beam welding gun.

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2.5.2 Processes for Electron Beam Welding of Ceramics to Metals ➀ Clean the surface of the weldment and place the assembled workpiece in the preheating furnace. ➁ After the vacuum degree in the vacuum chamber has reached 1.33 × 10–2 Pa, preheating of the workpiece with a tungsten thermal resistance furnace is started. ➂ The welding is started by allowing the electron beam to sweep the metal side of the workpiece to be welded at a preheated constant temperature. ➃ Cooling and annealing after welding, the voltage value drop to zero within 10 min in preheating furnace, then make the welded parts in the vacuum furnace natural cooling for 1 h. It can be taken out of the furnace after slow cooling. The main welding parameters of electron beam welding are: accelerating voltage, electron beam current, working distance (distance from the welded workpiece to the bottom of the focusing cylinder), focusing current and welding speed. The process parameters in vacuum electron beam welding of ceramics and metals have a great influence on the quality of the joint, especially on the weld depth and width of the weld, which are more sensitive. And this is an important indicator of the quality of electron beam welding. The selection of suitable welding parameters can make the weld shape, strength, gas tightness, etc. meet the design requirements. For electron beam welding between alumina ceramics (85%, 95% Al2 O3 ), high purity Al2 O3 and translucent Al2 O3 ceramics, the following process parameters can be selected: power of 3 kW, acceleration voltage of 150 kV, maximum electron beam current of 20 mA. Direct welding by high voltage electron beam weld machine with an electron beam focusing diameter of 0.25 to 0.27 mm, the high welding quality can be obtained. Electron beam welding of high purity Al2 O3 ceramics with refractory metals (W, Mo, Nb, Fe-Co–Ni alloys) can also be performed with the above process parameters using a high voltage electron beam welder. Electron beam welding of two translucent Al2 O3 ceramic butt joints can also be performed with Nb sheets of 0.5 mm thickness as an transition. Electron beam welding can also be performed with a 1.0 mm diameter metal molybdenum needle to alumina ceramics. Vacuum electron beam welding is currently mostly used for welding refractory metals (W, Mo, Ta, Nb, etc.) to ceramics, and it is important to make the linear expansion coefficient of the ceramics similar to that of the metal, so as to achieve a matched joint. As the heating spot of electron beam is very small, it can be concentrated on a very small area of heating. Qualified welded joints can be obtained by taking measures such as preheating before welding, slow cooling after welding and reasonable design of joint form.

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2.5.3 Example of Electron Beam Welding of Ceramics to Metals Some sensors used in petrochemical and other departments sectors require work in strongly leaching media. These sensors often use alumina series ceramics as the insulating material, while the conductor is 18–8 stainless steel. There should be a reliable connection between the stainless steel and the ceramic, and the weld must be heat resistant, corrosion resistant, strong and reliable and dense without leakage. For example, the ceramic part is a tube with a length of 15 mm, an outer diameter of 10 mm and a wall thickness of 3 mm. A dynamic fit is used between the ceramic tube and the metal tube. The ceramic tube is left with a heating expansion gap of 0.3 to 1.0 mm at each end to prevent the ceramic tube from bursting due to stresses generated during welding heating. Vacuum electron beam welding method was used to weld the 18–8 stainless steel tube with the ceramic tube, the joint was a lap weld, and the process parameters for electron beam welding are shown in Table 2.28. First of all, the surface of the ceramic and metal welded parts are cleaned, and the pickling method is taken to remove grease and dirt. Before electron beam welding, the workpiece were graded heated to 1200 °C at the heating rate of 40 to 50 °C/min, and holding 4 ~ 5 min. Then turn off the preheating power to make the ceramic parts preheat uniformly. When the joint temperature is lowered, one of the ends of the workpiece were weld and heated evenly while welding. After the first weld is made, the workpiece should be reheated to 1200 °C before the second weld is made. After the joint is welded, it should be cooled at the rate of 20 to 25 °C/min within the furnace, and the cooling rate should not be too fast. In the cooling process after welding, due to the role of shrinkage forces, the axial extrusion force is first produced in ceramics. So the workpiece should be slowly cooled to below 300 °C before it can be removed from the heating furnace, to prevent excessive squeezing pressure which may cause squeeze cracking in ceramic. Compared to metals and plastics, ceramic materials are hard, non-flammable and non-reactive. Therefore, ceramics can be used in high temperature, corrosive and high friction environments, including: lasting stability of various physical properties at high temperatures, low coefficient of friction (especially under heavy loads and low lubrication conditions), low linear expansion coefficient, corrosion resistance, thermal insulation, electrical insulation, and low density. Engineering ceramics are used in many industrial sectors to manufacture components, including ceramic substrates for electronic devices, rotors for turbochargers and tappet heads in automotive engines. Other examples of modern ceramic applications are oil-free lubricated bearings used in food processing equipment, aero turbine blades, atomic nuclear fuel rods, lightweight armour plates, cutting tools, abrasives, thermal spacers and heat-resistant components for furnaces and kilns. Future developments are likely to come from improved processing and manufacturing techniques for ceramic-to-metal materials to reduce component costs and improve performance, demanding higher standards for high-performance materials

8 8 8 10 12

4+4

5+5

6+6

8+8

10 + 10

18–8 Steel/ceramic

18–8 Steel/ceramic

18–8 Steel/ceramic

18–8 Steel/ceramic

18–8 Steel/ceramic

Electron-beam current/mA

Process parameters

Parent metal thickness/mm

Materials

14

13

12

11

10

Acceleration voltage/kV

55

58

60

62

62

Welding speed/m·min−1

Table 2.28 Process parameters for vacuum electron beam welding of 18–8 stainless steel to ceramics

1200

1200

1200

1200

1250

Preheating temperature/°C

25

23

22

22

20

Cooling rate/°C·min−1

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and requiring the use of more ceramic-to-metal composites. The joining of ceramics to metal materials will bea development field of great concern.

Bibliography 1. Ren J-L, Wu A-P (2000) Advanced materials joining. Machinery Industry Press, Beijing 2. Kou S (2002) Welding metallurgy. A Wiley-Interscience Publication, America 3. Wu A-P, Zou G-S, Ren J-L (2002) Development of advanced structural ceramics and its progress in brazing joining technology. Mater Sci Eng 20(1):104–106 4. Li C-Y, Qian E-Y, Zhang J-H et al (2000) Advanced joining methods. Machinery Industry Press, Beijing 5. Hongyuan F, Jicai F (2005) Interfacial behavior during material joining. Harbin Institute of Technology Press, Harbin 6. Welding Society of Chinese Mechanical Engineering Society (2008) Welding of materials, 3rd edn. Welding handbook, vol 2. Machinery Industry Press, Beijing 7. Peisheng C (1999) Handbook of brazing. Machinery Industry Press, Beijing 8. Cawley JD (1991) Introduction to ceramic-metal joining. The Minerals, Metals and Materials Society, pp 3–11 9. Tanaka T, Morimoto H, Homma H (1988) Joining of ceramics to metals. Nippon Steel Techn Rep 37:31–38 10. Jicai F, Xiangmeng J, Lixia Z et al (2006) Interface structure and bonding strength of brazed joint of TiC cermet/steel. Trans China Weld Institut 27(1):5–8 11. Wang SM, Lu SKN, Hua Z et al (2003) Study on the Sintering Process of Ti-Al/TiC Composites. J Mater Sci Eng 21(4):565–568 12. Ying W, Jian C, Jicai F et al (2009) TLP bonding of alumina ceramic and 5A05 aluminum alloy using Ag–Cu–Ti interlayer. China Weld 18(4):39–42 13. Zhang Y, He CY, Feng J (2007) Recent progress of interlayer used to join metals and ceramics. J Iron Steel Res 19(2):1–4, 34 14. Lemus J, Drew RAL (2003) Joining of silicon nitride with a titanium foil interlayer. Mater Sci Eng A 352(1–2):169–178

Chapter 3

Diffusion Welding of Composite Ceramics to Steel

Composite ceramics can be applied in the preparation of cutting tools due to the addition of reinforcing particles (e.g. TiC) to the matrix (e.g. Al2 O3 ), which gives them higher hardness, strength, and fracture toughness. The joining of Al2 O3 –TiC composite ceramics with carbon steel, stainless steel or tool steel (e.g. W18Cr4V steel) by diffusion welding to make composite members is important to improve the internal stress distribution state of structural members under stress and to broaden the use of Al2 O3 –TiC composite ceramics. The diffusion welding of composite ceramics to steel has also received much attention.

3.1 Diffusion Welding Process of Composite Ceramics to Steel 3.1.1 Basic Properties of Al2 O3 –TiC Composite Ceramics The Al2 O3 –TiC composite ceramics for the test were made by hot press sintering (HPS) in the form of circular specimens with a size of φ 52 mm × 3.5 mm. The chemical composition, thermophysical properties and mechanical properties of Al2 O3 – TiC composite ceramics are shown in Table 3.1. Al2 O3 –TiC composite ceramics are composed of Al2 O3 matrix and TiC reinforced particles which size is about 2.0 μm, and there is also a trace of adhesive phase which can increase the bonding strength and the toughness of the ceramic matrix. The microstructure morphology of Al2 O3 –TiC composite ceramics is shown in Fig. 3.1a. When observed under scanning electron microscope (SEM), the dark gray matrix is Al2 O3 , and the black particles are TiC reinforced phase, see Fig. 3.1b.

© Chemical Industry Press 2023 Y. Li, Joining Technology and Application of Advanced Materials, Advanced and Intelligent Manufacturing in China, https://doi.org/10.1007/978-981-19-9689-4_3

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Table 3.1 Chemical composition of Al2 O3 –TiC composite ceramics (mass fraction) % Element

C

O

Al

Ti

Cr

Fe

Ni

Mo

W

Total

Mass fraction (%)

18.74

29.12

13.51

28.86

0.22

1.34

2.41

3.86

1.95

100

(a)

(b)

10μm

25μm Fig. 3.1 Microstructure of Al2 O3 –TiC composite ceramics

3.1.2 Process Characteristics of Composite Ceramic to Steel Diffusion Welding (1) Diffusion welding equipment For vacuum diffusion joining of Al2 O3 –TiC ceramics and W18Cr4V steel, Workhorse II vacuum diffusion welding equipment manufactured by American Vacuum Industries was used, and its main performance indicators are shown in Table 3.2. Workhorse II vacuum diffusion welding equipment for testing mainly consists of fully automatic vacuum pumping system, vacuum furnace body, pressurization system, heating system, water circulation system and control system. Since the whole set of equipment is computer-controlled, the vacuum diffusion welding process can be operated automatically and the process parameters can be controlled with high precision. The heating temperature, connection pressure, holding time, vacuum degree and other parameters of the vacuum diffusion welding can be controlled by the Table 3.2 Main performance indicators of Workhorse II vacuum diffusion welding equipment Model

3033-1305-30 T

Main performance indicators Vacuum-chamber dimensions/mm

Maximum heating temperature/K

Maximum pressure/T

Limiting vacuum degree/Pa

Heating power/kVA

304 × 304 × 457

1623

30

1.33 × 10–5

45

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pre-programmed welding procedure, which improves the reliability of the welding process. (2) Interlayer materials The interlayer plays a key role both in the diffusion welding process of composite ceramics and metals and in ensuring the performance of the joint. This is because the final microstructure of the ceramic–metal diffusion joint depends mainly on the interlayer. And the diffusion properties and diffusion mode of the interlayer alloying elements are the key conditions for the transient liquid phase solidification and composition homogenization. The main roles of the interlayer material are as follows. ➀ Relieving residual stresses between ceramics and metals due to their different coefficients of thermal expansion and improving the joint strength. ➁ Promoting interfacial wetting and diffusion through melting or reaction with ceramics to form a strong metallurgical bond. ➂ Controlling interfacial reactions, altering or inhibiting interfacial products and leaving the interface in a more stable thermodynamic state. Intermediate layer materials also help to eliminate holes at the joining interface and form a better sealing ceramic/metal bonding joint. Ti is an active element with good wettability to composite ceramics and can form stable interface connections after reaction with ceramics. Ti can be applied to ceramic/ceramic or ceramic/metal connections. Cu is a softer material and is a good buffer layer material, which can reduce the residual stress at the diffusion interface. A eutectic reaction will occur between Ti and Cu above the eutectic temperature to form Cu–Ti liquid alloy, which plays the role of promoting wettability. During the joining process, the Cu–Ti liquid phase alloy and the base material on both sides diffuse each other, and the interfacial reaction occur to form the interface connection. Therefore, it is appropriate to design a Ti–Cu–Ti multi-interlayer to perform diffusion connection of Al2 O3 –TiC composite ceramics with steel. Using Ti–Cu–Ti multi-interlayer to realize the diffusion welding between Al2 O3 – TiC composite ceramics and steel can achieve a strong metallurgical bond. The chemical composition and thermo-physical properties of the Ti and Cu interlayer are shown in Table 3.3. The cleanliness and flatness of the specimen surface are important factors affecting the quality of the diffusion welding joint. Before diffusion welding, the surfaces of the specimen to be joined (Al2 O3 –TiC composite ceramic, Q235 steel, 18-8 steel or W18Cr4V steel) are polished smoothly with metallographic sandpaper, soaked in acetone, then scrubbed cleanly with alcohol, blown dry and left to use. The Cu foil and Ti powder were made into a composite interlayer material in the order of Ti–Cu–Ti. The assembled specimens were stacked into the vacuum chamber in the order of Al2 O3 –TiC/Ti–Cu–Ti/steel (Q235 steel, 18-8 steel or W18Cr4V steel)/Ti–Cu– Ti/Al2 O3 –TiC. The assembly of the diffusion welding specimen is shown in Fig. 3.2.

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-

-

Cu

Density/g·cm-3

4.5

8.92

Melting point/K

1913–1943

1338–1355

Ti

Cu

-

-

N 0.05

O

0.25

Material

Thermo-physical properties

Ti

C

0.10

H

0.015

Material

Chemical compositions (wt%) S 0.005

-

f.c.c

h.c.p

Crystal structure

0.001

-

Bi

Sb 0.002

-

16.92

8.2

Coefficient of thermal expansion/10-6 ·K-1 (20–100 °C)

0.005

0.30

Fe

Table 3.3 Chemical compositions and thermo-physical properties of Ti and Cu interlayer As 0.002

-

Pb

Ti or Cu Bal

Bal

125

115

Young’s modulus/GPa

0.005

-

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Fig. 3.2 Illustration of the assembly of the diffusion welding specimen

(3) Control of heating and cooling temperatures After the diffusion welding specimen is placed in the vacuum chamber, the vacuum chamber is evacuated to 1.33 × 10–4 –1.33 × 10–5 Pa. Then the running program is started and the temperature rise begins according to the setup program. After reaching the predetermined diffusion joining temperature, hold it until the set joining time so that the interlayer materials react fully with the joined materials. During the whole joining process, the vacuum chamber of the furnace chamber is kept in a high vacuum degree. To ensure good contact between the materials, pressure needs to be applied during the diffusion joining process. It is very important to control the heating and cooling speed during the joining process. Since Al2 O3 –TiC ceramics are brittle materials with poor resistance to cold and thermal shock, excessive heating or cooling speed may induce cracks within the ceramic and affect the final performance of the joint. If the vacuum chamber size of the diffusion welding equipment is large, it will take some time to make the temperature uniform, so the average temperature platform should be set up during heating and cooling to ensure uniform temperature in the furnace. After the vacuum chamber of the diffusion welding equipment is evacuated to 1.33 × 10–4 Pa, the temperature begins to rise with graded heating method. Several temperature holding platforms are set up. ➀ Heat to 350 °C at a rate of 15 °C/min from 20 °C room temperature and hold for 10 min. ➁ Heat to 600–650 °C at 15 °C/min and hold for 10 min. ➂ Heat to 900–950 °C at 15 °C/min and hold for 10 min. ➃ Heat to 1100–1150 °C at 10 °C/min and hold for 45–60 min. After holding at 1100–1150 °C for 45–60 min, the specimens were cooled to 950 °C at 10 °C/min and held for 2 min. Then they were cooled to 100 °C using circulating water and cooled to room temperature with the furnace at last. Before heating to diffusion welding temperature (1100–1150 °C), the hydraulic system is pressurized to about 10–15 MPa and held for 45–60 min. The pressure is removed and the cooling stage starts. The typical curve of the diffusion joining process is shown in Fig. 3.3. The protective gases are nitrogen (N2 ) and argon (Ar).

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T

Holding temperature

Heating temperature /ºC

1200 1000 800

Cooling rate

600 400

Heating rate

200 t

0 0

100

200

300

400

500

600

700

800

900 1000

Time /min Fig. 3.3 Typical curve of Al2 O3 –TiC/steel diffusion welding process

When Al2 O3 –TiC composite ceramic is diffusion connected with steel by adding an Ti–Cu–Ti interlayer, there is a diffusion reaction between the elements of the interlayer and the matrix (Al2 O3 –TiC composite ceramic, Q235 steel, 18-8 steel and W18Cr4V steel) on both sides. A diffusion welding interface transition zone with complex microstructure is formed. The microstructure of the transition zone determines performance of the Al2 O3 –TiC/steel diffusion welding joint.

3.1.3 Specimen Preparation and Test Methods for Diffusion Joints To analyze the microstructure and the properties of the Al2 O3 –TiC composite ceramics and steel (Q235, 18-8 or W18Cr4V) diffusion welded joints, the joints were cut firstly. The Al2 O3 –TiC/steel diffusion welded joints were cut into approximately 10 mm × 10 mm × 10 mm specimens perpendicular to the bonded interface by a wire-cutting machine. The schematic diagram of specimen cutting is shown in Fig. 3.4. To observe the microstructure and test the microhardness of the Al2 O3 –TiC/steel diffusion welded joint, specimen cross-sections were smoothed with different grit sizes metallographic sandpaper in turn, then polished mechanically on a polishing machine using diamond abrasive polishing paste with a particle diameter of 2.5 μm and tweed polishing cloth, and etched finally. The etching solution and etching time of Al2 O3 –TiC/ Q235 and Al2 O3 –TiC/18-8 diffusion bonded joints are shown in Table 3.4.

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Taking nine middle spcimens for test analysis Fig. 3.4 Schematic diagram of specimen cutting

Table 3.4 Etching solution and etching time used for diffusion bonded joints Diffusion joint

Etching solution

Etching time/s

Ambient temperature/°C

Al2 O3 –TiC/Q235

4% nitric acid alcohol solution

10

20

Al2 O3– TiC/18-8

aqua regia solution, HCl: HNO3 = 3: 1

8

20

The mechanical properties of ceramic-to-metal diffusion bonded joints are generally evaluated by tensile strength, shear strength and three-point or four-point bending strength, for which there is no unified standard. Room temperature shear strength is used to evaluate the performance of Al2 O3 –TiC composite ceramic and steel diffusion bonded joints, which has the characteristics of simple test method and reliable data. During the test process, the diffusion joint to be tested is put into a special fixture, and the WEW-600t microcomputer screen display hydraulic universal test machine can be used for the shear test. The shear test device is schematically shown in Fig. 3.5. The load applied on the joint at the time of fracture is recorded during the shear test. The shear strength of the joint is calculated according to Eq. (3.1). τ=

P S

(3.1)

where τ is the shear strength, MPa; P is the fracture load, N; and S is the crosssectional area, mm2 . The microstructure of the diffusion welded joint is observed in two ways. One way is to observe the microstructure on the cross-section of the joint, so the specimen is ground perpendicular to the interface direction. The other way is to observe the microstructure on the longitudinal section, so the specimen is ground parallel to the interface direction. The Al2 O3 –TiC composite ceramic/steel diffusion welded joint specimens were cut longitudinally from the central part of the steel (Q235, 18-8 or W18Cr4V) using the wire cutting method, as shown in Fig. 3.6. The specimens were ground along the direction parallel to the interface.

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Pressure Clamp bolt

head

Al2O3-TiC Fixed support

Al2O3-TiC

Fixture

W18Cr4V

Fig. 3.5 Schematic of Al2 O3 –TiC/W18Cr4V joint shear strength test Fig. 3.6 Schematic diagram of the microstructure analysis layer-by-layer in the Al2 O3 –TiC/steel interfacial transition zone

Wire-cutting

Al2O3-TiC Steel

Al2O3-TiC

Steel

Grinding the sample layer by layer along the direction of the arrow

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121

To determine the microstructure properties of the Al2 O3 –TiC composite ceramic and steel diffusion joint, the microhardness near the Al2 O3 –TiC/steel interface was measured perpendicular to the interface using a microhardness tester. The test load is 100 gf (1 gf = 9.80665 × 10–3 N) and the load time is 10 s. For the specimens used for the electron microprobe (EPMA) composition analysis, it was not necessary to perform etching. The surface of the specimen is polished after smoothing by metallographic sandpaper. To determine the generated phases at the Al2 O3 –TiC/steel diffusion interface, typical Al2 O3 –TiC/steel diffusion joint specimens were selected. The cut Al2 O3 – TiC/steel specimens used for phase analysis were peeled layer by layer along the direction parallel to the interface to analyze the phase structure at each layer. For fracture morphology analysis of the specimens, Al2 O3 –TiC/steel (Q235, 18-8 or W18Cr4V) diffusion welded joints were sheared and broken. Then the fracture surface was blown off with a hair dryer to avoid contamination of the fracture face. The fracture morphology was observed under scanning electron microscopy (SEM).

3.2 Diffusion Welding of Al2 O3 –TiC Composite Ceramics and Q235 Steel The Q235 steel specimen size is φ 52 mm × 1.5 mm and the Al2 O3 –TiC composite ceramic specimen size is φ52 mm × 3.0 mm. In the test, transition liquid phase diffusion welding (TLP) with the addition of an intermediate alloy was used to connect the composite ceramic with the Q235 steel. The intermediate alloy is 10 μm Ti/40 μm Cu/10 μm Ti multi-interlayer, and the purpose is to achieve a strong metallurgical bond between Al2 O3 –TiC and Q235 steel. The process parameters of Al2 O3 –TiC composite ceramic diffusion connection with Q235 steel are: the heating temperature of 1130–1160 °C, the holding time of 40–60 min, the pressure of 10–15 MPa and the vacuum degree of 1.33 × 10–4 to 1.33 × 10–5 Pa. The whole heating, pressurization and cooling process is digitally controlled by the American Honeywell DCP-550 instrument Program control. Q235 steel has good plasticity and toughness, and the matrix microstructure is mainly ferrite, as shown in Fig. 3.7.

3.2.1 Interfacial Characteristics and Microhardness of Al2 O3 –TiC/Q235 Diffusion Welded Joint Al2 O3 –TiC/Q235 diffusion welded joint were cut with a wire cutting machine and prepared as a series of metallographic specimens. A 4% alcoholic nitric acid solution was used for etching. Al2 O3 –TiC ceramics were used for transition liquid phase

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(a)

(b)

50μm   (a) OM

(b) SEM

Fig. 3.7 Microstructure of Q235 steel

Al2O3-TiC

Q235

Interfaceial transition zone  

100μm

Interfacial transition zone

Al2O3-TiC

Fig. 3.8 Microstructure of Al2 O3 –TiC/Q235 diffusion welded joint

diffusion welding (TLP) with Q235 steel using Al2 O3 –TiC/Ti–Cu–Ti/Q235 steel/Ti– Cu–Ti/Al2 O3 –TiC stacking. The microstructure of the Al2 O3 –TiC/Q235 diffusion welded joint is shown in Fig. 3.8. The properties of Al2 O3 –TiC composite ceramics and Q235 steel are very different. Al2 O3 –TiC composite ceramics and Q235 steel were diffusion welded with addition of a Ti–Cu–Ti multi-interlayer. The elements in the interlayer diffuse reaction with the elements in the two side substrates (Al2 O3 –TiC composite ceramics and Q235 steel), forming a diffusion welding interfacial transition zone with different microstructural characteristics from the two side substrates. Microhardness tests were performed on Al2 O3 –TiC/Q235 diffusion joints bonded at a holding time of 45 min and different heating temperatures (1130 °C, 1140 °C, 1160 °C). The microhardness measuring position and test results are shown in Figs. 3.9, 3.10, and 3.11. Microhardness test results show that the microhardness of Al2 O3 –TiC/Q235 diffusion joints welded at the same holding time (45 min) and different heating temperatures of 1130 °C, 1140 °C, and 1160 °C have similar variation rules. The microhardness of the interfacial transition zone is much lower than that of Al2 O3 –TiC

0.1

3.2 Diffusion Welding of Al2 O3 –TiC Composite Ceramics and Q235 Steel (a)

123

Al2O3-TiC

(b)

32 1 Transition zone

Al2O3-TiC

Q235

Q235

50μm

(a) Measuring position

(b) Microhardness distribution

Fig. 3.9 Microhardness distribution in the diffusion welded joint of Al2 O3 –TiC/Q235 (1130 °C × 45 min)

Fig. 3.10 Microhardness distribution in the diffusion welded joint of Al2 O3 –TiC/Q235 (1140 °C × 45 min)

ceramics and slightly higher than that of the Q235 steel parent material. It indicates that the there is no high hardness (higher than the hardness of Al2 O3 –TiC ceramics) brittle phase generation in the interfacial transition zone, which is conducive to the improvement of Al2 O3 –TiC/Q235 diffusion welded joint properties. Table 3.5 shows the microhardness near Al2 O3 –TiC/Q235 diffusion interface. The average values of microhardness in the interfacial transition zone are 188HV0.1 , 214HV0.1 and 245HV0.1 at heating temperature of 1130 °C, 1140 °C and 1160 °C respectively. With the increase of the diffusion welding heating temperature, the microhardnes of Al2 O3 –TiC/Q235 interfacial transition zone also increases. This is because the elements diffuse sufficiently and the interfacial reaction intensifies with the increase of the heating temperature. The microhardness values of the reaction products in the transition zone of the Al2 O3 –TiC/Q235 interface increase.

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Fig. 3.11 Microhardness distribution in the diffusion welded joint of Al2 O3 –TiC/Q235 (1160 °C × 45 min)

Table 3.5 Microhardness of Al2 O3 –TiC/Q235 diffusion joint interfacial transition zone at different heating temperatures

Reaction layer

Microhardness at different heating temperatures/HV0.1 1130 °C

1140 °C

1160 °C

Layer 1

120

165

127

Layer 2

209

210

213

Layer 3

236

249

243

Layer 4

–

233

398

Average microhardness

188

214

245

The microhardness of layer 1 in the interface transition zone near the Q235 steel side was lower than that of the other reaction layers at different heating temperatures. The microhardness of the reaction layer near the Al2 O3 –TiC ceramic side was higher. In the Al2 O3 –TiC/Q235 steel diffusion welding process, the diffusion reaction layer grows deeper towards the Q235 steel side and shallower towards the Al2 O3 –TiC ceramic side. Layer 1 near the Q235 steel side is farther away from the original interface (Ti/Cu/Ti interlayer interface). The number of atoms diffusing to layer 1 is less and the interfacial reaction is weaker, so the reaction products microhardness is not high. The reaction layers (layers 3 and layer 4) near Al2 O3 –TiC ceramic side are closer to the original interface. The number of atoms diffusing to the layers is higher and the interfacial reaction is intensified, so the microhardness value of the reaction products is higher. Thus, the microhardness of the transition zone from the Q235 steel side to the Al2 O3 –TiC ceramic side increases gradually, forming a good transition from soft to hard.

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3.2.2 Shear Strength of Al2 O3 –TiC/Q235 Diffusion Joint 10 mm × 10 mm × 8.5 mm specimens (2 specimens for each process parameter) were cut from the Al2 O3 –TiC/Q235 diffusion welded joint under different process parameters using wire cutting method. The specimen surfaces were ground. Then the interfacial shear strength test was performed on a microcomputer screened hydraulic universal testing machine. The load–displacement relationship curve during the shear test is shown in Fig. 3.12.The test and calculation results of the shear strength of Al2 O3 –TiC/Q235 diffusion welded joints are shown in Table 3.6. As can be seen from Fig. 3.12, only elastic deformation occurs during the shear process of the Al2 O3 –TiC/Q235 diffusion welded joint, with no yield point, and no plastic deformation. The load–displacement curve is the typical curve of brittle fracture of ceramic bodies. The load and displacement maintain a good linear relationship before fracture, and the load curve drops abruptly after fracture. The shear test results show (Table 3.6) that the shear strength of the Al2 O3 – TiC/Q235 steel joint increases from 94 to 143 MPa when the heating temperature increases from 1100 °C to 1140 °C at a pressure of 15 MPa (Fig. 3.13). This is due to the more adequate interfacial diffusion reaction with the increase of heating temperature, the tighter metallurgical bond between Al2 O3 –TiC and Q235 interface, and the increase of Al2 O3 –TiC/Q235 interfacial bond strength. When the heating temperature increases to 1180 °C, the shear strength of the Al2 O3 –TiC/Q235 interface decreases to 111 MPa instead. It is due to the microstructure coarsening near the Al2 O3 –TiC/Q235 interface when the heating temperature is too high. This reduces the bond strength. The formation of brittle compounds (such as Fe–Ti phase) near the interface also increases the joint brittleness. The test results of the shear strength at the Al2 O3 –TiC/Q235 diffusion welded interface show that when the heating temperature is controlled at 1140–1160 °C, the holding time is 45–60 min, and the welding pressure is 12–15 MPa, Al2 O3 –TiC/Q235 diffusion welded joints with high shear strength can be obtained. Fig. 3.12 Load versus displacement during shearing of Al2 O3 –TiC/Q235 joints

Brittle fracture

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Table 3.6 Test and calculated results of the shear strength of the Al2 O3 –TiC/Q235 diffusion welded interface Number

Technological parameter (T × t, P)

Shearing area S/mm2

1

1100 °C × 60 min, 15 MPa

10 × 10

2

1100 °C × 60 min, 15 MPa

3

Maximum load F max /kN

Shear strength σ τ /MPa

Average shear strength σ τ /MPa

Fracture type

9.0

90

94

Interface fracture

10 × 10

9.8

98

1120 °C × 60 min, 15 MPa

10 × 10

11.2

112

116

Type I mixed fracture

4

1120 °C × 60 min, 15 MPa

10 × 10

12.0

120

5

1140 °C × 60 min, 15 MPa

10 × 10

13.7

137

143

Type II mixed fracture

6

1140 °C × 60 min, 15 MPa

10 × 10

14.9

149

7

1160 °C × 45 min, 15 MPa

10 × 10

12.7

127

131

Type I mixed fracture

8

1160 °C × 45 min, 15 MPa

10 × 10

13.5

135

9

1180 °C × 45 min, 15 MPa

10 × 10

11.4

114

111

Ceramic fracture

10

1180 °C × 45 min, 15 MPa

10 × 10

10.8

108

Fig. 3.13 Variation of shear strength of Al2 O3 –TiC/Q235 diffusion welded interface with heating temperature

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127

3.2.3 Microstructure of Al2 O3 –TiC/Q235 Diffusion Welded Joint (1) Interface microstructure characteristics The microstructures of Al2 O3 –TiC/Q235 diffusion welded joint were observed using optical microscope (OM) and JXA-840 scanning electron microscope (SEM). The microstructure characteristics of the Al2 O3 –TiC/Q235 diffusion welded joint taken under OM and SEM are shown in Figs. 3.14 and 3.15 respectively. It can be seen that there are obvious diffusion characteristics at the interface between Al2 O3 –TiC composite ceramic and Q235 steel. The microstructure characteristic of the interface transition zone is different from the substrates on both sides. The Al2 O3 –TiC composite ceramic and Q235 steel diffusion interface is tightly bonded, and there are no microscopic pores, cracks and unconnected areas. On the Al2 O3 –TiC composite ceramic side, the boundary of the interfacial transition zone

Al2O3-Ti

B

(b)

A

Q235

B Q235

100μm

A

Al2O3-TiC

(a)

25μm

Fig. 3.14 Microstructure characteristics of Al2 O3 –TiC/Q235 diffusion welded joints by OM

Fig. 3.15 Microstructure characteristics of Al2 O3 –TiC/Q235 diffusion welded joints by SEM

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Fig. 3.16 Columnar crystal microstructure near the diffusion transition zone of Al2 O3 –TiC/Q235 joint

and the Al2 O3 –TiC matrix is flat and continuous. On the Q235 steel side, the boundary of the interfacial transition zone and the Q235 steel matrix is not obvious and mainly consists of fine, discontinuously distributed granular precipitation phases. The diffusion reaction layer grows deeper into the Q235 steel and shallower into the Al2 O3 –TiC composite ceramic, because the atoms diffuse more easily in metal than in ceramic. The intermediate layer diffusion reaction zone A is narrow, but it plays an important role in the reliable connection between Al2 O3 –TiC and Q235 steel. The boundary between the intermediate layer diffusion reaction zone A and the Al2 O3 –TiC ceramic and the boundary between reaction zone A and the diffusion reaction zone B near Q235 steel side are both quiet obvious. There are clear common grains between the diffusion reaction zone B near Q235 steel side and the Q235 steel substrate. The interface exhibits continuous growth with large interfacial interlocking characteristics, as shown in Fig. 3.14. Different from fusion welding, common grains are generally not easily formed at the interface of the joint in the process of diffusion joining, since the base material does not melt. Only a metallurgical diffusion bond with mutual atomic penetration is formed between the intermediate layer (brazing material) and the base material. Al2 O3 –TiC/Q235 diffusion transition zone has obvious common grains (with Q235 steel), and the grain morphology is columnar throughout the interfacial transition zone, as shown in Fig. 3.16. During the diffusion connection process, the microstructure of the Q235 steel matrix near the interface transition zone grows from the original equiaxed crystals to columnar crystals. The microstructure near the interface and the grain growth direction both change. Due to the narrow diffusion connection interface transition zone and slow cooling rate, the grains at the interface of the base material are most suitable as ready surfaces for diffusion connection interface crystallization, which is most favorable for grain growth. The diffusion bonding interface microstructure is easily formed on the basis of the parent phase, which is the epitaxial growth of the microstructure, and grows preferentially into columnar crystals along the heat conduction direction.

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(2) Microstructure of longitudinal section The Al2 O3 –TiC/Q235 diffusion welded joint specimens were cut longitudinally from the central part of the Q235 steel (Fig. 3.6) and the specimens were ground along the direction parallel to the interface. The microstructure of the Al2 O3 –TiC/Q235 diffusion joints were analyzed with layer-by-layer method. The cross-sectional microstructure of the Al2 O3 –TiC/Q235 steel diffusion joints corroded with 4% nitric acid alcohol is shown in Fig. 3.17, where: zone A is the intermediate layer reaction zone, zone B is the Q235 steel side diffusion reaction zone, and the microstructure of layers a to d is shown in Fig. 3.18. The observation and analysis of the longitudinal cross-sectional microstructure can provide an intuitive three-dimensional understanding of the microstructure distribution and characteristics for the interface transition zone of the Al2 O3 –TiC/Q235 diffusion welding joint, which can better analyze each layer microstructure characteristics in the interface transition zone of the Al2 O3 –TiC/Q235 diffusion welding joint. Layer a is located within the intermediate layer reaction zone (zone A). Gray microstructures with smooth border are distributed over the light copper-coloured microstructures. Layer a is mainly the area formed by the diffusion reaction of elements Ti with Cu in the Ti–Cu–Ti interlayer. Electron probe microanalyzer (EPMA) analysis shows that the element Fe in Q235 steel diffuses into the intermediate layer reaction zone (zone A). Layer b is the result of the diffusion reaction of Ti, Cu and Fe on the intersection of the reaction zone of the intermediate layer (zone A) and the reaction zone of the Q235 steel side (zone B). Layer c is located within the Q235 steel side diffusion reaction zone (zone B), adjacent to the intermediate layer reaction zone, and is characterized by bright white massive microstructure distributed along the grain boundaries.

Fig. 3.17 Microstructure analysis with layer-by-layer method for the transition zone of the Al2 O3 – TiC/Q235 joint

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Fig. 3.18 Microstructure analysis with layer-by-layer method for the transition zone of the Al2 O3 – TiC/Q235 joint (layers a to d)

Layer d is also located in zone B, showing a uniformly fine equiaxed crystal microstructure with fine granular microstructure distributed on the grain boundaries. The analysis suggests that the fine equiaxed crystal microstructure of layer d is the result of the elements lattice inner surface diffusion (also called grid diffusion). The diffusion of elements results in the division of the Q235 steel matrix grains, producing a phenomenon similar to the diffusion of grain boundaries. Layers e to g are all located in the Q235 steel side diffusion reaction zone, away from the intermediate layer reaction zone, adjacent to the original Q235 steel microstructure where the elemental diffusion reaction has not occurred, and are the end of the diffusion of Ti elements from the intermediate layer to the Q235 steel side. Influenced by diffusion welding parameters, layers e to g still have similar isometric crystal microstructures with layer d. The difference is as follows. On layer e the black fine dot microstructure uniformly diffuse distribution in the entire reaction layer. On layer f the black dot microstructure precipitation increases, with cluster aggregating or diffusing uniform distribution. Layer g is adjacent to the original Q235 steel microstructure without elemental diffusion reaction, and the microstructure morphology is similar to that of Q235 steel. The morphology and distribution characteristics of the precipitated phase are similar to those of layer e.

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3.2.4 Analysis of Precipitated Phases in the Interface Transition Zone Under scanning electron microscopy (SEM), there are some morphologically different precipitation phases in the interface transition zone of Al2 O3 –TiC/Q235 diffusion welding joint. By analyzing the microzone composition and precipitation phases in the interface transition zone, the microstructure composition of the interface transition zone of Al2 O3 –TiC/Q235 diffusion welding joint and its effect on the joint properties can be obtained. The microstructure characteristics and precipitation phases morphology in Al2 O3 –TiC/Q235 diffusion welding interface are shown in Fig. 3.19.

Fig. 3.19 Microstructure characteristics and precipitation phase morphology in Al2 O3 –TiC/Q235 diffusion welding interface

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Table 3.7 Tested results of composition for the Al2 O3 –TiC/Q235 interface transition zone (wt%) Point

Al

Si

1

0.49

0.60

2

0.43

0.65

–

3

1.15

–

5.97

4

0.47

0.42

–

5

0.53

6

W

Mo

8.30

Ti

Fe

27.07

69.04

27.49 55.03

Cu

Total

2.80

100

68.94

2.49

100

1.05

28.50

100

26.02

67.88

5.21

100

3.81

90.69

4.98

100

29.83

66.82

3.36

100

As seen in Fig. 3.19, the intermediate layer reaction zone A is a region consisting of light and dark gray massive microstructure, with black granular microstructure distributed on the light and dark gray matrix. The microstructure morphology of the Q235 steel side diffusion reaction zone B differs from that of zone A, with obvious grain boundary features. Q235 steel side diffusion reaction zone B is further subdivided into three small regions B1 , B2 , and B3 (Fig. 3.19a). There is a clear boundary between zone A and zone B1 , and the grains in zone B1 grow based on the undulating ready-made surface of zone A. Zone B3 , immediately adjacent to the Q235 steel side, is distributed with fine, diffuse granular precipitation phases. The histomorphology of zone B2 , located between zones B1 and B3 , is a natural transition between these two zones. The microzone compositions and precipitation phases in each region of the Al2 O3 – TiC/Q235 interface transition zone were analyzed using X-ray energy spectroscopy under SEM. The measurement point are shown in Fig. 3.19. The tested results of composition for the Al2 O3 –TiC/Q235 interface transition zone are shown in Table 3.7. From Table 3.7, it can be seen that the light gray matrix (measurement point 1) and gray-black matrix (measurement point 2) in the intermediate layer reaction zone A contain mainly Fe and Ti and a small amount of Cu, and the contents of Fe, Ti and Cu are almost the same in the two morphologically different microstructures. The phase compositions of these two microstructures are mainly α-Fe, Fe-Ti and a small amount of Cu–Ti compound phase. The black granular microstructure of the intermediate layer reaction zone A (measurement point 3) contains mainly Ti, Cu, Mo, W (from Al2 O3 –TiC ceramics) and small amounts of Al, Fe, Ti, Cu, which can easily combine into compounds. According to the Ti–Mo and Ti–W phase diagrams, Mo and W are infinitely miscible with Ti in the liquid state, and Mo and W can be dissolved in β-Ti in the solid state and keep the β-Ti phase to room temperature. Therefore, the main constituent phases of the black granular microstructure (measurement point 3) are the solid solution of Mo, W in β-Ti and the Ti–Cu compounds. The precipitates in the Q235 side diffusion reaction zone B1 (measurement point 4) mainly contain Fe, Ti and a small amount of Cu, and the content of Fe and Ti is almost the same as the content of Fe and Ti in the light gray matrix and gray-black matrix in the intermediate reaction zone A, with a slight increase in the content of Cu. The precipitates in the zone B1 (measurement point 4) are mainly α-Fe, Fe–Ti and a small amount of Cu–Ti compound phase.

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The microzone analysis results of the Q235 steel side diffusion reaction zone B2 (measurement point 5) show that, in addition to Fe elements, the precipitates contain a small amount of Cu and Ti, mainly α-Fe. The granular precipitation in the Q235 steel side diffusion reaction zone B3 (measurement point 6) contains mainly Fe, Ti and a small amount of Cu. It is mainly α-Fe, Fe–Ti and a small amount of Cu–Ti compound phase.

3.2.5 Effect of Process Parameters on the Microstructure of Al2 O3 –TiC/Q235 Diffusion Interface (1) Effect of heating temperature The most important process parameters of diffusion welding are heating temperature, holding time and welding pressure, which influence and restrict each other. The heating temperature T is one of the important parameters affecting the microstructure properties of the transition zone at the Al2 O3 –TiC/steel diffusion welding interface. The microstructure characteristics of Al2 O3 –TiC/Q235 diffusion welded interfacial transition zone at different heating temperatures are shown in Fig. 3.20. As can be seen from the figure, if the holding time (t = 60 min) and pressure (P = 15 MPa) are same, the width of the Al2 O3 –TiC/Q235 diffusion welding interfacial transition zone gradually increases when the heating temperature increases from 1120 °C to 1180 °C. The width of the intermediate layer reaction zone in the transition zone also gradually increases, and the microstructure morphology of the intermediate layer reaction zone changes greatly., This is due to the reaction degree increasing with the improvement of diffusion welding temperature, indicating that there is a surplus of Cu in the intermediate layer. When the temperature increases to 1160 °C, the elements of Ti and Cu in the intermediate layer reaction zone are fully reacted with the elements diffused from the both sides matrixs only remaining a small amount of Cu around the precipitation phase. The reaction products in the intermediate layer reaction zone are uniformly distributed. When the temperature increases to 1180 °C, the reaction layer was thickened due to over-reaction. It should be noted that the presence of high plasticity residual Cu in the reaction zone of the intermediate layer is beneficial for the mitigation of residual stresses in the joints. When the heating temperature was 1100 °C, the width of the interfacial transition zone of Al2 O3 –TiC/Q235 diffusion welding joint was 94 μm. As the heating temperature increased, the atomic diffusion rate increased, the interfacial reaction intensified, and the width of the interfacial transition zone increased. When the heating temperature increased to 1180 °C, the width of the interfacial transition zone of Al2 O3 –TiC/Q235 diffusion welding joint increased to 134 μm. The effect of heating temperature on the width of the transition zone is shown in Fig. 3.21. Based on the measured results, it can be inferred that the width of the interface transition zone of Al2 O3 –TiC/Q235 diffusion welding joint will continue to increase with the increase of the heating temperature.

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Fig. 3.20 Microstructure of Al2 O3 –TiC/Q235 diffusion welding interface transition zone at different heating temperatures

Fig. 3.21 Effect of heating temperature on the width of the transition zone at the Al2 O3 –TiC/Q235 diffusion welding interface

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However, a too high heating temperature will lead to coarsening of the microstructure near the transition zone of the diffusion welding interface. As the heating temperature increases, the interfacial reaction will be further intensified, and brittle intermetallic compounds are easily generated near the diffusion welding interface. Moreover, the heating temperature directly affects the residual stresses in the Al2 O3 –TiC composite ceramic and steel diffusion welded joints, that is, a higher heating temperature will produce larger residual stresses. Therefore, the heating temperature should be controlled under the premise of obtaining good microstructure and properties for Al2 O3 –TiC/Q235 diffusion welded joint. (2) Effect of holding time and pressure The holding time mainly determines the uniformity of element diffusion and the degree of interfacial reaction near the diffusion welding interface. Pressure is also an important parameter for diffusion joining, ensuring plastic deformation of the microscopic bulges on the joining surface, breaking the surface oxide film and inducing elemental diffusion. When the holding time is short (45 min) and the heating temperature is 1140 °C, the element diffusion in the transition zone of Al2 O3 –TiC/Q235 diffusion welding interface is not sufficient, and large black lumps are visible in the reaction zone of the middle layer. With the increase of holding time and pressure, the microscopic contact area of the diffusion welding interface increases, the number of atoms in a thermally activated state near the interface increases, the atomic diffusion distance also increases, the elemental diffusion and interface reaction are more sufficient, and thus the diffusion welding interface transition zone with uniform microstructure is formed. From the summary of the measured results, the effect of holding time on the width of the interfacial transition zone of Al2 O3 –TiC/Q235 diffusion welding joint is shown in Fig. 3.22. The width of the Al2 O3 –TiC/Q235 diffusion welding interfacial transition zone gradually increases with the extension of the holding time. In the initial stage of holding temperature (holding time less than 30 min), the width of the interfacial transition zone increases fast as the holding time increases. When the holding time exceeds 30 min, the width of the interfacial transition zone increases slowly. This is due to the fact that in the initial stage of holding temperature, the holding time has great influence on the diffusion migration of elements. The longer the holding time, the more adequate the diffusion migration of elements. When a certain time is reached, the influence of holding time on the diffusion migration of elements decreases, the diffusion of elements gradually reaches a quasi-equilibrium state, and the Al2 O3 –TiC/Q235 diffusion welding interface transition zone with stable microstructures is gradually formed. The interactions between heating temperature, holding time and pressure jointly affect the microstructure and properties of Al2 O3 –TiC/Q235 diffusion welded joints. In order to obtain Al2 O3 –TiC/Q235 diffusion welded joints with adequate diffusion, good interfacial bonding and excellent microstructure and properties, the appropriate heating temperature, holding time and pressure must be selected in coordination. The test results show that the suitable process parameters of Al2 O3 –TiC/Q235 diffusion welding are: heating temperature T = 1140–1160 °C holding time t = 45–60 min, pressure p = 12–15 MPa.

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Fig. 3.22 Effect of holding time on the width of the transition zone at the Al2 O3 –TiC/Q235 diffusion welding interface

3.3 Diffusion Welding of Al2 O3 –TiC Composite Ceramics with 18-8 Austenitic Steel The tested 18-8 steel is austenitic stainless steel (1Cr18Ni9Ti) with specimen size of φ 52 mm × 1.2 mm. The chemical composition and thermo-physical properties of the 18-8 austenitic steel are shown in Table 3.8.The microstructures of the 18-8 steel are γ-austenite and a small amount of δ-ferrite, as shown in Fig. 3.23. The process parameters for the diffusion connection of Al2 O3 –TiC composite ceramics with 18-8 steel are: heating temperature of 1090–1170 °C, holding time of 45–60 min, pressure of 10–20 MPa, and vacuum degree of 1.33 × 10–4 to 1.33 × 10–5 Pa.

3.3.1 Interfacial Characteristics and Microhardness of the Al2 O3 –TiC/18-8 Diffusion Welding Joint From the microstructure morphology of the Al2 O3 –TiC composite ceramics and 188 steel diffusion bonding joints, it can be seen that there exists diffusion interface transition zone between both side matrixs of Al2 O3 –TiC and 18-8 steel, as shown in Fig. 3.24. It is obvious that there are two distinct interfacial transition zones in the Al2 O3 –TiC/18-8 diffusion bonding joint, which can be divided into 2 regions: the intermediate layer diffusion reaction zone and the 18-8 steel side diffusion reaction zone.

8030

Density/g cm−3

0.50

Specific heat/J (g K)−1

16.0

Heat conductivity/W m−1 ·K−1

≤2.0

≤0.12

GB4237-2007

Thermo-physical properties

2.0

0.11

Measured value

Mn

C

Element

Chemical compositions/wt%

17.0–19.0

18

Cr

16.7

Coefficient of thermal expansion/10−6 ·K−1

≤1.0

0.8

Si

74

Resistivity/10−6 Ω·cm

8.0–11.0

9.5

Ni

Table 3.8 Chemical composition and thermo-physical properties of 18-8 steel

520

Tensile strength σ b /MPa

5(C%–0.02)–0.8

0.6

Ti

40

Elongationδ 5 /%

≤0.03

0.03

S

70

Hardness/HRB

≤0.035

0.03

P

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Fig. 3.23 Microstructure of 18-8 austenitic steel

Fig. 3.24 Schematic representation of the transition zone at the Al2 O3 –TiC/18-8 diffusion welding interface

➀ Intermediate layer diffusion reaction zone: formed by the diffusion reaction of the elements in the Ti–Cu–Ti intermediate layer with the matrix elements on both sides that diffuse into the zone (Zone A). ➁ 18-8 steel side diffusion reaction zone: Located within the 18-8 steel matrix, it is the zone formed by the diffusion of Ti elements in the intermediate layer into the 18-8 steel matrix at a certain distance and diffusion reaction with the elements in 18-8 steel (Zone B). In order to determine the microstructure and properties of Al2 O3 –TiC/18-8 diffusion welding joints, microhardness tests were performed near the diffusion welding

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interface of Al2 O3 –TiC/18-8 obtained at the same holding time (60 min) and different heating temperatures of 1140 °C and 1150 °C. The measuring position of the microhardness and the microhardness distribution are shown in Figs. 3.25 and 3.26, where the horizontal coordinate is perpendicular to the interface direction. The microhardness measurement results showed that the microhardness near the Al2 O3 –TiC/18-8 diffusion bonding interface obtained at different heating temperatures of 1140 °C and 1150 °C and the same holding time (60 min) had a similar variation regulation, that is, the microhardness in the interface transition zone was much lower than that of the Al2 O3 –TiC ceramics and slightly higher than that of the 18-8 stainless steel.

Fig. 3.25 Microhardness distribution in the transition zone of Al2 O3 –TiC/18-8 joint (1140 °C × 60 min)

Fig. 3.26 Microhardness distribution in the transition zone of Al2 O3 –TiC/18-8 joint (1150 °C × 60 min)

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140 Table 3.9 Microhardness of the Al2 O3 –TiC/18-8 transition zone at different heating temperatures

3 Diffusion Welding of Composite Ceramics to Steel Reaction zone

Microhardness at different heating temperatures /HV0.1 1140 °C 1150 °C

Intermediate layer diffusion reaction zone 441

452

18-8 steel side diffusion reaction zone

421

428

Mean microhardness

431

440

Table 3.9 shows the microhardness variation of Al2 O3 –TiC/18-8 diffusion welding interface transition zone at different heating temperatures. It can be seen that the microhardness values in the intermediate layer reaction zone and the 18-8 steel side reaction zone are higher than those at 1140 °C when the heating temperature is 1150 °C. As the diffusion welding heating temperature increases, the interfacial diffusion reaction intensifies and brittle intermetallic compounds are easily generated.

3.3.2 Shear Strength of Al2 O3 –TiC/18-8 Diffusion Joint Wire cutting method was used to cut 10 mm × 10 mm × 8.2 mm specimens from Al2 O3 –TiC/18-8 diffusion welded joint (2 specimens were cut under each process parameter). The surface of the cut down specimens were ground and then shear tested on a microcomputer screened hydraulic universal testing machine. The load– displacement relationship curve during the shear test is shown in Fig. 3.27. The test and calculation results of the shear strength of Al2 O3 –TiC/18-8 joints are shown in Table 3.10. Similar to the load–displacement graphs during shear tests of Al2 O3 –TiC/Q235 diffusion welded joints, the load–displacement graphs during the shear process of Al2 O3 –TiC/18-8 diffusion welded joints also exhibit typical ceramic body brittle fracture characteristics with a good linear relationship between load and displacement before fracture and an abrupt drop in the load curve after fracture. The results of the Al2 O3 –TiC/18-8 diffusion welded interface shear test (Table 3.10) showed that the shear strength of Al2 O3 –TiC/18-8 diffusion welded joints increased from 85 to 125 MPa when the heating temperature was increased from 1090 °C to 1130 °C, as shown in Fig. 3.28. This is due to the fact that as the heating temperature increases, the diffusion reaction of the elements at the Al2 O3 –TiC/18-8 diffusion welding interface is more adequate, the metallurgical diffusion bond at the Al2 O3 –TiC and 18-8 interface is tighter, and the bonding strength of Al2 O3 –TiC/18-8 joint increases. Then when the heating temperature continued to increase from 1130 °C to 1170 °C, the Al2 O3 –TiC/18-8 diffusion welded joint

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Brittle fracture

Fig. 3.27 Load–displacement relationship curve during the shear test of Al2 O3 –TiC/18-8 joints Table 3.10 Test results for the shear strength of the Al2 O3 –TiC/18-8 diffusion welded interface Number

Technological Shearing area Maximum parameter (T × S/mm2 load t, P) F max /kN

Shear strength σ τ /MPa

Average shear strength σ τ /MPa

Fracture type

1

1090 °C × 10 × 10 60 min, 15 MPa

8.2

82

85

Interface fracture

2

1090 °C × 10 × 10 60 min, 15 MPa

8.8

88

3

1110 °C × 10 × 10 60 min, 15 MPa

9.7

97

101

4

1110 °C × 10 × 10 60 min, 15 MPa

10.5

105

Type I mixed fracture

5

1130 °C × 10 × 10 60 min, 15 MPa

12.2

122

125

6

1130 °C × 10 × 10 60 min, 15 MPa

12.8

128

Type II mixed fracture

7

1150 °C × 10 × 10 45 min, 15 MPa

11.3

113

116

8

1150 °C × 10 × 10 45 min, 15 MPa

11.9

119

Type I mixed fracture

9

1170 °C × 10 × 10 45 min, 15 MPa

8.2

82

88

Ceramic fracture

10

1170 °C × 10 × 10 45 min, 15 MPa

9.4

94

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Fig. 3.28 Effect of heating temperature on the shear strength of Al2 O3 –TiC/18-8 diffusion welded interface

shear strength instead decreased to 88 MPa. This is due to the high heating temperature, Al2 O3 –TiC/18-8 interface transition zone microstructure coarsening to reduce the interfacial bonding strength, Al2 O3 –TiC/18-8 diffusion interface transition zone formation of brittle compounds (e.g. Fe–Ti phase) reduces the plasticity of the joint. The shear strength results of Al2 O3 –TiC/18-8 diffusion welded joints show that when the heating temperature is control at 1130–1150 °C, holding time is 45–60 min and welding pressure is 12–15 MPa, Al2 O3 –TiC composite ceramics and 18-8 steel diffusion welding joint with high shear strength can be obtained.

3.3.3 Microstructure of Al2 O3 –TiC/18-8 Diffusion Welded Joint (1) Microstructure characteristics The microstructure of Al2 O3 –TiC/18-8 diffusion welded joints was observed using an optical microscope (OM). Figure 3.29 illustrates the microstructure characteristics of the transition zone of the Al2 O3 –TiC/18-8 diffusion welded interface. Figure 3.29a shows the microstructure characteristics of the Al2 O3 –TiC/18-8 diffusion welded joint before etching. The metallographic specimens of Al2 O3 –TiC/18-8 diffusion welded joint were etched using aqua regia solution (HCl:HNO3 = 3:1), and the microstructure morphology of the Al2 O3 –TiC/18-8 transition zone after etching is shown in Fig. 3.29b.

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Fig. 3.29 Microstructural characteristics of the transition zone at the Al2 O3 –TiC/18-8 diffusion welded interface

The analysis shows that during the diffusion welding of Al2 O3 –TiC/18-8, the elements in the matrix continuously diffuse to the interface. The elements in the Ti– Cu–Ti intermediate layer diffuse into the matrix at both sides while diffusing into each other. When certain conditions are reached, the diffusion reaction occurs between the elements, forming an interfacial transition zone with different microstructure characteristics from the matrix on both sides. The interfacial transition zone between Al2 O3 –TiC composite ceramics and 18-8 steel consists of two regions: the interlayer diffusion reaction zone A (which can be further divided into two zones A1 and A2 ) and the 18-8 steel side diffusion reaction zone B, as shown in Fig. 3.29. It can be seen that there is a clear boundary between zone A and zone B, and the interfaces with the substrate on both sides are also flat and continuous, with no microscopic holes, cracks and other defects. The interface is well bonded. Zone A is formed by the diffusion reaction of the elements in the Ti–Cu–Ti interlayer with the matrix elements coming from both sides by diffusion. The metallic Cu-like color is clearly visible in the microstructure of zone A before corrosion (Fig. 3.29a), which may be the remaining Cu from the reaction (in the form of α-Cu) or a compound formed by Cu with other elements. From the microstructure of zone A after corrosion (Fig. 3.29b), it can be seen that most of zone A1 (Cu-like color area) changed in color under the aqua regia corrosion, while the other areas of zone A1 (except Cu-like color area) and zone A2 did not change much before and after corrosion. Zone B is the area formed by the diffusion reaction of Ti elements from the Ti–Cu– Ti interlayer diffusing into the 18-8 steel matrix with the 18-8 steel matrix elements. The morphology of the precipitated phase is visible in the microstructure before corrosion (Fig. 3.29a). The microstructure after corrosion is seen as fine uniform grainy microstructure distributed in the matrix microstructure and irregular massive microstructure distributed on the matrix microstructure.

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Fig. 3.30 Layer-by-layer microstructure analysis of the transition zone at the Al2 O3 –TiC/18-8 interface

(2) Microstructure of longitudinal section To further observe and analyze the microstructure characteristics of the Al2 O3 – TiC/18-8 diffusion welded interface transition zone, the specimens were cut and ground using the method shown in Fig. 3.6. The Al2 O3 –TiC/18-8 joint was ground layer by layer along the direction parallel to the interface from the 18-8 steel side to the Al2 O3 –TiC ceramic side. The microstructure of the Al2 O3 –TiC/18-8 interface transition zone was observed layer by layer. As shown in Fig. 3.30 (Zone A is the reaction zone of the intermediate layer, Zone B is the reaction zone of the 18-8 steel side), layer a is adjacent to Al2 O3 –TiC ceramic, layer f is adjacent to 18-8 steel. The microstructure is shown in Fig. 3.31, where the marked “corrosion” is the use of aqua regia solution corrosion, marked “ HCl + HNO3 + CH3 COOH corrosion” for the use of a mixture of hydrochloric acid, nitric acid and glacial acetic acid solution (HCl: HNO3 :CH3 COOH = 1:3:4) corrosion morphology, not marked these words for the microstructure morphology after polishing treatment. The layer a is immediately adjacent to the Al2 O3 –TiC ceramic. The light gray matrix is distributed with fine diffuse granular microstructure. This may be the oxide of Ti generated by Ti diffusion into the Al2 O3 –TiC surface to react with Al2 O3 . The Al ions produced by the reaction are dissolved in the Cu–Ti liquid phase. A small amount of spotted light copper color can be seen, indicating the diffusion of Cu into the vicinity of the Al2 O3 –TiC ceramic interface. The layer b is located in the reaction zone of the intermediate layer (zone A). The layer b microstructure consists of morphologically different microstructures: gray microstructure with rounded boundaries are distributed on the light copper-colored matrix, and small spherical purple-copper-colored microstructures are distributed on the gray and light copper-colored matrix. The boundaries of the various microstructures are clearer after corrosion. The layer b microstructure is formed mainly as a result of the diffusion reaction of elements Ti and Cu in the middle layer with elements from the matrix on both sides (small amounts of Al and O in Al2 O3 –TiC and Fe, Cr and Ni in 18-8 steel).

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Fig. 3.31 Layer-by-layer microstructure analysis of the transition zone at the Al2 O3 –TiC/18-8 interface (layers c to f)

The layer c is also located in the reaction zone of the intermediate layer and has a distinctly different microstructure morphology from layer b, being a gray matrix with a diffuse distribution of granular microstructure. The gray matrix of layer c is similar to the gray microstructure with rounded borders of layer b. The results show that both have better corrosion resistance than the other microstructure of layer a and the 18-8 steel side transition zone (B zone), with clearer borders after aqua regia corrosion. EPMA analysis shows that layer c is mainly the result of the diffusion reaction of element Ti in the intermediate layer with elements Cr, Ni, and Fe from 18-8 steel. The layer d is located at the intersection of the reaction zone of the intermediate layer (zone A) and the reaction zone of the 18-8 steel side (zone B), that is, the layer d is the original surface of the 18-8 steel. The layer d has the same fine diffuse granular microstructure as the layer c, but its gray matrix is distributed with white blocky microstructure, which is the same as that of the zone B. That is to say, the layer d presents the interface characteristics of the diffusion connection between the Ti–Cu–Ti intermediate layer and the 18-8 steel. The layer e represents the histomorphology of the 18-8 steel side reaction zone (zone B). The small white massive microstructure is uniformly distributed in the gray matrix (uncorroded). After aqua regia corrosion, the original white small massive microstructure becomes black and gray, while the original gray matrix becomes

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white, showing as the microstructure morphology of the layer e in Fig. 3.31. After another aqua regia corrosion (secondary corrosion), the grain boundary was faintly visible, with a small amount of grain boundary precipitation phase. A mixture of hydrochloric acid, nitric acid and glacial acetic acid was used to corrode the layer e. It was found that there was another microstructure in the layer e that was uniformly distributed in the matrix in the form of strips or blocks. The original gray matrix became light gray, and the white lumpy microstructure became bright white. That is, the microstructure morphology of the layer e was white small block microstructure and strips of lumpy microstructure uniformly distributed in the matrix. The EPMA analysis shows that the layer e (i.e. the reaction zone of 18-8 steel side) is mainly formed by the intermediate layer element Ti diffusing into 18-8 steel as the active element interaction with Fe, Cr and Ni in 18-8 steel. The result of elemental interaction forms a different microstructure morphology from that of 18-8 steel matrix. The layer f is located at the front of the 18-8 steel side reaction zone (zone B) immediately adjacent to the unchanged 18-8 steel. The layer f is faintly visible as fine dotted microstructure and aggregates at grain boundaries. The layer f is the end of Ti diffusion into 18-8 steel. As only a small amount of Ti diffusing into layer f, Ti aggregates mainly at grain boundaries because grain boundaries are fast diffusion channels.

3.3.4 Analysis of Precipitated Phases in the Interface Transition Zone The backscattered electron images of the interfacial transition zone of Al2 O3 –TiC/188 diffusion welded joint are shown in Figs. 3.32 and 3.33. The interfacial transition zone has different backscattered electron images from the substrates on both sides. Point composition analysis was performed on the intermediate layer reaction zone (zone A). The point composition analysis locations of the wave spectrum (measurement points 1, 2, 3, and 4 in Fig. 3.32) and the energy spectrum (points A, B, C, D, and E in Fig. 3.33) are shown in Figs. 3.32 and 3.33 respectively. The region A1 , immediately adjacent to the Al2 O3 –TiC ceramic, consists of four kinds of microstructures depending on the backscattered electron image morphology: large dark grey microstructure (measurement point 1 in Fig. 3.32), black granular microstructure (measurement point 2 in Fig. 3.32), white microstructure (point B in Fig. 3.33) and grey microstructure (measurement point 3 in Fig. 3.32). The results of the energy spectrum analysis of the interfacial phase components are shown in Table 3.11. The large dark gray microstructure (measurement point 1 in Fig. 3.32) contains mainly Ti, Fe, Ni and small amounts of Cr and Cu. The energy spectrum analysis shows that the dark gray microstructure contains some amount of O elements. Among them, Ti is from the Ti–Cu–Ti interlayer, Fe, Ni, and Cr are from the 18-8 steel matrix.

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Fig. 3.32 Backscattered electron image of the transition zone at the Al2 O3 –TiC/18-8 diffusion welded interface (I)

Fig. 3.33 Backscattered electron image of the transition zone at the Al2 O3 –TiC/18-8 diffusion welded interface (II)

It indicates that the 18-8 steel matrix elements dissolved out and diffused into the Cu– Ti liquid phase in the interlayer during the diffusion welding process, and diffused with the liquid phase to the vicinity of the ceramic matrix. O is from the Al2 O3 –TiC ceramic, indicating that the active element Ti “captures” the O in Al2 O3 –TiC ceramic and combines with O to form the corresponding reaction products retained in the intermediate layer reaction zone. The analysis shows that the dark grey microstructure

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Table 3.11 Energy spectrum analysis of interfacial phase composition (at. %) Position

C

O

Al

Ti

Cu

Ni

Cr

Fe

Total

A

—

60.82

0.59

27.90

2.71

4.40

2.91

Bal

100

B

—

14.34

4.95

8.09

68.46

2.34

1.03

Bal

100

C

42.40

—

—

54.88

2.35

—

—

0.37

100

D

46.58

—

—

48.70

1.30

0.40

0.45

2.57

100

E

—

—

0.98

34.31

—

—

7.83

56.88

100

in the zone A is Ti oxides, Fe–Ti–O compounds, Ni–Ti–O compounds, Ti compounds with Ni, Cr, Cu, etc. The black granular microstructure (measurement point 2 in Fig. 3.32) is TiC containing mainly Ti and C. The white microstructure (point B in Fig. 3.33) contains mainly Cu, O, Ti and small amounts of Al, Ni, Cr. It is most likely residual Cu, Cu–Ti and Cu–Ti–O compounds etc. in the form of α-Cu. The grey microstructure (measurement point 3 in Fig. 3.32) contains mainly Fe, Ti, Cr, Ni and small amounts of other elements. It may form γ-solid solution of Cr, Ni in Fe and compounds of Ti–Cr, Ti–Ni, etc. The zone A2 is adjacent to the surface of 18-8 steel. There is black granular microstructure distributed on the grey matrix. The energy spectrum analysis shows that the fine black granular microstructure (point D in Fig. 3.33) contains mainly Ti and C as the black granular microstructure in zone A1 . The atomic content ratio of Ti and C is close to 1:1. So the black granular microstructure is TiC phase. The black granular microstructure in zone A2 is clustered in bands on the gray matrix. The analysis concluded that the TiC phase precipitated in the A1 zone and A2 zone is not related to the enhanced phase TiC in Al2 O3 –TiC ceramics. The TiC phase precipitated in the intermediate layer reaction zone is the combination of element Ti in the intermediate layer and C in the 18-8 steel matrix diffusing into the intermediate layer reaction zone. The gray matrix microstructure in the zone A2 (measurement point 4 in Fig. 3.32) has a similar backscattered electron image morphology to that in the A1 zone (measurement point 3 in Fig. 3.32). Measurement point 4 has slightly lower content of major elements Ti and Ni and slightly higher content of Fe and Cr compared with measured point 3. The gray microstructures may be γ-solid solution of Cr and Ni in Fe and compounds of Ti–Cr and Ti–Ni, etc., with only slightly different content of each constituent phase. It can be seen that in the intermediate reaction zone, the Ni content in the precipitated phase, except for the TiC phase, is high, close to or even exceeding the Ni content of the 18-8 steel parent material, while the Cr content of the precipitated phase in the intermediate reaction zone is much lower than that of the 18-8 steel parent material. From the Cu–Ni phase diagram, it can be seen that Cu and Ni are infinitely interdigitated in both the liquid and solid states. In the diffusion welding process, once the Cu–Ti liquid phase appears, Ni is easily dissolved in the Cu–Ti liquid phase and diffuses to the whole intermediate reaction zone with the liquid phase.

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The 18-8 steel side diffusion reaction zone (Zone B) is actually a zone formed by the diffusion “loss” of elements from the 18-8 steel matrix into the intermediate layer reaction zone, where Ti elements from the intermediate layer diffuse into the 18-8 steel matrix. Depending on the morphology of the backscattered electron image (different gray levels of the composition image), it contains four main types of microstructure (measurement points 5, 6, 7 and 8 in Fig. 3.32). The measurement point 5 is immediately adjacent to the intermediate layer reaction zone. The measurement point 6 represents the darker grayscale microstructures in the reaction zone on the 18-8 steel side. The measurement point 7 represents the darkest grayscale microstructures in the composition image in the reaction zone on the 18-8 steel side. And measurement point 8 represents the bright white microstructures in the reaction zone on the 18-8 steel side. Measurement points 5 and 7 contain the same types of elements, mainly Fe, Cr, Ti and Ni, and the content of each element is not very different. Measurement point 6 contains mainly Fe, Cr and Ni, and is close to the composition of the 18-8 steel base material. Measurement point 8 contains mainly Fe and Cr, and its Ni content is lower than that of the parent material. The Ti content of measurement points 5 and 7 is higher than that of measurement points 6 and 8. The various microstructures with different backscattered electron images in the reaction zone of 18-8 steel side have higher Fe content, which is close to the Fe content of the 18-8 steel parent material, and the various precipitation phases (Fe–Ti, Cr–Ti, Ni–Ti) are distributed on the γ-Fe matrix.

3.3.5 Effect of Process Parameters on the Microstructure of the Al2 O3 –TiC/18-8 Diffusion Interface (1) Effect of heating temperature The influence of heating temperature on the diffusion process is significant, with small changes in heating temperature producing large changes in the diffusion rate. The heating process of Al2 O3 –TiC/steel diffusion welding is accompanied by a series of physical, chemical, mechanical and metallurgical changes that directly or indirectly affect the Al2 O3 –TiC/steel diffusion welding process and the quality of the joint. The microstructure characteristics of the transition zone of the Al2 O3 –TiC/18-8 diffusion welded interface at different heating temperatures are shown in Fig. 3.34. As seen in the figure, for Al2 O3 –TiC/18-8 diffusion welded joints, when the holding time (45 min) and the connection pressure (P = 12 MPa) are same, the width of the interfacial transition zone of Al2 O3 –TiC/18-8 diffusion welding joint gradually increases with the increase of the heating temperature. When the heating temperature is 1110 °C, the uneven element diffusion and insufficient interfacial reaction are clearly visible in the reaction zone of the middle layer, and the interfacial transition zone is relatively narrow. As the heating temperature increases, the elements

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diffuse uniformly and react adequately, and the width of the interfacial transition zone increases, especially the width of the middle layer reaction zone increases more obviously. When the heating temperature was increased to 1170 °C, the diffusion rate and the reaction degree of the elements were intensified due to the higher temperature, which further widened the interfacial transition zone. The effect of heating temperature on the width of the transition zone at the Al2 O3 – TiC/18-8 diffusion welding interface obtained from the measured results is shown in Fig. 3.35. When the heating temperature was 1090 °C, the width of the transition zone at the Al2 O3 –TiC/18-8 diffusion welding interface was 73 μm. When the heating temperature was increased to 1170 °C, the width of the transition zone at the Al2 O3 – TiC/18-8 joint increased to 109 μm. (2) Effect of holding time and pressure. The microstructure characteristics of the Al2 O3 –TiC/18-8 diffusion welding interface transition zone at a heating temperature of 1150 °C and at different holding times and welding pressures are shown in Fig. 3.36. When the heating temperature is 1150 °C and the holding time is shorter (30 min), the elemental diffusion inhomogeneity is clearly visible in the intermediate layer reaction zone of the Al2 O3 –TiC/18-8 joint,

Fig. 3.34 Microstructure of Al2 O3 –TiC/18-8 diffusion welding interface transition zone at different heating temperatures

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Fig. 3.35 Effect of heating temperature on the width of the transition zone at the Al2 O3 –TiC/18-8 diffusion welding interface

see Fig. 3.36(a). This is due to the shorter holding time, the smaller pressure, the smaller microscopic contact area of the Al2 O3 –TiC/18-8 diffusion welding interface, the number of atoms in the thermally activated state is less and the distance of element diffusion is shorter, so the interface reaction is not sufficient The effect of holding time on the width of the transition zone at the Al2 O3 –TiC/188 diffusion welding interface obtained from the actual results is shown in Fig. 3.37. The test shows that the holding time should not be too long. If the holding time is too long, the microstructure near the Al2 O3 –TiC/18-8 diffusion welding interface is

Fig. 3.36 Microstructure of Al2 O3 –TiC/18-8 diffusion welding interface transition zone at different holding times and pressures

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Fig. 3.37 Effect of holding time on the width of the transition zone at the Al2 O3 –TiC/18-8 diffusion welding interface

easy to coarsen. The degree of interfacial reaction is also intensified, easy to generate a large number of brittle precipitates or compounds, resulting in joint embrittlement. Therefore, the holding time of Al2 O3 –TiC/18-8 diffusion welding process should be controlled to ensure that the diffusion welding interface transition zone with a certain width can be obtained, but also not to make the interface transition zone microstructure precipitation embrittlement.

3.4 Diffusion Welding of Al2 O3 –TiC Composite Ceramics and W18Cr4V High-Speed Steel 3.4.1 Diffusion Process Characteristics and Specimen Preparation W18Cr4V is a tungsten-based general-purpose high-speed steel (HSS) with high tensile strength, thermal hardness, and wear resistance. It can be used to manufacture a variety of cutting tools, such as turning tools, planing tools, milling tools, etc. It is not suitable for manufacturing large-section and thermo-plastic forming tools. The chemical composition, thermo-physical properties, and mechanical properties of W18Cr4V high-speed steel are shown in Table 3.12. The microstructure of W18Cr4V steel consists of tempered martensite, a small amount of residual austenite and white carbide particles, etc., as shown in Fig. 3.38.

17.5–19.0

0.70–0.80

8.70

Density/g cm−3

3.80–4.40

≤0.30

27.21

10.4

63–66

30–35

0.3

≤0.030

S

2500–3500

Bending strength/MPa

0.10–0.40

Mn

Poisson ratio

0.20–0.40

Si

Impact toughness /J cm−2

1.00–1.40

V

Coefficient of Hardness/HRC thermal expansion/10–6 K−1

Cr

Mo

Heat conductivity/W m−1 K−1

Thermo-physical properties

W

C

Chemical compositions/wt%

Table 3.12 Chemical composition, thermophysical properties and mechanical properties of W18Cr4V HSS

225–230

Modulus of elasticity /GPa

≤0.030

P

3.4 Diffusion Welding of Al2 O3 –TiC Composite Ceramics and W18Cr4V …

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Fig. 3.38 Microstructure of W18Cr4V HSS

An important process measure in the diffusion connection of Al2 O3 –TiC composite ceramics and W18Cr4V steel is the use of an active interlayer, one to wet the ceramics and control the interfacial reaction, then to mitigate the stresses caused by the difference in physical and mechanical properties between the ceramics and the steel. For the Al2 O3 –TiC/W18Cr4V diffusion connection, a Ti–Cu–Ti interlayer is selected. The Ti in it is the active metal, which can wet the Al2 O3 –TiC ceramic and W18Cr4V steel well and react with the base materials to form a reaction layer during the joining process. Cu in the intermediate layer is a soft metal with good plasticity and low yield strength, which can relieve the stress in the joint through plastic deformation and creep deformation. The thickness of Ti and Cu layers in the Ti–Cu–Ti intermediate layer has a great influence on the performance of diffusion welded joints. The Ti layer is too thin and cannot fully react with Al2 O3 –TiC ceramics and W18Cr4V steel to form a continuous reaction layer, and the joint is easily disconnected from the interface. The Ti layer is too thick and will form an excessively thick reaction layer containing reaction products such as TiC, Ti3 Al, TiO, high hardness and brittleness, which increases the stress at the interface of the joining process and also reduces the strength of the joint. The use of an interlayer with Cu reduces the stress in the joint compared to the use of a Ti interlayer alone. The greater the thickness of the Cu layer, the more pronounced the stress reduction. However, when the Cu layer is too large, the interface transition zone formed is mainly a Cu solid solution with lower strength, which also reduces the strength of Al2 O3 –TiC/W18Cr4V diffusion connected joints. Al2 O3 –TiC/W18Cr4V diffusion joint with a shear strength of 154 MPa was obtained using Ti–Cu–Ti interlayer at a heating temperature of 1130 °C, a holding time of 45 min, and a joining pressure of 15 MPa, with a mixed fracture in shear fracture. Ti–Cu–Ti interlayer was used for diffusion joining of Al2 O3 –TiC/W18Cr4V steel. The interlayer materials were Cu foil and Ti powder with purity above 99.9%, where the Cu foil was 60–100 μm thick and the Ti powder particle size was 200–250 mesh.

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According to the phase diagram of Cu–Ti binary alloy, the minimum eutectic temperature of Cu–Ti binary alloy is 875 °C. According to the characteristics of diffusion joining and the nature of Al2 O3 –TiC ceramics and W18Cr4V steel, 1130 °C was chosen as the reference joining temperature. The diffusion joining process parameters range from 1080 to 1160 °C for heating temperature, 30 to 60 min for holding time, 10 to 20 MPa for pressure, and 1.33 × 10–4 to 1.33 × 10–5 Pa for vacuum degree. Experimental analysis of Al2 O3 –TiC/W18Cr4V diffusion joint properties and interfacial microstructure was carried out by first cutting and preparing specimens. The wire-cut diffusion joint specimens were ground off the sharp corners and burrs with a diamond grinding wheel and polished with metallographic sandpaper of different grit sizes from coarse to fine. Since the hardness of Al2 O3 –TiC composite ceramics is much greater than that of W18Cr4V steel, the specimen grinding process tends to lead to a high degree of unevenness on both sides of the specimen, which affects the observation of metallographic microstructure. In the process of specimen grinding, attention is paid to the appropriate strength and vertical interface direction grinding. After the specimens are ground on metallographic sandpaper, they are placed on a polishing machine and polished with Cr2O3 polishing powder solution until the surface is smooth and free of scratches. The interfacial microstructure characteristics and performance of the Al2 O3 –TiC/W18Cr4V diffusion joint are related to the elements contained in the interlayer material and the diffusion joining process. When diffusion joining Al2 O3 –TiC composite ceramics and W18Cr4V HSS using Ti–Cu– Ti intermediate layer, the Ti–Cu–Ti intermediate layer will melt to form a liquid phase to react with the base materials on both sides, forming interface microstructure with different properties from those of the materials being joined.

3.4.2 Interfacial Characteristics of Al2 O3 –TiC/W18Cr4V Diffusion Welded Joint Microstructural features near the Al2 O3 –TiC/W18Cr4V diffusion interface were observed by optical microscopy and scanning electron microscopy (SEM). Figure 3.39 illustrates the microstructural characteristics of the Al2 O3 –TiC and W18Cr4V diffusion joint observed under metallurgical microscopy. As can be seen from the figure, the Al2 O3 –TiC composite ceramic is tightly bonded with W18Cr4V steel at the interface of the diffusion joint at a heating temperature of 1130 °C, a holding time of 45 min, and a joining pressure of 20 MPa, and no defects such as poor bonding and voids are observed. The Ti–Cu–Ti intermediate layer between Al2 O3 –TiC and W18Cr4V has been completely melted down and forms a reaction layer with clearly different microstructure from the base materials on both sides. The edge of the reaction layer near the Al2 O3 –TiC ceramic side is relatively straight, while the interface near the W18Cr4V steel side is slightly undulating, which may be caused by the uneven surface of the W18Cr4V steel or the uneven spreading of the Ti powder in the intermediate layer before joining.

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Fig. 3.39 Microstructure characteristics near the interface of Al2 O3 –TiC/W18Cr4V diffusion welding joint (1130 °C × 45 min, P = 20 MPa)

3.4.3 Shear Strength of the Al2 O3 –TiC/W18Cr4V Diffusion Welded Interface The shear specimens were cut from the Al2 O3 –TiC/W18Cr4V diffusion joints obtained at different process parameters using wire cutting method. The specimen surface was ground and then clamped with a special fixture on a WEW-600t microcomputer screen display hydraulic universal test machine for shear testing. At the beginning of the shear test process, the displacement increased linearly with the increase of the load. When the load reached the maximum value, the displacement decreased rapidly and the joint fractured rapidly. It indicated that the plastic deformation of the joint was small and the joint underwent brittle fracture. The test results of the Al2 O3 –TiC/W18Cr4V diffusion joint interface shear strength are shown in Table 3.13. When the heating temperature increased from 1080 °C to 1130 °C and the joining pressure increased from 10 to 15 MPa, the Al2 O3 – TiC/W18Cr4V diffusion joint interface shear strength increased from 95 to 154 MPa. This is due to the fact that with the increase in heating temperature the intermediate layer reacts more fully with the base materials on both sides, and a good metallurgical bond is formed near the interface. The increased pressure can make the interface contact tighter and provide more channels for element diffusion. However, when the heating temperature was increased to 1160 °C, the Al2 O3 –TiC/W18Cr4V diffusion joint interface shear strength started to decrease instead, with a shear strength of 141 MPa. This is due to the formation of a thick TiC reaction layer by the interfacial reaction at too high a temperature, which reduces the strength of the joint.

3.4 Diffusion Welding of Al2 O3 –TiC Composite Ceramics and W18Cr4V … Table 3.13 Al2 O3 –TiC/W18Cr4V diffusion joint interface shear strength

Process parameters (T × t, P)

Shear area/mm2

Maximum average load/kN

157 Shear strength/MPa

1080 °C × 10 × 10 45 min, 10 MPa

9.53

95

1100 °C × 12 × 10 45 min, 10 MPa

14.57

121

1130 °C × 10 × 10 45 min, 15 MPa

15.40

154

1160 °C × 10 × 10 45 min, 15 MPa

14.10

141

3.4.4 Effect of Process Parameters on the Microstructure of the Interface Transition Zone During diffusion joining, process parameters (heating temperature, holding time, joining pressure, etc.) are key factors in determining the microstructure properties of the transition zone at the Al2 O3 –TiC/W18Cr4V interface. In order to obtain Al2 O3 – TiC/W18Cr4V diffusion joined joints with good interfacial bonding, diffusion joining processability tests were performed on Al2 O3 –TiC ceramics and W18Cr4V steel using different process parameters. The interfacial microstructure characteristics were observed and analyzed by metallographic microscopy and scanning electron microscopy (SEM). (1) Influence of heating temperature The heating temperature T is the most important process parameter for diffusion joining and determines the extent of diffusion of elements and interfacial reactions. Figure 3.40 illustrates the microstructure characteristics of the Al2 O3 –TiC/W18Cr4V diffusion interface transition zone at different heating temperatures. As can be seen from the figure, the higher the heating temperature is, the more adequate the interfacial reaction is, the width of the Al2 O3 –TiC/W18Cr4V interfacial transition zone gradually increases, and the microstructure of the interfacial transition zone gradually coarsens. The measured values of the width of Al2 O3 –TiC/W18Cr4V interfacial transition zone at different heating temperatures with the same holding time (t = 45 min) are listed in Table 3.14. From the table, it can be seen that the width of Al2 O3 – TiC/W18Cr4V diffusion joint interfacial transition zone is about 32 μm when the heating temperature is 1080 °C, and increases to 72 μm when the heating temperature increases to 1160 °C. The effect of heating temperature on the width of Al2 O3 – TiC/W18Cr4V interfacial transition zone obtained from the measured results is shown in Fig. 3.41. Based on the measured results, it is foreseen that the width of the transition zone at the Al2 O3 –TiC/W18Cr4V diffusion joint interface will increase by continuing to increase the heating temperature (Fig. 3.41). However, too high heating temperature

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Fig. 3.40 Microstructure of the Al2 O3 –TiC/W18Cr4V interface transition zone at different heating temperatures

Table 3.14 Width of Al2 O3 –TiC/W18Cr4V interface transition zone at different heating temperatures (t = 45 min) Heating temperature/°C

1080

1100

1130

1160

Width/μm

32

42

57

72

will lead to coarsening of the microstructure near the diffusion joint interface, which will have a detrimental effect on the microstructure and mechanical properties of the diffusion welded joints. Therefore, the heating temperature should be limited. (2) Effect of holding time and pressure The element diffusion distance is proportional to the square root of the diffusion time, i.e., it conforms to the parabolic law. Therefore, the longer the holding time is, the larger the element diffusion distance is. When the holding time is 30 min and the pressure is 10 MPa, the diffusion distance of the elements is very short and the diffusion reaction is not sufficient. Ti is mainly aggregated at the interface between the base material and the intermediate layer on both sides. The distribution of Ti is not uniform in the interface transition zone. When the holding time is 60 min

3.4 Diffusion Welding of Al2 O3 –TiC Composite Ceramics and W18Cr4V …

159

80

Width /µm

60 40 20 0 1060

1080

1100 1120 1140 1160 Heating temperature /°C

1180

Fig. 3.41 Effect of heating temperature on the width of Al2 O3 –TiC/W18Cr4V interface transition zone

and the joining pressure is 15 MPa, the diffusion distance of the elements is greatly improved and the elemental reaction is more adequate. An interfacial transition zone with uniform microstructure is formed at the Al2 O3 –TiC/W18Cr4V diffusion joint interface. The aggregation of Ti at the interface of the base material and the interfacial transition zone is no longer obvious. The effect of pressure on the microstructure of the transition zone of the Al2 O3 – TiC/W18Cr4V interface is mainly manifested in the early stage of the diffusion welding to promote the close contact between the interfaces. After the Ti–Cu–Ti intermediate layer melts to form the liquid phase, pressure mainly improves the wettability of the Cu–Ti liquid phase to the Al2 O3 –TiC ceramic and the W18Cr4V base material and promotes the interfacial diffusion, so that the intermediate layer and both sides of the base material reach atomic-level contact and form a large number of diffusion channels. Therefore, the pressure should be increased appropriately to promote interfacial diffusion.

3.4.5 Crack Extension and Fracture Characteristics at the Al2 O3 –TiC/W18Cr4V Diffusion Welded Interface There are inherent differences between Al2 O3 –TiC composite ceramics and W18Cr4V HSS. The presence of large stresses near the Al2 O3 –TiC/W18Cr4V interface can affect the bond strength of the diffusion joint, which can lead to cracking or

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fracture of the joint in severe cases. The fracture of the joint may occur at the interface between the base material and the reaction layer, in the reaction layer, in the base material, or at the interface between the reaction layer and the reaction layer. The analysis of the crack expansion path and fracture can determine the bonding properties between the various zones of the diffusion joint. (1) Crack extension at the diffusion interface Through the analysis of the shear fracture of Al2 O3 –TiC/W18Cr4V diffusion welded joints, it is found that the fracture of Al2 O3 –TiC/W18Cr4V diffusion joint is dominated by interfacial fracture and near-interfacial fracture. Fracture generally occurs from the interface, which extend along the interface or extend into the matrix near the interface to cause structural damage. Or cracks sprout near the interface, and extend along the near-interface or extend into the interface to cause fracture. The strength of ceramic-to-metal joints is jointly determined by the interfacial strength (the adhesion strength of the reaction layer to the ceramic) and the residual stresses generated during cooling process. It was found that the shear strength and shear fracture location of diffusion welded joints are related to the path of crack extension. According to the crack extension path, the fracture forms of Al2 O3 –TiC/W18Cr4V joints can be broadly classified into four categories. (i)

Interfacial fracture: the joint fractured at the Al2 O3 –TiC/W18Cr4V interface, see Fig. 3.42a. (ii) Ceramic (matrix) fracture: the fracture starts at the interface, arc to the Al2 O3 – TiC ceramic side, and then returns to the Al2 O3 –TiC/W18Cr4V interface, see Fig. 3.42b. (iii) Type I mixed fracture (near the interface): the fracture starts at the interface and then extends towards the Al2 O3 –TiC side, where the fracture occurs in the Al2 O3 –TiC ceramic, see Fig. 3.42c. (iv) Type II mixed fracture (near interface): the crack undergoes multiple transitions from interface → Al2 O3 –TiC → interface → Al2 O3 –TiC during extension, see Fig. 3.42d. The interfacial fracture mainly refers to the parallel fracture of the Al2 O3 – TiC/W18Cr4V diffusion welded joint along the interface between the reaction layer and the Al2 O3 –TiC ceramic. The fracture surface is flush, and the corresponding joint bonded strength is very low. The weak connection between the Al2 O3 –TiC ceramic and the interfacial reaction layer can be characterized from the fracture morphology. This fracture pattern was observed when the interface between Al2 O3 –TiC ceramics and the interfacial reaction layer was poorly bonded, there was an unwelded and weak joining. The interfacial strength was lower than that of the Al2 O3 –TiC ceramics in the near-interface region. Interfacial fracture occurred at low joining temperature or short holding time. Because the Ti–Cu–Ti interlayer did not react sufficiently with Al2 O3 –TiC ceramics at low joining temperature or short holding time, the reaction layer was very thin and did not form a good bond. The fracture morphology of the Al2 O3 –TiC/W18Cr4V diffusion welded interface at the joining temperature of 1080 °C and holding time

3.4 Diffusion Welding of Al2 O3 –TiC Composite Ceramics and W18Cr4V …

(a)

Al2O3-TiC

(b)

Crack expansion path

Al2O3-TiC Crack expansion path

W18Cr4V

W18Cr4V (a) Interfacial fracture (c)

161

Crack expansion path

Al2O3-TiC

(b) Ceramic (matrix) fracture (d)

Al2O3-TiC Crack expansion path

W18Cr4V

W18Cr4V

(c) Type I mixed fracture

(d) Type II mixed fracture

Fig. 3.42 Illustration of crack expansion path in Al2 O3 –TiC/W18Cr4V joint

of 30 min is shown in Fig. 3.43a. The fracture occurred completely at the Al2 O3 – TiC/W18Cr4V diffusion bonded interface, and there is a metallic luster of Cu on the Al2 O3 –TiC composite ceramics surface. Ceramic (matrix) fracture means that the fracture occurs at the Al2 O3 –TiC ceramic near the interface rather than at the joint. This form of fracture indicates that a strong connection is formed between the Al2 O3 –TiC ceramic and the interfacial reaction layer and that the interfacial strength is higher than the strength of the Al2 O3 –TiC ceramic which is weakened due to the action of stress. Ceramic (matrix) fracture occurs when the crack starts from the end of the interface (the intersection of the interface and the outer surface of the specimen) at an angle to the interface, then extends in the near-interface Al2 O3 –TiC ceramics, and finally returns to the Al2 O3 –TiC/reaction layer interface. The whole fracture path is arcshaped. Most of this fracture occurs in Al2 O3 –TiC ceramics. The crack extension path is generally consistent with the stress distribution in the Al2 O3 –TiC/W18Cr4V diffusion joint. This phenomenon occurs mainly in the case of too long holding time. The interface fracture morphology for the Al2 O3 –TiC/W18Cr4V joint at a heating temperature of 1160 °C and a holding time of 60 min is shown in Fig. 3.43b. The fracture starts at the interface, extends a section along the interface and then extends into the Al2 O3 –TiC ceramic, where some of the ceramic peels off, and finally returns

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(a)

Interface

(b)

Al2O3-TiC

Interface Interface (a) 32×

(b) 8× (d)

(c)

Al2O3-TiC

Al2O3-TiC

Interface

Interface Al2O3-TiC

(c)

32×

(d) 16×

Fig. 3.43 Macroscopic fracture morphology of the Al2 O3 –TiC/W18Cr4V interface

to the interface. Although this fracture is partially fractured in the ceramic, the joint strength is much lower than that of the ceramic. The analysis of the microzone composition of the shear fracture at the ceramic fracture interface of the Al2 O3 –TiC/W18Cr4V joint shows that the shear fracture contains mainly Al and Ti elements, which indicates that the fracture of the Al2 O3 – TiC/W18Cr4V joint occurs mainly at the interface near the Al2 O3 –TiC ceramic side. This is due to the poor toughness of the Al2 O3 –TiC ceramic. After the crack opens from the Al2 O3 –TiC/W18Cr4V interface under shear load, the crack is very easy to extend along the Al2 O3 –TiC ceramic which has poor fracture toughness. Mixed fracture occurs at the interface and within the Al2 O3 –TiC ceramic near the interface. The strength of the mixed fracture joint is higher than that of the interfacial fracture and ceramic fracture. This is due to the tortuous fracture path and the need to consume more energy for crack extension. Type I mixed fracture starts from the interface, and then extends at an angle into the Al2 O3 –TiC ceramic and eventually fractures in the Al2 O3 –TiC ceramic, see Fig. 3.43c. This fracture is prone to occur when the joining process is close to the optimum process parameters. The joining strength of the joint is also higher. This fracture occurs in Al2 O3 –TiC/W18Cr4V joints when the joining temperature T =

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1130 °C and the joining time t = 30 min. After the joint fractures, the Al2 O3 –TiC ceramic of the joint are broken into many small pieces, indicating that the interface reaction of this joint is very adequate and the interface formes a good connection. The crack of type II mixed fracture turns from interface to Al2 O3 –TiC to interface to Al2 O3 –TiC for many times in the process of propagation, that is, the path of crack extension is a fold line. Several small pieces of Al2 O3 –TiC with scattered distribution are adhered to the golden yellow substrate on the metal side surface of the fracture. Type II mixed fracture corresponds to the highest Al2 O3 –TiC/W18Cr4V joint strength. After the shear fracture occurres, some of the Al2 O3 –TiC ceramics broke into many small pieces, indicating that the Al2 O3 –TiC/W18Cr4V interfacial bond strength is high at this time. The fracture morphology of the Al2 O3 –TiC/W18Cr4V diffusion interface at a joining temperature of 1130 °C and a joining time of 45 min is shown in Fig. 3.43d, where the fracture occurs alternately between the Al2 O3 –TiC ceramics and the interface. All of the four fracture types mentioned above, except the first one, which is an interfacial fracture, are mixed fractures, that is, partly fractured at the interface and partly fractured at the ceramic. The joint strength of interfacial fracture is less than that of mixed fracture. For Al2 O3 –TiC/Q235 and Al2 O3 –TiC/18-8 diffusion welded joints, when the heating temperature is low (1100 °C for Al2 O3 –TiC/Q235 joint and 1090 °C for Al2 O3 –TiC/18-8 joint), the joint shear fracture is an interface fracture; when the heating temperature is high (1180 °C for Al2 O3 –TiC/Q235 joint and 1170 °C for Al2 O3 –TiC/18-8 joint), shear fracture is ceramic fracture within the Al2 O3 –TiC ceramic. When the heating temperature is moderate, the Al2 O3 –TiC/steel joint shear fracture shows a mixed fracture. (2) Joint shear fracture morphology The shear fracture of the Al2 O3 –TiC/W18Cr4V diffusion interface was observed using scanning electron microscopy (SEM), and the fracture morphological characteristics are shown in Fig. 3.44. When the shear fracture of the Al2 O3 –TiC/W18Cr4V diffusion interface was observed at low magnification of SEM, the fracture is flush, with no obvious plastic deformation and obvious fracture steps, see Fig. 3.44a. SEM high magnification observation of the fracture shows the presence of a small amount of tearing ridges in the fracture where plastic deformation occurred, see Fig. 3.44b. The shear fracture morphology analysis of the Al2 O3 –TiC/W18Cr4V diffusion interface shows that the Al2 O3 –TiC/W18Cr4V diffusion interface fracture mainly occurs at the interface near the Al2 O3 –TiC composite ceramic side. Since the Ti– Cu–Ti interlayer reacts with the Al2 O3 –TiC ceramic to form brittle compounds such as TiC and Ti3 Al during the diffusion welding process, the fracture tends to occur at the Al2 O3 –TiC/Ti interface reaction layer, extending towards the Al2 O3 –TiC ceramic side. The TiC particle-reinforced phase in Al2 O3 –TiC ceramics prevents crack extension and improves the fracture toughness of the material. The fracture of Al2 O3 –TiC ceramics is mainly transcrystalline fracture, with a small amount of intercrystalline fracture. The fracture morphology of the Al2 O3 –TiC/W18Cr4V diffusion interface,

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(b)

(a)

(a) Fracture step

(b) Tearing morphology

Fig. 3.44 Shear fracture morphology of the Al2 O3 –TiC/W18Cr4V diffusion welding interface

in general, shows the characteristic of brittle cleavage fracture (Fig. 3.44), with obvious cleavage steps. Cleavage fracture is the fracture of crystalline material caused by breaking of atomic bonds along the cleavage plane in the crystal. The cleavage surface is very flat, so the crack within the grain is flat. When a cleavage crack expands from one grain across a grain boundary into an adjacent grain, the direction of crack extension changes, as shown in Fig. 3.45a. A single cleavage crack within a grain can expand on two parallel cleavage surfaces at the same time, as shown in Fig. 3.45b. Two parallel cleavage cracks overlap and form a cleavage step by the action of two cleavage or shear stress, as shown in Fig. 3.45c. When the cleavage crack extension meets a spiral dislocation in the perpendicular direction of the crack plane, the cleavage cracks continue to extend and form a cleavage step along two parallel cleavage planes with one atomic distance apart, as shown in Fig. 3.45d. Analysis of the Al2 O3 –TiC/W18Cr4V diffusion interface shear fracture revealed the presence of some granular inclusions on the shear fracture. These granular inclusions can easily cause stress concentration near the interface. When the stress value exceeds the limit value of interfacial fracture stress, it will cause cracking at the interface. When ceramic fracture occurres at the Al2 O3 –TiC ceramic with steel diffusion welded joint shear test, the fracture morphology shows mirror area—mist area— serrated area. The Al2 O3 –TiC/Q235 diffusion welded joint with heating temperature of 1180 °C, holding time of 45 min and pressure of 15 MPa is broken within the Al2 O3 –TiC ceramic near the interface. The shear strength is 111 MPa, which is a ceramic fracture pattern with typical characteristics of a brittle material fracture. Ceramic fracture begins with a defect and expands after the crack nucleation. The crack expansion is slow at the beginning, t and the energy release rate increases as the crack expansion accelerates. The crack bifurcation arises when a critical rate is reached. The bifurcation process continues to repeat itself, the crack families form. The crack expansion process interacts with the microstructure of the material,

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Cleavage step (b) Extension on two parallel cleavage surfaces

(a) Transcrystalline extension

Spiral dislocation

Cleavage step

Cleavage step

(d) Meeting with spiral dislocation

(c) Cleavage step Fig. 3.45 Cleavage fracture process

the stresses and the resulting elastic waves. Such interactions form unique fracture morphologies on the fracture surface. These morphological features can provide important information about the initial location of the crack, i.e., the source of the crack. (3) Microscopic mechanism of joint fracture Al2 O3 –TiC/W18Cr4V diffusion joint fracture is brittle cleavage fracture. The cleavage fracture process can be divided into two stages: crack initiation and instability expansion, and can only occur when the Griffith energy condition is satisfied. For the brittle fracture problem of brittle materials, Griffith proposed the crack instability expansion condition from the energy viewpoint. He stated that the elastic strain energy released by crack expansion can overcome the work done by the material resistance, then the crack instability expands. The condition for crack expansion in a plane stress state is as follows. σ ≥ σc σc =

2γ E (1/2) πa

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(3.2)

(3.3)

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where γ —the surface energy of the material. a—half-length of the crack. E—the elasticity modulus of the material. σ c —critical stress. When a cleavage fracture of the Al2 O3 –TiC/W18Cr4V joint occurs, it is not completely brittle fracture, but is accompanied by a small amount of ductile fracture. Thus the plastic deformation work γ p at the crack tip must also be overcomed. Orowan corrected the Griffth energy condition to obtain the condition for cleavage crack extension in the planar state as follows (2γp E) πa)(1/2)

(3.4)

(2γ p E) πa1 − υ (1/2)

(3.5)

Planar stress state σc = Planar strain state σc =

where γ p —work of plastic deformation of the material. υ—Poisson’s ratio. On the basis of satisfying the energy conditions, the initiation process of general cleavage fracture can be divided into two steps: the first step is the nucleation of microcracks, and the second step is the extension of microcracks, thus initiating cleavage fracture. During the Al2 O3 –TiC/W18Cr4V diffusion welding process, the stress concentrations tend to form at the contact interface, which makes the diffusion joint interface sproduce a large shrinkage during the cooling phase, and can easily cause microcracks. These microcracks continue to expand under external loads, eventually leading to the fracture of the Al2 O3 –TiC/W18Cr4V diffusion interface. Figure 3.46a shows microcracks on the Al2 O3 –TiC ceramic side at the Al2 O3 – TiC/W18Cr4V diffusion interface. Figure 3.46b shows microcracks in the reaction layer. In addition, there are some impurities at the interface that also tend to cause stress concentration and become a source of microcracks. The formation of a microcrack does not necessarily cause cleavage fracture, but only when the local stress applied to it exceeds a critical stress can the microcrack

(a)

Al2O3-TiC Microcrack

Interfacial transition zone

(b)

Al2O3-TiC

Microcrack

Interfacial transition zone

Fig. 3.46 Microcracks near the Al2 O3 –TiC/W18Cr4V interface

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be extended. Because the cleavage is the breaking of atomic bonds along a certain crystal plane, the tip of the microcrack that causes the cleavage fracture should have the sharp degrees of the order of atomic spacing. If the tip of the microcrack is blunted for some reason it will not initiate cleavage fracture. In the shear test, only when the shear stress causes the Al2 O3 –TiC/steel interface microcrack to expand to a large enough length crack,the cleavage fracture of the Al2 O3 –TiC/steel joint can be caused. (4) Influencing factors of joint fracture. The factors influencing the fracture of Al2 O3 –TiC/W18Cr4V diffusion joints are complex, such as the formation of multiple compounds by interfacial reactions between Al2 O3 –TiC ceramics and W18Cr4V steel, the presence of large stresses at the interface, etc. All these influencing factors are related to the nature of the base material, the joining process parameters and the interlayer material. With a certain base material, the diffusion joining process parameters and the interlayer material are the most important factors affecting the interfacial reaction and the stress distribution in the joint, thus affecting the fracture performance of the joint. ➀ Heating temperature Heating temperature is the most important parameter for Al2 O3 –TiC/W18Cr4V diffusion welding. During thermal activation, temperature has a significant effect on the kinetics of the process. There is an optimal range of heating temperature in the diffusion joining process. If the temperature is too low, the Ti–Cu–Ti intermediate layer cannot be melted sufficiently, and the interfacial reaction between the active element Ti and Al2 O3 –TiC ceramics and W18Cr4V is not sufficient, and the interfacial reaction layer is too narrow, so the good bonding cannot be formed between the interfaces. If the temperature is too high, both the interfacial reaction layer microstructure and the base material microstructure will be coarsened, which will likewise reduce the joint strength. The interfacial transition zone was narrow at temperatures of 1080 °C and 1100 °C when diffusion joining Al2 O3 –TiC andW18Cr4V with Ti–Cu–Ti interlayer. When the heating temperature was 1080 °C and the holding time was 45 min, the interfacial transition zone was only 32 μm, and the interfacial shear fracture strength was low, and the joint was easy to break from the Al2 O3 – TiC/W18Cr4V connection interface and the interfacial fracture occurred. When the temperature was increased to 1130 °C, the interfacial reaction was more adequate, the width of the interfacial transition zone reached 57 μm, the shear strength reached 154 MPa, and the joint fracture alternated between the interface and the Al2 O3 –TiC ceramic. When the temperature was increased to 1160 °C, the width of the interfacial transition zone increased to 72 μm, but the shear strength instead decreased to 141 MPa, which was lower than the shear strength when the heating temperature was 1130 °C. The joints were easy to crack from the interface to the ceramic side and finally broke on the ceramic side. High temperatures also increase the axial tensile stress in the joint, which also reduces the strength of the joint and causes fracture.

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➁ Holding time The holding time mainly determines the degree of elemental diffusion and interfacial reaction, and different holding times result in different interfacial products and interfacial microstructures. The selection of the holding time must take into account the heating temperature. Under a certain heating temperature, the thickness of the reaction layer increases parabolically with the increase of the holding time. At the same time, there is an optimum value for the effect of the holding time on the joint properties. The strength of the reaction phase gradually decreases with the increase of the holding time. While the interfacial strength shows an increasing trend with the increase of holding time at the initial moment, and stops increasing when a certain holding time is exceeded. The joint presents a macroscopic strength that is a combination of both. During Al2 O3 –TiC/W18Cr4V diffusion welding, if the holding time is too short, the interfacial reaction layer is too thin to form a good connection. When the holding time is less than 30 min, the interface is easily subjected to interfacial fracture. With the extension of the holding time, the width of the interfacial transition zone increases. When the holding time is extended to 45 min, the Al2 O3 –TiC/W18Cr4V diffusion bonding interfacial transition zone reaches a more suitable width with the highest interfacial strength. When the heating temperature was 1130 °C and the joining time was 45 min, the shear strength of the Al2 O3 –TiC/W18Cr4V joint reached 154 MPa and the joint fracture was mixed fracture. When the holding time continued to be extended to 60 min, the interfacial transition zone continued to widen, but the strength of the joint started to decrease. Because the holding time is too long, the thickness of the reactive layers such as TiC and Ti3 Al in the interfacial transition zone increases, which in turn increases the interfacial stress and thus reduces the strength of the joint. ➂ Welding pressure Welding pressure plays an important role in the process of Al2 O3 –TiC/W18Cr4V diffusion welding. The welding pressure has four main effects: first, it can promote close contact of the welding surface, increase the contact area, reduce porosity defects, increase the diffusion channels of the elements and improve the microstructure of the joint; second, it can promote the spreading of the Cu–Ti liquid alloy on the surface of Al2 O3 –TiC ceramics, thus promoting the occurrence of interfacial reactions and forming a continuous dense reaction layer; third, it can exclude the excess of the Cu–Ti liquid phase, reducing the maximum width of the liquid phase zone at the welding temperature, thus reducing the time required for isothermal solidification and improving the connection efficiency. The reduction of liquid Cu is also conducive to improving the high-temperature resistance of the joint; fourth, the increase in pressure reduces the amount of Cu–Ti liquid metal in the connection zone, which can reduce the shrinkage of the liquid metal during solidification, reduce the probability of defects during solidification, and prevent cracking due to stress concentration and cracking.

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When the welding pressure is small (less than 10 MPa), the Al2 O3 –TiC/W18Cr4V diffusion interface has cavity defects, the residual stress in the joint is large, and the interface fracture is easily produced, and the joint strength is low. With the increase of welding pressure, the Al2 O3 –TiC/W18Cr4V diffusion joining interface defects gradually disappear and a reaction layer of suitable thickness is formed, the residual stress in the joint decreases, and the joint strength gradually increases. However, it is not better to have a higher welding pressure. Too high pressure will make too much liquid metal extrusion between the joining surfaces, instead, the Al2 O3 – TiC/W18Cr4V interface cannot form a suitable reaction layer, and it is also easy to produce microcracks in Al2 O3 –TiC ceramics. So when the pressure exceeds the optimum value, the joint strength decreases instead by continuing to increase the welding pressure. Al2 O3 –TiC/steel diffusion welded joints have the highest shear strength when type II mixed fracture occurs, whether it is Al2 O3 –TiC/Q235 steel joints or Al2 O3 – TiC/18-8 steel joints, compared to other fracture types.Al2 O3 –TiC/Q235 steel diffusion welded joints have a shear strength of 143 MPa when type II mixed fracture occurs, and Al2 O3 –TiC/18-8 steel diffusion welded joints have a shear strength of 125 MPa when type II mixed fracture occurs. The shear strength of Al2 O3 –TiC/18-8 steel diffusion welded joints is lower than that of Al2 O3 –TiC/Q235 steel diffusion welded joints regardless of any type of fracture. That occurs because of the large residual stresses in the joint during diffusion welding of Al2 O3 –TiC ceramics to 18-8 steel and the brittleness of the interfacial reaction products. Analysis of the shear fracture morphology of the Al2 O3 –TiC/steel diffusion welded joint shows that the weak areas of the Al2 O3 –TiC/steel diffusion welded joint are: first, the Al2 O3 –TiC ceramic near the interface; and second, the Cu–Ti and Fe–Ti intermetallic compound layers in the reaction zone of the intermediate layer immediately adjacent to the Al2 O3 –TiC ceramic.

Bibliography 1. Wanqun H, Yajiang L, Juan W, et al (2010) Element distribution and phase constitution of Al2 O3 –TiC/W18Cr4V vacuum diffusion bonded joint. Vacuum 85(2):327–331 2. Wanqun H, Yajiang L, Juan W, et al (2010) Microstructure of Al2 O3–TiC composite ceramics and Q235 steel diffusion bonded interface. Trans China Weld Inst 31(8):101–104 3. Barrena MI, Matesanz L, Gómez de Salazar JM (2009) Al2 O3 /Ti6Al4V diffusion bonding joints using Ag–Cu interlayer. Mater Charact 60(11):1263–1267 4. Song SX, Ai X, Zhao J, et al (2003) Mechanical properties and toughening and strengthening mechanism of Al2 O3 /TiC nanocomposite tool materials. Mech Eng Mater 27(12):35–37, 41 5. Shen X, Li Y, Putchkov UA, et al (2009) Finite-element analysis of residual stresses in Al2 O3 – TiC/W18Cr4V diffusion bonded joints. Comput Mater Sci 45(2):407–410.

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6. Xiaoqin S, Yajiang L, Juan W, et al (2009) Diffusion bonding of Al2 O3 –TiC composite ceramic and W18Cr4V high speed steel in vacuum. Vacuum 84(3):378–381 7. Xiaoqin S, Yajiang L, Puehkov UA, et al (2008) Stress distribution in Al2 O3 -TiC/1Crl8 Ni9Ti diffusion bonded joint. Trans China Weld Inst 29(10):41–44. 8. Juan W, Yajiang L, Haijun M, et al (2006) Microstructure in diffusion bonded TiC–Al2 O3 W18 Cr4 V joint with Ti/ Cu/ Ti composite interlayer. Trans China Weld Inst 27(9):9–12 9. Xiaoqin S, Yajiang L, Juan W, et al (2008) Numerical simulation of stress distribution in Al2 O3 – TiC/Q235 diffusion bonded joints. China Weld 17(4):47–51

Chapter 4

Joining of Nickel–Aluminium and Titanium–Aluminium Intermetallic Compounds

An intermetallic compound is made up of two or more metallic elements in proportion, with a long-range ordered crystal structure and basic metallic properties different from those of their constituent elements. The metallic elements are bonded to each other by a mixture of covalent and metallic bonds coexisting, and the properties are between ceramics and metals. Since the 1980s, toughening research of Ni3 Al, toughness improvement of Ti3 Al and TiAl-based alloy and performance improvement of Fe3 Al, have made significant progress on the research and development of intermetallic compounds high temperature structural materials and applications. Also, welding of Ni–Al and Ti–Al intermetallic compounds is increasingly drawing the attention of many researchers.

4.1 Development and Properties of Intermetallic Compounds 4.1.1 Development of Intermetallic Compounds for Structures Intermetallic compounds have a long-range ordered superdot structure and maintain the coexistence of metallic and covalent bonds between atoms, enabling them to combine both the plasticity of metals and high-temperature strength of ceramics. Intermetallic compounds containing Al and Si elements also have good oxidation resistance and low density. Composition of the intermetallic compounds can deviate from the stoichiometry within a certain range and still maintain their structural stability, which is expressed as ordered solid solution on the alloy state diagram. The long-range ordered superdot structure of intermetallic compounds maintains strong metal bonding and covalent bonding, which make them have special physical © Chemical Industry Press 2023 Y. Li, Joining Technology and Application of Advanced Materials, Advanced and Intelligent Manufacturing in China, https://doi.org/10.1007/978-981-19-9689-4_4

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and chemical properties and mechanical properties, such as special electrical properties, magnetic properties and high-temperature properties, etc. They are a promising new type of high temperature structural materials. Research of intermetallic compounds began in the 1930s, and the intermetallic compounds used for structural materials are mainly concentrated in the three major alloy systems Ni–Al, Ti–Al and Fe–Al. Ni–Al and Ti–Al intermetallic compounds have excellent high temperature performance, but are expensive and mainly used in aerospace and other fields. Compared with Ni–Al and Ti–Al intermetallic compounds, Fe–Al intermetallic compounds have the advantages of high strength and corrosion resistance, as well as low cost and low density, and have broad application prospects. When heated, steel gradually turn red and soften (until they melt into liquid). High temperatures are the enemy of most metals, and metals lose their strength at high temperatures and become “unbearable”. This does not happen with intermetallic compounds. At temperatures above 700 °C, some intermetallic compounds become harder and even stronger. The “heroic” nature of intermetallic compounds can be seen under high temperatures. Intermetallic compounds has this special property, which is related to their internal atomic structure. By intermetallic compound, we mean a compound generated by the combination of metal and metal, and metal-like material and metal in the form of covalent bonds, with a highly ordered pattern of atomic arrangement. When it is present in the organization of the alloy in the form of tiny particles, it increases the overall strength of the alloy. In particular, the strength of the alloy increases with temperature in a certain temperature range, which gives the Ni–Al and Ti–Al intermetallic compounds great potential advantages for high-temperature structural applications. However, intermetallics has greater room-temperature brittleness, along with the high-temperature strength. When intermetallics were first discovered in the 1930s, they had mostly zero room-temperature ductility, meaning that they would break upon folding. So, many people predicted that intermetallic compounds would have no practical value as a bulk material. In the mid-1980s, American scientists made a breakthrough in the study of roomtemperature brittleness of intermetallic compounds. They added a small amount of boron (B) to a Ni–Al intermetallic compound, which increased its room-temperature elongation to 50%, comparable to the ductility of pure aluminum. This important discovery and the prospects it holds have attracted material scientists from all over the world to conduct in-depth research on intermetallic compounds, which are beginning to take on a new look in the field of new materials. In the past 20 years, people have started to pay attention to the development and application of intermetallic compounds, which is an important shift in the field of materials and one of the important directions for future materials development. Intermetallic compounds have properties that other solid solution materials do not have due to their special crystal structure. In particular, solid solution materials usually decrease in strength as the temperature raised, but the strength of certain intermetallic compounds increases in a certain range instead as the temperature raised, and this

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is the basis for its potential as a new type of high-temperature structural material. In addition, intermetallic compounds have some properties that are several times or even 20 or 30 times those of solid solution materials. At present, in addition to being high-temperature structural materials, other functions of intermetallic compounds have been developed one after another. Rare earth compound permanent magnetic materials, hydrogen storage materials, super magnetostrictive materials, and functional sensitive materials have been introduced one after another. The application of intermetallic compounds has greatly contributed to the progress and development of high technology, promoting the miniaturization, lightweight, integration and intelligence of structures and components, leading to the continuous emergence of a new generation of components. The greatest use of intermetallic compounds, “high-temperature materials”, is in the aerospace field, such as low density, high melting point, high-temperature performance of Ti–Al intermetallic compounds have extremely attractive application prospects.

4.1.2 Basic Properties of Intermetallic Compounds Intermetallic compounds are compound phases formed between metals and metals or metalloids with long-range ordered superdotted crystal structures, strong atomic bonding, high modulus of elasticity at high temperature, and good oxidation resistance, thus forming a series of new structural materials, such as aluminide materials of titanium, nickel, and iron, which have promising applications. Intermetallic compounds do not follow the traditional laws of valency and have the properties of metals. But the crystal structure differs from that of the two elements that make it up in that the atoms of each occupy certain dotted positions in an ordered arrangement. Typical long-range ordered structures are formed mainly on the basis of the three main crystal structures of metals: face-centered cubic, body-centered cubic and dense-row hexagonal. For example, Ni3 Al is a face-centered cubic ordered superdot structure, Ti3 Al is a dense-rowed hexagonal ordered superdot structure, and Fe3 Al is a body-centered cubic ordered superdot structure. Many intermetallic compounds can maintain structural stability within a certain range and appear as ordered solid solutions on the phase diagram. The main factors that determine the phase structure of intermetallic compounds are electronegativity, size factor and electron concentration. The crystal structure of intermetallic compounds, although more complex or ordered, still has metallic properties in terms of atomic bonding, metallic luster, electrical and thermal conductivity, etc. However, their electron cloud distribution is not completely uniform, and there is a certain degree of directionality, with some degree of covalent bonding characteristics, leading to higher melting points and directional interatomic bonding. Intermetallic compounds can be divided into two categories, structural and functional. The former are materials used as load-bearing structures with good room temperature and high temperature mechanical properties, such as high temperature

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ordered intermetallic compounds Ni3 Al, NiAl, Fe3 Al, FeAl, Ti3 Al, TiAl, etc. The latter have some special physical or chemical properties, such as magnetic material YCo5 , shape memory alloy NiTi, superconducting material Nb3 Sn, hydrogen storage material Mg2 Ni, etc. Compared with disordered alloys, the long-range ordered superdot structure of intermetallic compounds maintains very strong metal bonding and has many special physical and chemical properties, such as electrical properties, magnetic properties and high temperature mechanical properties. The intermetallic compounds containing Al and Si also have high resistance to oxidation and corrosion. Intermetallic compounds composed of light metals have low density and high specific strength, which are suitable for the application requirements of the aerospace industry. Research and development of intermetallic compounds for applications have been of great importance. Among A3B-type intermetallic compounds, Ti3 Al, Ni3 Al and Fe3 Al-based alloys have become increasingly mature. The brittleness problem has been solved and is entering the industrial application stage. Among the AB intermetallic compounds, the room-temperature brittleness of TiAl-based alloys has been enhanced, cast TiAl alloys are initially entering industrial applications, and deformed TiAl alloys are under intensive research and development. As the room-temperature brittleness of NiAl alloys is still to be solved and the strength above 500 °C is low, a lot of research work is still needed for their engineering applications. Research of FeAl alloys is getting more and more advanced and their industrial applications are being explored.

4.1.3 Three Promising Intermetallic Compounds Intermetallic compounds based on aluminides are new high temperature structural materials with promising applications. Recent years, the intermetallic compounds that have been studied and made significant progress at home and abroad are mainly Ti–Al, Ni–Al and Fe–Al three systems of A3 B and AB type intermetallic compounds, of which A3 B type intermetallic compounds are mainly Ti3 Al, Ni3 Al and Fe3 Al; AB type intermetallic compounds are mainly TiAl, NiAl and FeAl. Especially Ni–Al and Ti–Al system intermetallic compounds, due to higher high-temperature strength, excellent oxidation resistance and corrosion resistance than nickel-based alloys, lower density and higher melting point, can work in higher temperature and harsh environment, and have broad application prospects in aerospace, energy and other high-tech fields. The physical properties of several important intermetallic compounds are shown in Table 4.1. Ni–Al and Ti–Al intermetallic compounds are suitable for aerospace materials and have good application potential, which have been generally valued by Europe, the United States and other developed countries. Some Ni–Al alloys have been applied or tried, such as for diesel engine components, electric heating components, aerospace

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Table 4.1 Physical properties of several important intermetallic compounds Intermetallic compound

Structure Density/g × Melting cm−3 point/°C

Young’s Coefficient Order critical modulus/GPa of linear temperature/°C expansion/× 10–6 °C−1

Ni–Al alloy

Ni3 Al

Ll2

7.40

1397

178

16.0

1390

NiAl

B2

5.90

1638

293

14.0

1640

Ti–Al alloy

Ti3 Al

DO19

4.50

1680

110–145

12.0

1100

TiAl

Ll0

3.80

1480

176

11.0

1460

Fe–Al alloy

Fe3 Al

DO3

6.72

1540

140

–

540

FeAl

B2

5.56

1250–1400 259

–

1250–1400

fasteners, etc. Ti–Al alloys can replace nickel-based alloys made of aerospace engine high-pressure turbine stator support ring, high-pressure compressor box, engine combustion chamber expansion nozzle, etc.. China’s aerospace industry is trying out these alloys to manufacture engine hot-end components, with great prospects. As structural materials, the most promising applications are Ni–Al, Ti–Al, Fe–Al system intermetallic compounds, such as Ni3 Al, NiAl, Ti3 Al, TiAl, Fe3 Al, FeAl, etc. Researchers around the world have conducted weldability studies on Ni–Al, Ti–Al and Fe–Al intermetallic compounds and have made promising progress. Fe3 Al intermetallic compound can replace stainless steel, heat-resistant steel or high-temperature alloy in many occasions due to its high oxidation resistance and wear resistance for the manufacture of corrosion-resistant parts, heat-resistant parts and wear-resistant parts, and its good resistance to sulfide is suitable for applications under severe conditions (such as high-temperature corrosive environments). For example, structural parts of thermal power plants, structural parts for carburizing furnace atmosphere work, chemical devices, automobile exhaust gas exhaust, petrochemical catalytic cracking device, guide of heating furnace, high temperature grates, etc. In addition, Fe3 Al intermetallic compound has excellent high-temperature oxidation resistance and high resistivity, which may be developed into a new type of electric heating material. Fe3 Al can also be made into a composite structure with WC, TiC, TiB, ZrB and other ceramic materials, which has a broader application prospect. (1) Ni–Al system intermetallic compounds The Ni–Al family of intermetallic compounds mainly includes Ni3 Al and NiAl. Ni3 Al has a melting point of 1395 °C and has a face-centered cubic ordered L12 superdot structure below the melting point. The phase diagram of Ni–Al binary alloy is shown in Fig. 4.1. In the Ni–Al binary system, in addition to the solid solution of Ni and Al, five stable binary compounds exist, namely Ni3 Al, NiAl, Ni5 Al3 , Al3 Ni2 , and Al3 Ni, where Ni3 Al, Al3 Ni2 , and Al3 Ni are formed by peritectic reactions, Ni5 Al3 by peritectoid reaction, and NiAl by homogeneous crystal transformation. Except for the NiAl single-phase region where exists a wide composition range of 45 to 60% Ni (molar fraction), the other

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Fig. 4.1 Phase diagram of Ni–Al binary alloy

compounds have a narrow composition range, for example, the low-temperature Ni3 Al phase has a composition range of 73 to 75% Ni (molar fraction). It was shown that among the Ni–Al family of alloys, only Ni3 Al and NiAlbased alloys have potential for application as structural materials. The other three compounds have difficulty competing with high-temperature alloys because of their very low melting points. ➀ Ni3 Al intermetallic compounds Ni3 Al is an intermetallic compound in the range of 4.5% solid solution on both sides of a fixed proportion of chemical composition. Ni3 Al has a melting point of 1397 °C, a lattice constant of 0.3565–0.3580 nm, a density of 7.4 g/cm3 , and a face-centered cubic ordered LI2 type structure below the melting point. Ni3 Al has unique high-temperature properties, with its yield strength increasing with temperature up to 800 °C, but is very brittle at room temperature with a significant tendency of intergranular fracture. Tests have shown that the room temperature plasticity of Ni3 Al can be improved by microalloying. The improving effect of trace element B on the room-temperature plasticity of polycrystalline Ni3 Al is closely

4.1 Development and Properties of Intermetallic Compounds 500

Yield strength σs / MPa

Elongation δ/%

60 50 40 30 20 10 0 0.00

177

0.04

0.08

0.12

0.16

0.20

400 300 200 100 0 0.00

0.04

0.08

0.12

0.16

0.20

B /%

B /% (a)Elongation

(b) Yield strength

Fig. 4.2 Effect of boron on the elongation and yield strength of Ni3 Al

related to the Al content. Only when the Al content is less than 25% of the molar fraction, trace element B can effectively improve the room temperature plasticity of Ni3 Al and suppress the tendency of intergranular fracture. The effect of boron (B) content on the elongation (δ) and yield strength (σ s ) of Ni3 Al is shown in Fig. 4.2. The room temperature elongation increases from 0 to 40– 50% with the addition of 0.02 to 0.05% of element B to Ni3 Al. However, when the molar fraction of Al in the Ni3 Al matrix is higher than 25%, the plasticity decreases sharply with increasing Al content and shifts the fracture mode from transgranular fracture to intergranular fracture. The addition of Fe and Mn to the Ni3 Al matrix can also improve the room temperature plasticity of the alloy by displacing Ni and Al and changing the interatomic bonding state and charge distribution. For example, the addition of Fe at a mass fraction of 15% (or Mn at 9%) is more effective, and its fracture elongation after room temperature can reach 8% and 15%, respectively. However, specific strength decreases after macro-alloying. In addition, the room temperature and high temperature strength of Ni3 Al can be further improved by solid solution strengthening, but usually only those solid solution elements that displace the Al sub-dot positions produce the strengthening effect. The addition of the alloying element hafnium (Hf) can also significantly increase the strength of Ni3 Al, especially at high temperature. The chemical composition of five Ni3 Al alloys from the United States is shown in Table 4.2. These materials have applications, such as IC-396 for diesel engine parts and IC-50 has been used in electrical heating elements and fasteners for aerospace. ➀ NiAl intermetallic compounds NiAl intermetallic compounds have a high melting point (1600 °C), a density of 5.9 g/cm3 , a body-centered cubic ordered B2 superdot structure, and high oxidation resistance, and are a promising high-temperature intermetallic compound for applications. The main problems affecting the practical use of NiAl intermetallic

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Table 4.2 Chemical compositions of five Ni3 Al alloys from the United States Serial number

Material name

Chemical composition (mole fraction)

1

IC-50

Ni –Al23% ± 0.5%–Hf(Zr) ± 0.3%–B0.1% ± 0.05%

2

IC-218

Ni–Al16.7% ± 0.3%–Cr8%–Zr0.5% ± 0.3%–B0.1% ± 0.05%

3

IC-328

Ni–Al17.0% ± 0.3%–Cr8%–Zr0.2% ± 0.1%–Ti0.3% ± 0.1%–B0.1% ± 0.05%

4

IC-396

Ni–Al16.1% ± 0.3%–Cr8%–Zr0.25% ± 0.15%–Mo1.7% ± 0.3%–B0.1% ± 0.07%

5

IC-405

Ni–Al18% ± 0.5%–Cr8%–Zr0.2% ± 0.1%–Fe12.2% ± 0.5 5–B0.1% ± 0.05%

compounds are few independent slip systems at room temperature, low plasticity, high brittleness, and low strength above 500 °C. Since NiAl intermetallic compounds can remain stable over a wide range of compositions, it is possible to improve their mechanical properties through alloying. For example, the addition of Fe to NiAl can increase strength and improve elongation by forming a two-phase organization (Ni, Fe) (Fe, Ni) and (Ni, Fe)3 (Fe, Ni), and the addition of Ta or Nb improves creep strength by precipitating second-phase particle strengthening. In addition, creep strength and high-temperature strength can be improved by mechanically alloying with the addition of Al2 O3 , Y2 O3 and ThO2 dispersive particles, but room-temperature strength is reduced. Plasticity can also be improved by grain refinement, but the critical grain size required to significantly improve room-temperature plasticity is small (less than 3 μm in diameter). And although fine grain organization can be obtained by new processes such as rapid solidification and powder metallurgy, it affects its creep resistance. (2) Ti–Al system intermetallic compounds There are two intermetallic compounds (Ti3 Al, TiAl) in the Ti–Al system that have received attention for their development. The Ti3 Al intermetallic compound-based alloy is called Ti3 Al-based alloy and the TiAl intermetallic compound-based alloy is called γ-TiAl-based alloy (referred to as TiAl alloy).The binary phase diagram of the Ti–Al system is shown in Fig. 4.3. In the early 1950s, American scholars studied the properties of Ti–50Al alloy, which was abandoned because the plasticity of the alloy was too poor. 15 years later, Professor M. Blackburn of the United States studied about 100 TiAl alloys of different compositions and found the alloy with the best properties, Ti–48Al–1V–0.3C, the first generation TiAl alloy. The room temperature plasticity could reach 2%, but TiAlbased alloys were not developed as engineering alloys. It was not until the late 1980s that the second generation TiAl alloy (Ti–48Al–2Cr–2Nb) was developed by GE in the United States and proved its good comprehensive performance that people’s interest in TiAl alloy was aroused. After a lot of research, the third generation TiAl alloy has been developed. The comparison of Ti3 Al and TiAl alloys with Ti-based alloys and Ni-based alloys is shown in Table 4.3. As can be seen from the table, Ti3 Al and TiAl-based alloys have a density similar to that of Ti-based alloys; excellent high-temperature

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179

Fig. 4.3 Binary phase diagram of Ti–Al system

performance similar to that of Ni-based alloys, but the density is only half that of Ni-based high-temperature alloys. This is a highly promising alternative to Ni-based alloys for high-temperature structural materials that can be used in high temperature parts of aviation engines (such as turbine discs, blades and valve valves, etc.). The main application advantages of TiAl-based alloysn. are list as follow. (i)

TiAl-based alloys have about 50% higher specific stiffness than other commonly used structural materials for aero-engines, which is beneficial for components requiring low clearance, such as cases, members and supports, and can extend the life of components such as blades by shifting noise vibrations to higher frequencies. (ii) The good creep resistance of TiAl-based alloys at 600–700 °C, making it possible to replace some Ni-based high-temperature alloy components (half the weight). (iii) Good flame retardancy, replacing some expensive Ti alloys of flame retardant design.

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Table 4.3 Comparison of the properties of Ti3 Al and TiAl alloys with Ti-based alloys and Ni-based alloys Ti-based alloy

Ti3 Al-based alloy

TiAl-based alloy

Ni-based alloy

4.5

4.1–4.7

3.7–3.9

7.9–9.5

95–115

100–145

160–180

206

Yield strength/MPa 380–1150

700–990

350–600

800–1200

Tensile strength/MPa

480–1200

800–1140

440–700

1250–1450

Creep limit/°C

600

750

750➀ –950➁

800–1090

Oxidation limit/°C

600

650

800➂ –950➃

870–1090

Linear expansion coefficient/× 10–6 °C−1

9.1

12.0

11.0

13.3

Room temperature plasticity/%

10–25

2––10

1–4

3–25

High temperature plasticity/%

12–50

10–40

10–20

20–80

Room temperature fracture toughness/MPa × m1/2

12–80

13–35

10–30

30–100

Crystal structure

hcp/bcc

DO19

LI0

Fcc/LI2

Performance Density/g ×

cm−3

Elastic modulus/GPa

➀ Bimorphic organization ➁ Full lamellar sheet organization ➂ No coating ➃ Coating/controlled cooling

The disadvantage of TiAl alloy components is their lower resistance to damage, and their lower room temperature plasticity, fracture toughness and high temperature crack expansion increases the possibility of failure. Ti3 Al belongs to a close-packed hexagonal ordered DO19 super dotted structure, with a small density (4.1 to 4.7 g/cm3 ) and a high elastic modulus (100 to 145 GPa). The mass can be reduced by 40% compared to nickel-based high-temperature alloys, and it has good high-temperature properties at high temperature (800 to 850 °C), but room-temperature plasticity is very low and processing and forming is difficult. The solution to these problems is to add β-phase stabilizing elements, such as Nb, V, Mo, etc. for alloying, with Nb playing the most significant role. Mainly by lowering the martensite transformation point (M s ), refining the α2 phase and reducing the slip length, in addition to inducing the formation of a two-phase organization with better plasticity and strength α2 + β. TiAl has a face-centered tetragonal ordered L10 superdot structure. In addition to its very good high temperature strength and creep resistance, TiAl also has low density (3.7–3.9 g/cm3 ), high elastic modulus (160–180 GPa) and good oxidation

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181

resistance, making it an attractive high temperature structural material for aviation and aerospace applications. The composition of Ti3 Al-based alloy and Ti2 AlNb-based alloy developed in China is shown in Table 4.4. Among them, the parts made of TAC–1B alloy successfully participated in the flight of the Shenzhou spacecraft, and a variety of important structural parts of aerospace engines developed by China have also completed flight tests. The turbine guide plate of aero-engine made of TD2 alloy has also been tested in engine test. The mechanical properties and high-temperature endurance of some typical Ti3 Al alloys are shown in Table 4.5. The aerospace industry in China is experimenting with these alloys as partial replacements for nickel-based high-temperature alloys in the manufacture of engine hot-end components. The room temperature plasticity of TiAl can be improved by alloying and controlling the microstructure. Alloys containing a biphasic (α2 + γ) lamellar organization have Table 4.4 Composition of Ti3 Al-based alloys and Ti2 AlNb-based alloys developed in China Brand number Alloy categories

Alloy composition (molar fraction)/%

Phase composition

24-11 25-11 8-2-2

Ti3 Al-based alloy (belonging to the first category)

Ti–24Al–11Nb Ti–25Al–11Nb Ti–25Al–8Nb–2Mo–2Ta

α2 and B2/β two-phase structure

TAC-1 TAC-1B TD2 TD3

Ti3 Al-based alloy (belonging to the second category)

Ti–24Al–14Nb–3V–(0–0.5)Mo Ti–23Al–17Nb Ti–24.5Al–10Nb–3V–1Mo Ti–24Al–15Nb–1.5Mo

Solid solution α2 + B2 two-phase structure or steady state α2 + B2 + O three-phase structure

TAC-3A TAC-3B TAC-3C TAC-3D

Ti2 AlNb-based alloy Ti–22Al–25Nb (belonging to the third Ti–22Al–27Nb category) Ti–22Al–24Nb–3Ta Ti–22Al–20Ni–7Ta

O-phase alloy (ortho-phase) contains a small amount of B2/β phase

Table 4.5 Mechanical properties and high temperature endurance life of typical Ti3 Al alloys Alloy

Yield strength/MPa

Tensile strength/MPa

Elongation/%

High temperature endurance life➀ /h

Ti–24Al–11Nb

761

967

4.8

–

Ti–24Al–14Nb

790–831

977

2.1–3.3

59.5–60

Ti–25Al–10Nb–3V–1Mo

825

1042

2.2

–

Ti–24.5Al–17Nb

952

1010

5.8

>360

Ti–24.5Al–17Nb–1Mo

980

1133

3.4

476

➀ 650 °C, 380 MPa

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better plasticity and strength than those with a single phase (γ) organization. Tests on the alloying elements V, Cr, Mn, Nb, Ta, W, and Mo showed that plasticity could be improved (elongation ≥ 3%) when 1% to 3% V, Mn, or Cr was added to the Ti– Al48 alloy. Increasing the purity of the alloy also helps to improve its plasticity, for example, when the oxygen content is reduced from 0.08% to 0.03%, the elongation in tension of Ti–Al48 alloy increases from 1.9% to 2.7%. Alloying is the basic way to plasticize and toughen Ti3 Al alloys. Adding Nb can improve the strength, plasticity and toughness of Ti3 Al alloy. V can also make the plasticity of the alloy improved, but it is not good for the strength and oxidation resistance of the alloy. Increasing the content of Al, Mo and Ta is good for improving the high temperature strength and creep resistance of the alloy, etc. (3) Fe–Al intermetallic compounds It mainly includes Fe3 Al and FeAl. Fe3 Al has a DO3 type ordered superdotted structure with a higher modulus of elasticity, higher melting point and lower density. It is ferromagnetic at room temperature, and the saturation magnetization intensity of the ordered DO3 superdotted structure is 10% lower than that of the disordered α-phase. The Fe–Al binary alloy shows excellent resistance to high temperature oxidation due to the ability of Fe3 Al to form a dense protective film of alumina at very low oxygen partial pressures.The Fe–Al binary alloy phase diagram is shown in Fig. 4.4.

Fig. 4.4 Phase diagram of Fe–Al binary alloy

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183

The Al-stabilized α-Fe phase with Al atomic percentage content below 18% to 20% is a disordered α-Fe(Al) solid solution phase at room temperature and high temperature. At 25% to 35% Al atomic percentage content, the Fe–Al intermetallic compound has a DO3 type ordered structure with a dotting constant of 0.578 nm, and gradually transforms to a partially ordered B2 structure and a disordered α-Fe(Al) structure with changes in temperature and Al content. The ordering temperature for the transition from DO3 to B2-type structure is about 550 °C; the transition temperature for B2 and α-Fe(Al) structure is about 750 °C. The stable FeAl alloy at room temperature has B2-type ordered structure when the percentage of Al atoms is 36.5–50%, and the dot-site constant is 0.289–0.291 nm depending on the Al content and heat treatment process. In the Fe–Al binary alloy state diagram, the three brittle intermetallic compounds, FeAl2 (49.2% to 50% mass fraction of Al), Fe2 Al5 (54.9% to 56.2% mass fraction of Al), and FeAl3 (59.2% to 59.6% mass fraction of Al), have a narrow composition range, while Fe3 Al and the nearby α-Fe(Al) solid solution have a wide compositional range, which is conducive to the stability of the properties of Fe3 Al-based alloys. The composition and high temperature mechanical properties of several typical Fe3 Al-based alloys are shown in Table 4.6.

4.1.4 Superplasticity of Ni–Al and Ti–Al Intermetallic Compounds Ni–Al and Ti–Al intermetallic compounds are a class of high-temperature structural material with promising application, including Ni3 Al, NiAl, Ti3 Al, TiAl. Since these materials have the characteristics of ceramic (covalent bonding) and metal (metallic bonding), they become a bridge between metal and inorganic nonmetal (ceramic). Superplasticity of intermetallic compounds is a coordinated process consisting of grain boundary sliding mechanism and accompanying dynamic recrystallization and dislocation slippage. The effect of plasticity for fine grain is similar to that of general alloys. While for large grain intermetallic compounds superplasticity is somewhat universal, and its superplasticity is a continuous dynamic reversion and recrystallization process. Subcrystals do not exist in the original large grains before superplastic deformation, and during deformation dislocations form unstable subgrain boundaries by slipping or climbing, and these subgrain boundaries are formed within the original boundaries by absorbing the slipped dislocations within the grain boundaries, thus in situ recrystallization occurs. The continuation of this process leads to a macroscopic superplastic behavior of the material. (1) Superplasticity of Ni–Al intermetallic compounds ➀ Superplasticity of Ni3 Al intermetallic compounds Single-crystal Ni3 Al has good toughness, but polycrystalline Ni3 Al has poor toughness and exhibits intergranular fracture. In the experiment, it was found that the use

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➀ Test temperature 593 °C

208 202

Fe–28Al–2Mo–0.1Zr

Fe–28Al–5Cr–0.5Mo–0.1Zr-0.05B

FA-91

FA-130

61

55

34 49

2 13

Fe–28Al

Fe–28Al–5Cr–0.1Zr–-0.05B

Elongation δ/%

Time/h

207 MPa Persistent strength➀

FA-61

Composition at./%

FA-122

Alloy

Table 4.6 Composition and properties of typical Fe3 Al-based alloys

554

698

480

393

Yield Strengthσ0.2 /MPa

12.6

5.7

16.4

4.3

Elongation δ/%

Room temperature tensile property

527

567

474

345

Yield Strength σ0.2 /MPa

31.2

20.9

31.9

33.4

Elongation δ/%

Tensile property at 600°C

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185

of boron (B) alloying can effectively prevent Ni3 Al from intergranular fracture and greatly improve the plasticity. The IC-218 intermetallic compound Ni–8.5Al–7.8Cr–0.8Zr–0.02B (mass fraction, %) obtained by powder metallurgy, has a grain diameter of 6 μm when the ordered γ, phase contains a disordered γ phase with a volume fraction of 10% to 15%, shows superplasticity at 950–1100 °C and a deformation rate of 10–5 to 10–2 /s. A elongation after fracture of 640% was obtained at 1100 °C and a deformation rate of 8.94 × 10–4 /s, and the deformation mechanism was grain boundary slip. A large number of voids were found in the superplastic deformation zone and it was intergranular fracture. Nanoscale (grain diameter of 50 nm) Ni3 Al intermetallic compounds (IC-218) also exhibit superplasticity at 650–750 °C, showing superplasticity at 650 °C and 725 °C and a deformation rate of 10–3 /s, with elongation after fracture of 380% and 750%, respectively. Ni3 Al intermetallics with slightly larger grains (grain diameter of 10–30 μm) also exhibit superplasticity. ➁ Superplasticity of NiAl intermetallic compounds Although NiAl intermetallic compounds have many excellent properties, severe room temperature brittleness hinders its application. The plasticity and toughness of NiAl intermetallic compounds can be improved by adding of a large amount of Fe elements to the NiAl intermetallic compound to introduce the plastic γ phase. For example, the NiAl–20Fe–YCe alloy in the cast extruded state with a mass fraction (%) of Ni–28.5Al–20.4Fe–0.003Y–0.003Ce shows superplasticity at 850–980 °C and a deformation rate of 1.04 × 10–4 to 10–2 /s. The Ni–50Al (molar fraction, %) intermetallic compound (grain diameter of 200 μm) at 900–1100 °C and deformation rate of 1.67 × 10–4 –10–2 /s, the elongation after fracture can reach 210%. The NiAl–25Cr (molar fraction, %) intermetallic compound (grain diameter of 3–5 μm) at 850–950 °C and deformation rate of 2.2 × 10–4 –3.3 × 10–2 /s, the elongation after fracture can reach 480%, showing superplasticity. Eutectic alloys of NiAl–9Mo type also show superplasticity at 1050–1100 °C and deformation rates of 5.55 × 10–5 to 1.11 × 10–4 /s. (2) Superplasticity of Ti–Al intermetallic compounds (1) Superplasticity of Ti3 Al intermetallic compounds Ti3 Al intermetallic compound is α2 + β organization, Ti–24Al–11Nb alloy (mass fraction, %) can obtain superplasticity of 810% elongation after fracture at 980 °C. Ti– 25Al–10Nb–3V–1Mo alloy (mass fraction, %) can obtain superplasticity of 570% elongation after fracture at 980 °C. Ti–24Al–14Nb–3V–0.5Mo alloy (mass fraction, %) has good low-temperature plasticity and high-temperature strength, and a superplasticity of 818% elongation after fracture can be obtained at 980 °C and a deformation rate of 3.5 × 10–4 /s.

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Fig. 4.5 Effect of test temperature on superplasticity of TiAl intermetallic compounds

(2) Superplasticity of TiAl intermetallic compounds (i) Effect of test temperature on the superplasticity of TiAl intermetallic compound For Ti–47.3Al–1.9Nb–1.6Cr–0.5Si–0.4Mn alloy (mass fraction, %) with a grain diameter of 20 μm and tissue γ + α2 at a strain rate of 8.0 × 10-5 /s, the effect of the test temperature on the superplasticity of the coarse-grained TiAl intermetallic compound is shown in Fig. 4.5a. It can be seen that although the fracture strength increases with increasing test temperature, the fracture stress is lower and plasticity increases, the true stress-deformation curve also changes from a softening type (stress decreases with increasing deformation) to a hardening type (stress also increases with increasing deformation). (ii) Effect of grain size on the superplasticity of TiAl intermetallic compounds The superplasticity of Ti–Al intermetallic compounds is largely influenced by the grain size. For Ti–48Al–2Nb–2Cr alloy (mass fraction, %) with a grain diameter of 0.3 μm and a tissue γ + α2 at a strain rate of 8.3 × 10-4 /s−1 , the effect of test temperature on the superplasticity of fine-grained TiAl intermetallic compounds is shown in Fig. 4.5b. It can be seen that, in the true stress-deformation curve, hardening-type temperature decreases after grain refinement and the degree of hardening intensifies as the test temperature increases. (iii) Effect of alloying elements V, Cr, and Mn elements can increase the plasticity of Ti–Al intermetallic compounds, while interstitial elements O, C, N, and B can decrease the plasticity of Ti–Al intermetallic compounds.

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4.2 Welding of Ni–Al Intermetallic Compounds IC is the abbreviation of “intermetallic compounds”, Ni3 Al-based alloy is called IC alloy in the United States. The main problem of welding Ni–Al intermetallic compoundis welding cracking. The alloy contains 10% Fe and 5% Hf can improve the weld cracking tendency. Adjusting the content of grain boundary element B in Ni3 Al-based alloys is also beneficial in eliminating weld thermal cracking in the alloy.

4.2.1 Diffusion Bonding of NiAl Alloys NiAl alloys have poor room temperature plasticity and toughness, and tend to form a continuous Al2 O3 film on the surface using melt welding method, making its weldability very poor. So NiAl alloys are often diffusion brazed or transition liquid phase diffusion bonded. (1) Diffusion brazing of NiAl and Ni In many cases NiAl is used in Ni-based alloys as the main structure, and the joining of NiAl with Ni-based alloys can be achieved by diffusion brazing. For diffusion brazing of Ni–48Al alloy with industrially pure Ni (mass fraction of Ni 99.5%), an amorphous braze BNi-3 with a thickness of 51 μm can be used as an interlayer. The composition of BNi-3 braze is Ni–Si4.5%–B3.2% (molar fraction), and the solidphase line temperature of the brazing material is 984 °C, and the liquid-phase line temperature is 1054 °C. The diffusion brazing temperature was 1065 °C. When the heating temperature reaches 1065 °C, the brazing material melts to form a transition liquid phase and no diffusion (or only little diffusion) occurs between the liquid phase and the solid phase matrix. The distribution of elements in the brazed joint is shown in Fig. 4.6, at which point the brazed joint is composed entirely of eutectic phases. With the increase of holding time, the matrix NiAl starts to dissolve into the liquid phase continuously, so that the original Al-free Ni–Si–B eutectic liquid phase starts to contain Al, and continuously increases its Al content. When the holding time is 5 min, the average Al content of the eutectic tissue in the NiAl/Ni–Si–B/Ni brazed joints is about 2% (Fig. 4.7) and starts to grow epitaxially from the Ni matrix to the liquid phase for isothermal solidification. Due to the short holding time, the resulting joints are still mainly eutectic except for some of the isothermal solidification tissue with Ni epitaxial growth. A boride zone is formed in the Ni matrix near the interface due to the diffusion of B, and its width is equivalent to the diffusion depth of B in the Ni matrix. It can also be seen from Fig. 4.7 that an Al-poor zone is formed in NiAl near the interface due to the diffusion of Al into the liquid phase. The composition distribution of NiAl/Ni diffusion bonded joints after 2 h holding time is shown in Fig. 4.8, at

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Fig. 4.6 Composition distribution of NiAl to Ni brazed joints (after holding at 1065 °C for 0 min). M—Ni matrix; E—Eutectic; I— NiAl matrix

Fig. 4.7 Composition distribution of NiAl and Ni brazed joints (after holding at 1065 °C for 5 min). M—Ni substrate; P—epitaxially grown pre-eutectic; I—NiAl substrate; T—Al-poor transition zone

which time the eutectic phase in the joint has completely disappeared (isothermal solidification phase has ended), but the borides in the Ni matrix near the interface are still present. Test results show that it is difficult to obtain a joint without borides even after a longer holding time, which indicates that the homogenization process is controlled by the diffusion of B elements. The diffusion of element B is related to the composition of the matrix, and the ability of B to diffuse in the NiAl intermetallic compound is much slower than in B. Therefore, epitaxial growth of the NiAl matrix into the liquid phase is also much more difficult than in Ni, as shown in Fig. 4.9. In addition, since there is no Al in the original composition of the eutectic liquid phase, the epitaxial growth of NiAl into the liquid phase must be preceded by a sufficient amount of Al into the liquid phase. The epitaxial growth of Ni into the liquid phase is much easier because that it is not necessary to have Al in the liquid phase as a prerequisite, since a large amount of Ni is already present in the liquid phase. So it is more difficult to diffusion braze

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189

Fig. 4.8 Composition distribution of NiAl/Ni diffusion welded joints (after holding at 1065 °C for 2 h). M—Ni substrate; P—epitaxially grown pre-eutectic; E—eutectic; I—NiAl substrate; T—Al-poor transition zone

Fig. 4.9 Effect of holding time on the width of eutectic zone in different diffusion connected joints. 1—Ni/Ni–Si–B/Ni; 2—NiAl/Ni–Si–B/Ni; 3-NiAl/Ni–Si–B/NiAl

NiAl/NiAl with amorphous BNi-3 brazing material than to braze Ni/NiAl, while diffusion brazing Ni/NiAl is more difficult than brazing Ni/Ni. (2) Transition liquid phase diffusion bonding of NiAl When transition liquid-phase diffusion bonding of domestic NiAl intermetallic compounds (such as IC-6 alloy), the interlayer composition was adjusted on the basis of the base material by removing Al from the base material, adding about 7% Cr to improve oxidation resistance, and also adding 3.5% to 4.5% B to make a powder layer of 0.1 mm. The diffusion heating temperature was 1260 °C and the isothermal solidification and composition homogenization time was 36 h. The resulting diffusion bonded joints can last up to 100 h under 980 °C and 100 MPa tensile force. A typical application of the transition liquid-phase diffusion bonding method is the development of a NiAl single crystal foliate blade for GE company in the USA, with the manufacturing process shown in Fig. 4.10. The solid blade is cast first, and the blade is cut from the middle using EDM wire cutting. Then the two halves are machined with a cavity structure inside the blade, and the final process is to join the

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Fig. 4.10 Fabrication process of NiAl single crystal alloy blades. a solid casting b wire cutting c machining of internal cavities. d machining of the connecting intermediate layer e instantaneous liquid phase diffusion welding f final machining

two halves together. A transition liquid phase diffusion bonding technique is used to obtain joints with comparable mechanical properties to NiAl single crystals.

4.2.2 Fusion Welding of Ni3 Al Intermetallic Compounds (1) Electron beam welding of Ni3 Al When using the energy-controlled electron beam welding of Ni3 Al-based alloys, welded joints without cracks can be obtained at low welding speeds. Chemical compositions of the two Ni3 Al-based alloys containing Fe used in the tests are shown in Table 4.7. The generation of weld cracks is mainly related to the welding speed and the B content in Ni3 Al-based alloys, and the weld crack rate increases significantly with Table 4.7 Chemical composition of Ni3 Al-based alloys containing Fe Alloy

Chemical composition (molar fraction)/% Ni

Al

Fe

B

Other

IC-25

69.9

18.9

10.0

0.24 (0.05%)

Ti 0.5 + Mn 0.5

IC-103

70.0

18.9

10.0

0.10 (0.02%)

Ti 0.5 + Mn 0.5

Note: Numbers in parentheses are mass fractions

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191

increasing welding speed. Electron beam welding speed on two Ni3 Al-based alloys (IC-103, IC-25) crack rate is as shown in Fig. 4.11. When welding speed exceeds 13 mm/s, IC-25 alloy is sensitive to cracking. Element B plays a favorable role in improving the room temperature plasticity of Ni3 Al. The addition of B can improve the bonding of grain boundaries. But the tendency to crack increases when the B content exceeds a certain limit (see Fig. 4.12), and the B content at the lowest weld cracking rate is about 0.02%. As seen in Fig. 4.11, when the B content was reduced from 0.05% in IC-25 alloy to 0.02% in IC-103 alloy, the weld crack was completely eliminated. The weld crack Fig. 4.11 Effect of welding speed on cracking during electron beam welding of Ni–Fe aluminides

Fig. 4.12 Effect of B on Ni3 Al welding thermal cracking tendency

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was never present in IC-103 alloy when the welding speed was all the way up to 50 mm/s. B has a similar role in Ni-based high-temperature alloys. The addition of trace amounts of B in Ni-based high-temperature alloys can strengthen grain boundaries and improve high-temperature strength. But excess B tends to form brittle compounds at grain boundaries and may be low-melting, which can lead to local melting and reduced thermoplasticity in the heat affected zone and cause liquefaction cracks in the heat affected zone. However, no local melting was found in the heat affected zone of Ni3 Al and no liquid phase was observed on the crack surface. Therefore, a moderate reduction in B content is necessary to improve the weldability of Ni3 Al alloys, although it has some effect on room temperature plasticity. Based on the thermoplastic change curves of the two alloys IC-25 and IC-103 during heating process measured from the Gleeble-1500 thermal simulation tester (Figs. 4.13 and 4.14), it can be seen that there is a great difference between 1200 and 1250 °C. The elongation in tension at 1200 °C for IC-25 and IC-103 is 0.5% and 16.1%. The fracture morphology of IC-25 alloy is brittle intergranular fracture, but the fracture of IC-103 alloy is characterized by plastic fracture and exhibits high tensile ductility. The fracture morphology of Ni3 Al-based alloys is closely related to the bond strength at the grain boundaries. When the grain boundary bond strength is lower than the yield strength of the material, the fracture morphology is a non-ductile intergranular fracture, and the fracture strain increases with the increase of the grain boundary bond strength. The fracture strain of IC-103 alloy at 1200 °C is much higher than that of IC-25 alloy, and the grain boundary bond strength of IC-103 alloy is much higher than that of IC-25 alloy at this time. This also suggests that the effect of B on the high temperature plasticity of the Ni3 Al-based alloy containing Fe is not consistent with its effect on room temperature plasticity.

Fig. 4.13 Tensile plasticity of IC-25 alloy at elevated temperatures as a function of temperature

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Fig. 4.14 Tensile plasticity versus temperature for IC-103 alloy at elevated temperature

Although boron (B) significantly increases the room-temperature plasticity of Ni3 Al, the effect is not significant at high temperature, especially in the midtemperature range of 600 to 800 °C. The boron-containing Ni3 Al-based alloy has a brittle temperature region, which is a dynamic embrittlement phenomenon related to the oxygen content of the test ambient atmosphere. Thus, the higher heat affected zone cracking tendency exhibited by IC-25 alloys with high B content in electron beam weld joints with welding speed above 13 mm/s is due to their intergranular embrittlement at high temperature and the action of thermal stresse. (2) Electrode arc welding of Ni3 Al alloy When Ni3 Al alloy is welded by electrode arc welding, the selection of welding consumables is very important, and the selection of a reasonable welding consumable can compensate for the poor weldability of Ni3 Al alloy and reduce or eliminate weld cracks. The Ni3 Al base material cannot be used as a welding material because it is highly susceptible to cracking during welding. The high temperature alloy Ni818 is a more suitable welding material for Ni3 Al alloy, which can achieve crack-free welding of Ni3 Al structural parts. The main composition of this welding material is based on the Ni base, adding 0.04C–15Cr–7Fe–15Mo–3.5W–1Mn–0.25 V. In order to ensure the stability of the welding process and prevent the emergence of welding cracks, the surface oxides, oil, etc. of the welded parts must be removed before welding, to avoid foreign non-metallic inclusions mixed into the welding pool. Under the premise of ensuring the metallurgical requirements of welding, small groove welding should be considered to minimize the size of the weld and control the welding heat affected zone as small as possible. The welding process uses low current and low speed to control the welding heat input and enhance heat dissipation to prevent overheating of the weld pool and coarse phase of the joint area after welding.

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The process parameters for welding of Ni3 Al-based IC-218 alloy using Ni818 welding consumables are shown in Table 4.8. The weld surface was well formed and no surface cracks were found from macroscopic observation after chemical corrosion, and no defects such as weld cracks, internal porosity or slagging were found when the weld was dissected. The strength of the weld reached 450 MPa and was pulled off at the fusion zone, which is a ductile fracture. The alloying of the fusion zone was very complex, thus making the strength of the weld (essentially the strength of the fusion zone) lower than that of the base material and the weld material.

4.2.3 Diffusion Bonding of Ni3 Al to Carbon Steel (or Stainless Steel) (1) Diffusion bonding of Ni3 Al to carbon steel By adding alloying elements such as B, Mn, Cr, Ti and V, Ni–Al intermetallic compounds have good room temperature plasticity and high temperature strength. Joining of Ni3 Al and steel dissimilar material, heat affected zone is prone to cracking using fusion welding methods. Most of the current Ni–Al intermetallic compounds dissimilar material are joined using diffusion bonding and brazing. The content of alloying elements in carbon steel is low, and Ni3 Al and carbon steel can be vacuum diffusion bonded directly without an interlayer. The welding process parameters are shown in Table 4.9. Wettability and compatibility between Ni3 Al and carbon steel are good, and can be combined tightly at the diffusion interface, forming a diffusion transition zone with a thickness of about 20–40 μm. The microhardness distribution of Ni3 Al and carbon steel diffusion bonded joints at a heating temperature of 1400 °C and holding time of 30 min and a heating temperature of 1200 °C and holding time of 60 min is shown in Fig. 4.15. Ni3 Al intermetallic compound has a microhardness of about 400HM. The closer to the Ni3 Al and carbon steel diffusion bonding interface, due to the existence of diffusion microscopic cavities and the diffusion of different element content, resulting in the Ni3 Al crystal structure of disordered transformation, microhardness decreased to 230HM. And in the intermediate zone of Ni3 Al and carbon steel diffusion bonded joints, due to a certain element diffusion during the diffusion bonded process, the organization fine and dense, and the microhardness increases to 500HM, then the microhardness decreases to 200HM (the microhardness of the carbon steel) after diffusion welding. Whether the diffusion bonded joint of Ni3 Al and carbon steel can meet the performance in working conditions mainly depends on the distribution of the various elements from base materials near the interface. The elemental distribution of Ni3 Al and carbon steel diffusion bonded joint under the conditions of heating temperature of 1200 °C, holding time of 60 min and heating temperature of 1000 °C, holding time of 60 min and welding pressure of 2 MPa is shown in Fig. 4.16.

Groove angle/(°)

45

Base material

IC-218

Mechanical polishing

Cleaning before welding 3.2

Welding rod diameter/mm 200

Preheating temperature/°C

Table 4.8 Process parameters for welding IC-218 casting alloy with Ni818 welding consumables

130

Welding current/A

750 °C x 2 h annealed

Post-weld treatment

Multi-storey stacking

Process features

4.2 Welding of Ni–Al Intermetallic Compounds 195

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Table 4.9 Process parameters for diffusion bonding of Ni3 Al to carbon steel Heating temperature/°C

Holding time/min

Heating speed/°C x min−1

Cooling speed/°C x min−1

Bonding pressure/MPa

Vacuum level/Pa

1200–1400

30–60

5

10

2

3 × 10–3

Fig. 4.15 Microhardness distribution of Ni3 Al diffusion bonded joints with carbon steel. 1–1400°C × 30 min; 2–1200°C × 60 min

Fig. 4.16 Elements distribution of Ni3 Al to carbon steel diffusion bonded joints. 1—Ni; 2—Al; 3—Fe

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Table 4.10 Process parameters for diffusion welding of Ni3 Al to stainless steel Heating temperature/°C

Holding time/min

Heating speed/°C x min−1

Cooling speed/°C x min−1

Bonding pressure/MPa

Vacuum level/Pa

1200–1380

30–60

20

30

0–9

3.4 × 10–3

While heating temperature of 1200 °C, holding time of 60 min, Ni, Al, Fe element concentration changes of Ni3 Al and carbon steel diffusion bonded joint are mainly reflected in the grain boundary. Diffusion in the grain boundary plays a major role. At the diffusion interface, the recrystallized grains are larger and the elemental concentration fluctuates less, only in a tiny area of the joint near the carbon steel side, the Ni, Al, Fe elemental concentration changes abruptly to the initial concentration values of the elements in the carbon steel base material. While heating temperature of 1000 °C, holding time of 60 min, welding pressure of 2 MPa, the temperature is lower, recrystallization phenomenon is less likely to occur, grain growth is slower, and the role of pressure make the volume of diffusion dominates between Ni3 Al and carbon steel grains, the element concentration changes fluctuate more. (2) Diffusion welding of Ni3 Al with stainless steel Ni3 Al intermetallic compounds have higher resistance to high temperature and corrosion than stainless steel. In some cases where high temperature corrosion resistance is required for the parts, sometimes Ni3 Al intermetallic compounds are to be joined with stainless steel. Studies have shown that Ni3 Al and stainless steel can be vacuum diffusion bonded directly without adding an interlayer, and the process parameters are shown in Table 4.10. The microhardness distribution of Ni3 Al diffusion bonded joints with stainless steel at a heating temperature of 1380 °C and a holding time of 30 min versus a heating temperature of 1200 °C and a holding time of 60 min is shown in Fig. 4.17. Fig. 4.17 Microhardness distribution of Ni3 Al diffusion bonded joints with stainless steel. 1–1380 °C × 30 min; 2–1200°C × 60 min

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Fig. 4.18 Elemental distribution of Ni3 Al with stainless steel diffusion welded joints (1200 °C x 60 min). 1—Ni; 2—Al; 3—Fe

The microhardness of diffusion bonded joint of Ni3 Al and stainless steel increases to a maximum of 450HM, and near the stainless steel side, the microhardness decreases to the microhardness value of 220HM of the stainless steel base material. The microhardness of the entire diffusion bonded joint of Ni3 Al and stainless steel varies continuously, which is mainly related to the continuity of microstructure at the joint, the continuous growth of grains and changes in element concentration. The elemental distribution of diffusion bonded joint of Ni3 Al and stainless steel at a heating temperature of 1200 °C and holding time of 60 min is shown in Fig. 4.18. More alloying elements existed in stainless steel. In the diffusion bonding process of Ni3 Al and stainless steel, the diffusion path of elements is more complex, the interaction between the elements is large. So in the diffusion bonded joint of Ni3 Al and stainless steel, element concentration changes fluctuate, the formation of intermediate compounds is also more complex.

4.2.4 Diffusion Bonding and Vacuum Brazing of Ni3 Al (IC10) Alloys The Ni3 Al-based IC10 alloy is a directional solidification multifaceted composite reinforced high-temperature alloy developed in China, mainly used in the guide vanes of aero-engines, which requires welded connections during its manufacture. The chemical composition and high-temperature mechanical properties of the Ni3 Albased IC10 alloy are shown in Table 4.11. (1) TLP diffusion bonding of Ni3 Al-based IC10 alloy This IC10 alloy is cast by the directional solidification method, with γ + γ’, biphase, of which the γ, phase in a blocky distribution and the γ phase in a reticulation around

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Table 4.11 Chemical composition and high temperature mechanical properties of Ni3 Al-based IC10 alloy Chemical composition (mass fraction, %) Co

Cr

Al

W

Mo

Ta

Hf

B

Ni

11.5–12.5

6.5–7.5

5.6–6.2

4.8–5.2

1.5–5.0

6.5–7.5

1.3 to 1.7

≤ 0.02 Residuals

High temperature mechanical properties Status

Solid solution

Durable strength at 980 °C R100 /MPa

1100 °C lasting strength R100 /MPa

Vertical

Horizontal

Vertical

Horizontal

160

80

70

40

the γ, phase. After holding at (1260 ± 10)°C for 4 h and then oil cooling, or air cooling treatment for structure homogenization, it is still γ phase in a reticulation around γ, phase. γ phase is about 20–30%, γ, phase is about 65–75%, and there are small amounts of borides and carbides. KNi-3, YL alloy are used as intermediate layers in the diffusion bonding process. (1) Use of KNi-3 as an interlayer Diffusion bonding process parameters: bonding temperature of 1230 to 1250 °C, holding time of 4 h and 10 h, respectively. ➀ Microstructure of the joint When the holding time is 4 h, the diffusion bonded joint consists of γ-phase matrix, large γ, phase, massive borides and a small amount of carbides. When the holding time is 10 h, microstructure of the joint is large γ ‘phase, massive borides and a small amount fine carbide, uniformly distributed in the γ phase matrix, microstructure of the joint and the parent material is basically similar, the bonding form is good. ➁ Mechanical properties of joints The strength of joints at room temperature is 705–894 MPa (average value is 772 MPa); the strength of joints at 980°C is 530–584 MPa (average value is 561 MPa), and the elongation after fracture is 1.2–2.8% (average value is 2.23%.) The high temperature lasting strength at 980°C and 100 h is 120 MPa, which reaches 80% of the base material. The fracture morphology at room temperature is dominated by small tough nests, with cleavage surfaces distributed in the nests and little macroscopic fracture undulation, with more W, Mo, Co and Hf elements in the cleavage surfaces and less W, Mo, Co and Hf elements in the tough nests. The fracture morphology at high temperature is dominated by small tough nests, with large macroscopic fracture undulation. (2) Use of YL alloy as an interlayer YL alloy is used as a special interlayer material for transition liquid phase diffusion bonding (TLP) of Ni3 Al-based IC10 alloy with a chemical composition similar to

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that of IC10 alloy with the removal of Hf and C and the addition of B, which is added to lower its melting point. ➀ Diffusion bonding process. Diffusion bonding process parameters: bonding temperature is 1270 °C (solid solution temperature of base material), holding time is 5 min, 2 h, 8 h and 24 h respectively. ➁ Microstructure characteristics of the joint. (i) Changes in the microstructure during diffusion bonding Using YL alloy as an interlayer, a good diffusion bonded joint can be formed with a wide weld seam under very short holding time, and a flower cluster γ + γ’, eutectic (black phase in the center of the weld) is formed at the interface contacted with the IC10 base material, as well as a fishbone compound (borides) and a large mesh-like phases (Ni-Hf eutectic). After 2 h of holding temperature, the weld structure has been basically consistent with the base material and the weld width has become narrower, except for some borides at the edges of the γ + γ, eutectic. After 8 h of holding temperature, the weld width was further narrowed. After 24 h of holding time, the joint structure has been homogenized and no junction between the weld and the base material can be seen. The formation process of Ni3 Al-based IC10 alloy transition liquid-phase diffusion bonded joint is as follows: firstly, the interlayer alloy melts, and because the interlayer alloy contains γ, phase forming elements such as Al and Ta, and elements such as Hf and B that lower the melting point can promote the formation of eutectic, a large number of continuous flake-like γ + γ, eutectic is formed at the interface between the interlayer and the parent material close to both sides, thus discharging Cr, Mo, W and other elements, and borides of Cr and Ta are formed around the eutectic. The time of this process is very short and the weld width has exceeded the thickness of the interlayer, indicating that some of the base material has been dissolved. At the same time, the mutual diffusion of elements occurs between the interlayer and the base material, and the B in the interlayer diffuses into the base material, which makes the base material melt with a lower melting point, and a large number of borides are formed during cooling. With the increase of holding time, the B content in the near seam area gradually decreases and the structure tends to be homogenized because of the small diameter of B atoms and easy diffusion. (ii) Changes in the morphology of the γ, phase When the holding time is short, the shape of γ, phase is nearly spherical. After increasing the holding time, it gradually becomes tetragonal and some field shape, and the grains also grow up. This is because the precipitation of γ, phase is controlled by the interfacial energy and co-grid deformation energy, and the short holding time is too late to grow, so it is spherical. With the extension of the holding time, the γ, phase grows up and will destroy the co-grid, while forming part of the co-grid interface, and the shape tends to be square to reduce the co-grid elastic energy.

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In high temperature alloys, Al, Ti, Nb, Ta, V, Zr and Hf are γ, phase forming elements, while Co, Cr and Mo are γ, phase forming elements, and W is roughly distributed in both γ, and γ phases, so the mass fraction of (Al + Ti + Nb + Ta + V + Zr + Hf + 1/2W) can be used as the γ, phase forming factor. The bigger the γ, phase forming factor is, the more γ, phase is. Since Hf is removed from the interlayer, the formation factor of γ, phase is only Al, Ta and W, so the formation factor of γ, phase is not big. While the holding time is 5 min, a large number of borides are formed in the weld, and the Hf in the base material diffuses into the weld, forming Ni-Hf eutectic, so the formation factor of γ, phase in the weld is smaller, and the γ, phase content is less, and the size is also smaller, and it is easy to become spherical. After increasing the holding time, the composition of the weld tends to be uniform and basically consistent with the base material, the formation factor of γ, phase increases, the content of γ, phase also increases, and the size becomes larger and becomes quadrilateral. (2) Vacuum brazing of Ni3 Al-based IC10 alloy with nickel-based alloys Ni3 Al-based IC10 alloy was joined to GH3039 nickel-based alloy by vacuum brazing. The chemical composition and high temperature mechanical properties of this GH3039 nickel-based alloy are shown in Table 4.12. Since Ni3 Al-based IC10 alloy is produced by the cast method and has an uneven surface, filling its large gaps requires the use of Rene’95 high temperature alloy powder. The chemical composition of this alloy powder see Table 4.13. The brazing material is Co50CrNiWB. (1) Brazing process The process parameters for brazing Ni3 Al-based IC10 alloy and GH3039 nickelbased alloy were: heating temperature of 1180 °C, holding time of 30 min, gaps of 0.1 mm and 0.5 mm, and vacuum of 5 × 10–2 Pa. (2) Microstructure of the brazed joint ➀ Narrow gap (0.1 mm) brazing There is no longer a clear boundary between the solid solution matrix of the braze and the base material of GH3039 alloy, and skeletal borides are distributed continuously on the solid solution matrix of the braze seam. The continuously distributed skeletal gray phase is the Cr-rich boride phase, and the black massive phase is probably TiN. ➁ Large gap (0.5 mm) brazing There is no boundary between the braze seam and the base material of GH3039 alloy. The braze seam between Rene’95 high-temperature alloy powder is a solid solution matrix with a large number of skeletal boride phases, which are divided into white and gray skeletal boride phases, with the white skeletal boride phase being a W-rich boride phase and the gray skeletal boride phase being a Cr-rich boride phase. The brazed joints between Rene’95 high temperature alloy powders are Ni–Cr solid solution matrix. The structure of the large gap brazed joints of Ni3 Al and GH3039 nickel based alloy are a little more complex.

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0.35–0.75

1.80–2.30

19–22

0.90–1.30

Nb

Fe ≤3.0

C ≤0.08

68 –

161 –

Elongation/%

Solid solution 1170 °C

Tensile strength/MPa

Tensile properties at 900 °C

0.35–0.75

Ti

Solid solution 1080 °C, air cooled

Status

High temperature mechanical properties

Al

Mo

Cr

Chemical composition (mass fraction, %) ≤0.40

Mn

≤0.80

Si

Residuals

Ni

39

34

Durable strength at 900 °C R100 /MPa

Table 4.12 Chemical composition and high temperature mechanical properties of GH3039 nickel-based alloy

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Table 4.13 Chemical composition of Rene’95 high temperature alloy powder C

Cr

Co

W

Al

Ti

Mo

Nb

Zr

B

Ni

0.15

14.0

8.0

3.5

3.5

2.5

3.5

3.5

0.15

0.01

Residuals

(3) Mechanical properties of the brazed joints The tensile properties and 900 °C high temperature endurance properties of brazed joints of Ni3 Al-based IC10 alloy with GH3039 nickel-based alloy using 50CoCrNiWB brazing material, prefilled with Rene’95 high temperature alloy powder, and brazed at 1180 °C, holding time 30 min are shown in Tables 4.14 and 4.15 for normal gap (0.1 mm) and large gap (0.5 mm). The test results showed that in the tensile properties of the brazed joints of Ni3 Albased IC10 alloy with GH3039 nickel-based alloy for normal gap (0.1 mm) and large gap (0.5 mm), the tensile strength of the brazed joints exceeded the tensile strength of the parent material of GH3039 nickel-based alloy (161 MPa), and only specimen No. 903 broke on the brazed joint, the rest broke on the GH3039 nickel The rest were broken on the GH3039 Ni-based alloy base material. The high temperature durability of the brazed joints at 900 °C was well over 100 h, and all of them were also broken Table 4.14 Tensile properties of Ni3 Al-based IC10 alloy brazed joints with GH3039 nickel-based alloy Specimen No Clearance/mm Tensile strength/MPa Elongation/% Remark 901

0.1

185

31

Broken from GH3039, IC10 with minimal elongation

902

0.1

173

21

Mainly GH3039 elongated

903

0.1

180

4.7

Break in brazing seam

907

0.5

169

58

Broken from GH3039, IC10 with minimal elongation

908

0.5

178

55

Mainly GH3039 elongated

Table 4.15 High temperature durability at 900 °C of brazed joints of Ni3 Al-based IC10 alloy with GH3039 nickel-based alloy Specimen no.

Clearance/mm

Test stress/MPa

Longevity/h

Breaking area

904

0.1

40

178.4

GH3039

905

0.1

40

159.8

GH3039

906

0.1

40

199.8

GH3039

909

0.5

40

214.2

GH3039

910

0.5

40

215.5

GH3039

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on the base material of GH3039 Ni-based alloy, indicating that the brazed joints have good high temperature endurance performance.

4.3 Welding of Ti–Al Intermetallic Compounds Ti–Al series intermetallic compounds are valued for their low density and high specific strength, especially for aerospace vehicles. The three intermetallic compounds of this series, Ti3 Al, TiAl and TiAl3 , all have development and application prospects. TiAl intermetallic compounds have a density of 3.9 g/cm3 and can be used at temperatures up to 900 °C, which is very attractive for aerospace applications. TiAl3 intermetallic compound is one of the lowest density (3.45 g/cm3 ) materials in the Ti–Al system, which has high strength and good oxidation resistance at higher temperature, thus attracting attention. Ti–Al intermetallic compounds can be joined by argon arc welding, electron beam welding, diffusion welding, brazing, and other methods.

4.3.1 Welding Characteristics of Ti–Al Intermetallic Compounds TiAl and Ti3 Al have poorer weldability and room temperature plasticity than titanium alloys, and in order to obtain good defect-free welded joints, welding of these alloys should be done with the following in mind. ➀ TiAl and Ti3 Al are highly susceptible to adsorption of interstitial elements such as oxygen and nitrogen, resulting in significant degradation of alloy properties. Therefore the welding melting, solidification crystallization and solid state cooling processes must be carried out in an inert atmosphere or vacuum; compared with the local protection of argon arc welding and laser welding, the high vacuum chamber of electron beam welding and diffusion welding provides a good protective environment. ➁ To prevent contamination of the welded parts, it is important to clean and sanitize the surface of the welded parts. ➂ According to the size of the welded parts and the complexity of the structure, adopt the corresponding welding process. For example, argon arc welding and laser welding can be used for thin parts or parts of medium thickness; electron beam welding and diffusion welding should be used for large-section parts to ensure welding quality. ➃ Considering the residual stresses after welding, a welding process with high energy density should be used to achieve full penetration and complete welding in one pass, avoiding multiple passes of argon arc welding process.

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➄ Thorough understanding of the weld metallurgical process that affect the microstructure and properties of welded joints is required, such as melting of the welded alloy, solidification and crystallization, the continuous cooling pattern of the phase change, precipitates and the effect of post-weld heat treatment. Proper welding processes and post-weld heat treatment are essential to obtain strong welded parts. (1) Phase transformation of Ti–Al intermetallic compounds during heating and cooling. TiAl intermetallic compound is a material with poor room temperature plasticity, but its room temperature elongation can be increased to 2–4% by adding Cr, Mn, V, Mo and other elements for alloying and tissue adjustment to form a certain proportion and morphology of (γ + α2 ) two-phase organization. Therefore, some alloy compositions of TiAl intermetallic compounds are designed to have a lamellar biphasic organization with (γ + α) two phases at room temperature, with the α2 phase in a lamellar shape across the γ phase grains. This biphasic organization is obtained by the co-precipitation reaction of the α phase → (α2 + γ) two phases during cooling. In the Ti–48Al (molar fraction, %) intermetallic compound, the γ phase transforms to α phase in the high temperature range of 1130 to 1375 °C, but the α phase transforms to γ phase very quickly during cooling. For example, in the Ti–48Al–2Cr–2Nb (molar fraction, %) intermetallic compound with Cr and Nb added, quenching from the αphase region at 1400 °C resulted in a transition to the γ-phase, and a bulk γ-phase was obtained, with lamellar organization obtained only at slow cooling. Therefore, the faster cooling rate under welding conditions will disrupt the ideal tissue state of the TiAl intermetallic compound, causing it to transform into a brittle tissue prone to solid-phase (cold) crack formation. In the Ti3 Al intermetallic compound, in addition to the ordered α2 phase, there is a small amount of disordered body-centered cubic β phase, which improves the room temperature plasticity of the Ti3 Al intermetallic compound. Microscopic morphology of the fracture shows a destructive fracture across the grains of the α2 phase, but shows a plastic tearing morphology due to the presence of the β phase at the grain boundaries. Therefore, in order to improve the room temperature plasticity of Ti3 Al intermetallic compounds, some β phase should be retained at the grain boundaries. However, welding thermal cycling tends to destroy this favorable (α2 + β) two-phase structure, which deteriorates the plasticity of the joint after welding. That is, the β-phase obtained at high temperature for Ti3 Al intermetallic compounds will undergo a transformation when cooled to low temperature. Figure 4.19 shows the CCT (continuous cooling) curves of Ti3 Al intermetallic compound and Fig. 4.20 shows the CCT (continuous cooling) curves of a simple α2 phase and a super α2 phase. The organization of the Ti3 Al intermetallic compound after cooling can be predicted using these continuous cooling curves. The room temperature phase of Ti3 Al intermetallic compound in equilibrium should be (α2 + β), which becomes β-phase when heated to high temperature. During the subsequent cooling, the decomposition process of the β-phase is very slow, too

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Fig. 4.19 Ti3 Al intermetallic compound CCT (continuous cooling) graph. ➀ β → BP2 ; ➁ β → α’ + 2 BP2 ; ➂ β → α’ + 2 α2P ; ➃ β → α’2 + α2 + β/B2; ➄ β → (α2 + β) + α’2 + α2 + β

Fig. 4.20 CCT (continuous cooling) profiles for one α2 phase and super α2 phase. ➀ cooling rate 1 °C/s; ➁ cooling rate 10 °C/s;➂ cooling rate 100 °C/s

late for the β-phase → α2 phase transition, and the resulting tissue is a sub-stable body-centered cubic β-phase ordered B2 structure. This tissue is softer and more ductile, but due to the instability of the B2 structure, it may transform into the hard and brittle α2 phase of the martensitic α’2 phase at the normal arc welding cooling rate, while the plasticity of this fine needle-like tissue is almost zero. The TEM morphology of the Ti3 Al intermetallic compound (Ti–14Al–21Nb) at a cooling rate of 100 °C/s is shown in Fig. 4.21. Obviously, to obtain a more desirable (α2 + β) two-phase, the weld must be cooled more slowly, which requires preheating of the workpiece. For example, for a thin plate with a thickness of 3 mm preheating up to 600 °C with a cooling rate below 25 °C/s or post-weld heat treatment is required. Therefore, the cooling rate during

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207

Fig. 4.21 TEM of Ti3 Al intermetallic compound (Ti–14Al–21Nb) at a cooling rate of 100 °C/s

continuous post-weld cooling has a decisive influence on the properties of the Ti3 Al intermetallic compound joint zone. (2) Cracking tendency of Ti3 Al intermetallic compounds (1) Cold cracking of Ti3 Al intermetallic compounds Ti3 Al intermetallic compounds differ from Ni3 Al intermetallic compounds in that Ti3 Al intermetallic compounds have a very narrow critical stress range for thermal crack and therefore a low propensity for thermal crack. Moreover, Ti3 Al intermetallic compounds are more plastic at high temperatures and do not produce liquefaction cracks in the heat affected zone. The main problem in welding Ti3 Al intermetallic compounds is the low plasticity at room temperature and the resulting cold crack. (2) Factors affecting the cold crack sensitivity of Ti3 Al intermetallic compounds. ➀ Base material state and welding method The base material is Ti–24Al–14Nb–1Mo (TD3 alloy), which has a solid solution temperature of 950 °C. The three states are as follows. a. 980 °C + 1 h after forging, air-cooled. b. 980 °C + 1 h after hot rolling, air-cooled. c. 950 °C + 1 h after hot rolling, air-cooled. Their room temperature mechanical properties are shown in Table 4.16. The Ti–Al–Nb alloy wire with high Nb content was used for manual GTAW welding in an argon filled tank and automatic GTAW welding in an atmospheric environment, respectively. (i) Cold crack sensitivity. In manual GTAW welding, no cracking occurred in condition a, while cold cracking accompanied by a loud noise sometimes occurred in condition c. Cold cracks were produced in both a and c states during automatic GTAW welding. The cracks originated in the fusion zone and extended

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Table 4.16 Room temperature mechanical properties of three states of Ti3 Al-based TD3 alloy Material status

After forging, After hot rolling at 980°C After hot rolling, 950°C + 980°C + 1 h, air + 1 h, air cooling treatment 1 h, air cooling treatment cooling treatment

Direction

—

Rolling direction

Vertical rolling direction

Rolling direction

Vertical rolling direction

Tensile strength/MPa

1052

975

921

984

1064

Elongation/%

5.8

10.1

2.3

9.7

3.8

perpendicular to the weld towards the base material on both sides, which was apparently related to the plasticity of the base material, which was better in state a than in state c (the elongation after fracture of the base material was 5.8% and 3.8%, respectively). The effect of the welding method is related to the cooling rate after welding, and since manual GTAW welding is carried out in an argonfilled tank, the cooling rate is slower than that of automatic GTAW welding in the atmosphere, and the residual stresses in the weld zone are smaller than in the latter, so that the former is less susceptible to cold cracking than the latter. (ii) Mechanical properties of joints. The mechanical properties of the manual GTAW joints are: a state with a tensile strength of 919 MPa and an elongation after fracture of 3.1 per cent; and c state with a tensile strength of 817 MPa and an elongation after fracture of 1.2 per cent, both fractured in the heat affected zone near the fusion zone. ➁ Effect of preheating Welding Ti3 Al intermetallic compound (Ti–24Al–14Nb–4V) using manual GTAW, no cold cracks were observed in the preheated welded joints, while cold cracks formed in the unpreheated welded joints. The hardness of the preheated welded joints was lower than that of the unpreheated welded joints. Hydrogen contributes to the cold crack susceptibility of Ti3 Al intermetallic compound welded joints and preheating will promote the escape of hydrogen. Therefore preheating is also an effective measure to prevent cold cracking in Ti3 Al intermetallic compounds. Tests have shown that the fracture of unpreheated welded joints have larger cleavage surfaces and more dense and pronounced river patterns, indicating greater brittleness. (3) Effect of preheating and post-weld heat treatment on joint properties Preheating before welding was able to significantly reduce the crack susceptibility in the Ti3 Al intermetallic compound welded joint area, and the hardness peak in the heat affected zone after preheated welding was also moderated and the joint strength factor increased from about 30% to 78% from non-preheated. The base material had a tensile strength of 820 MPa, yield strength of 584 MPa and elongation after fracture of 17%. The tensile strength of the GTAW welded joint without preheating was only

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209

246 MPa and the tensile strength of the GTAW welded joint after preheating was up to 638 MPa. For welded joints of most metal, post-weld heat treatment is able to reduce residual stress and improve fracture toughness. Likewise, post-weld heat treatment can result in improved microstructure and mechanical properties of the Ti3 Al intermetallic compound weld and heat affected zone. The specific post-weld heat treatment parameters need to be determined according to the thickness of the welded part and the shape and size of the welded structure.

4.3.2 Arc Welding of TiAl Intermetallic Compounds Ti–Al intermetallic compounds can be fusion welded, but Ti–Al intermetallic compounds are prone to crystalline cracking in arc welded joints, and this material has a high tendency to harden, so the mechanical properties of arc welded joints are generally poor. Ti–Al intermetallic compound arc welding (commonly used is the GTAW method) has the advantage of low cost, ease of operation, high productivity and has a promising application in the repair of engineering structural parts, the main problem in welding is the avoidance of cracking. When welding Ti–48Al–2Cr–2Nb (molar fraction, %) intermetallic compounds by the tungsten argon arc welding (GTAW) method, the microstructure of the weld consists of columnar and equiaxed tissue, with a small amount of γ phase. Cracking can be avoided when welding with a larger heat input; however, it is highly susceptible to cracking when welding with a low current or smaller heat input. The hardness of GTAW weld metal is higher than that of the base material, and its room temperature plasticity and strength properties are lower than those of the base material. Cracking can be avoided by using a preheat welding process. If preheating is not performed, improper welding parameters can produce a lot of cracks. When welding cast Ti–48Al–2Cr–2Nb (molar fraction, %) and pressed Ti–48Al– 2Cr–2Nb–0.9Mo (molar fraction, %) using the GTAW method, the cooling rate in the welded joint area is controlled by adjusting the magnitude of the welding current (regulating the welding heat input), and the cracking tendency in the weld can be reduced with increasing heat input. The controlled welding heat input also results in a more desirable weld organization, with less α2 brittle organization and less tendency to dendrite segregation, which helps to optimize the organizational properties of the joint.

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4.3.3 Electron Beam Welding of TiAl Intermetallic Compounds (1) Cracking of welded joints The main problems of electron beam welding of TiAl intermetallic compounds are thermal cracking of the weld and degradation of the mechanical properties of the joint. Electron beam welding has a large depth of fusion and good atmosphere, and the tendency to weld cracking is greatly influenced by the faster cooling rate when using electron beam welding of TiAl alloys. The weld crack susceptibility of electron beam welding of TiAl alloy was studied, and the material used was TiB2 particlereinforced Ti–48Al alloy, containing a volume fraction of 6.5% of the reinforcing phase TiB2 and organized into lamellar α2 + γ clusters, equiaxed α2 and γ grains, and short, coarse TiB2 particles. The process parameters and heat affected zone cooling rate used for electron beam welding of TiAl alloy thin plates are shown in Table 4.17. The effect of cooling rate in the heat affected zone of electron beam welding on the cracking tendency of TiAl alloys is shown in Table 4.18 and Fig. 4.22. The average cooling rate (1400–800 °C) corresponding to the appropriate welding parameters is important in order to obtain a crack-free welded joint. Cracking is not sensitive when the cooling rate in the heat affected zone is below 300 K/s. After the cooling rate exceeds 300 K/s, the cracking sensitivity increases significantly with the increase in cooling rate. Transverse cracks are produced in the weld when the cooling rate exceeds 400 K/s and may extend into the base material on both sides. The fracture morphology of this type of cracking is solid cracking, and there is no sign of hot cracking, which is cold cracking.

Table 4.17 Welding parameters and HAZ cooling rate used for electron beam welding Preheat temperature/°C

Acceleration voltage/kV

Electron beam current/mA

Welding speed /mm x s−1

HAZ cooling rate /K x s−1

27

150

2.2

2

90

27

150

2.5

6

650

27

150

3.5

12

1320

27

150

4.0

12

1015

27

150

6.0

24

1800

170

150

2.5

6

400

300

150

2.2

2

35

335

150

2.5

6

200

335

150

4.0

12

310

470

150

2.0

6

325

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211

Table 4.18 Effect of cooling rate on cracking tendency in the heat affected zone HAZ cooling rate/K x s−1

0

300

700

1000

1800

2700

Crack rate/mm x bar−1

0

0

0.14

0.23

0.45

0.57

Fig. 4.22 Effect of cooling rate in the heat affected zone on cracking rate (cooling from 1400 °C to 800 °C)

Therefore, the cooling rate is the main factor affecting weld cracking when welding TiAl alloys by electron beam welding. When the welding parameters are selected appropriately, a crack-free joint can be obtained by electron beam welding of TiAl alloys. Related studies have shown that the preheating temperature required for electron beam welding to prevent crack generation is 250 °C when the welding speed is 6 mm/s (Fig. 4.23). (2) Tissue transformation of electron beam welding heads The tissue properties of TiAl alloy electron beam welded joints have a very important correlation with the heat input (cooling rate). When the cooling rate is slow, the transformation will occur according to the Ti–Al binary alloy phase diagram. At high temperature, the transformation of β-phase → α-phase occurs first. Then the γ-phase precipitates from the α-phase to form a laminar organization. Finally, the (α2 + γ) biphasic laminar organization and the biphasic organization of the equiaxed γphase are obtained. From the phase diagram of Ti–Al binary alloy, it can be seen that Fig. 4.23 Relationship between preheating temperature and crack rate (welding speed of 0.6 cm/s and 1.2 cm/s)

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the co-precipitation reaction α phase → (α2 + γ) two-phase occurs at a temperature of 1125 °C. When the cooling rate is faster, it transforms into a granular γm phase. The granular transformation is from the α phase to the γ phase of the same composition but different crystal structure, and this granular γm phase is irregular in shape. When the cooling rate is very fast, most of the β-phase crystallized in the weld pool is retained, transforming into an ordered β2 phase retained to room temperature. β2 phase is predominantly light-colored under light microscopy, which is due to the rapid cooling rate, so that impurities and low-melting point eutectics do not have time to migrate to the grain boundaries, so the grain boundaries are not obvious. When welding Ti–48Al–2Cr–2Nb alloy with a thickness of 10 mm using electron beam welding, preheating at 750 °C transforms the weld into a lamellar structure. But during rapid cooling without preheating, the weld is mainly a lumpy transformation structure. Under these high cooling rate conditions, the weld is highly susceptible to cracking and therefore the welding thermal process must be strictly controlled. TiAl alloys also suffer from hydrogen embrittlement, and since the welding methods used now are low hydrogen, hydrogen is not a major problem affecting weld cracking. For electron beam welding of Ti3 Al–Nb (Ti–14Al–21Nb) intermetallic compounds, the effect of different welding parameters (different cooling rates) on the hardness of the welded joint is shown in Fig. 4.24. (3) Example of vacuum electron beam welding of TiAl alloy ➀ Formation of weld The heat input of vacuum electron beam welding is 1.15–2.48 kJ/cm for TiAl alloy with a thickness of 3 mm and a uniform surface melt width, uniform and detailed arc pattern, slight collapse of the weld, and localized macroscopic transverse microcracks, especially at the closing arc crater. The weld width increases with the increase of electron beam welding beam current and decreases with the increase of welding speed. ➁ Mechanical properties of welded joints The effect of hardness distribution and heat input of the electron beam welding on the strength of the joint is shown in Figs. 4.25 and 4.26. When the acceleration voltage is 55 kV, the electron beam current is 24 mA, and the welding speed is 400 mm/min (welding heat input is 1.98 kJ/cm) the electron beam welded joint has the highest strength of 221 MPa, reaching 50.5% of the strength of the TiAl alloy base material (438 MPa). In the molten pool, the metal crystallizes out mainly the β phase, which then transforms into the β2 phase and the ductile (α2 + γ) structure. The welding heat input has a significant effect on the weld structure and therefore on the joint strength. When the welding heat input decreases, the above transformation is not sufficient, the plastic toughness is not good, and therefore the joint strength is not high. As the welding heat input increases, the cooling rate decreases and the β-phase transforms into granular γm phase and (α2 + γ) biphasic laminar structure with increased strength.

4.3 Welding of Ti–Al Intermetallic Compounds Fig. 4.24 Effect of electron beam welding parameters (different cooling rates) on the hardness of Ti3 Al–Nb (Ti–14Al–21Nb) intermetallic compound welded joints

Fig. 4.25 Hardness distribution of the electron beam welding joint.

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213

214

4 Joining of Nickel–Aluminium and Titanium–Aluminium Intermetallic …

Fig. 4.26 Effect of welding heat input on joint strength

The welding heat input is further increased, and due to the increase in melt pool temperature, the alloying elements burn and volatilize seriously, resulting in coarse phase and excessive collapse, leading to a decrease in joint strength. (iii) Fracture path of the joint The test results show that most of the microcracks in electron beam welding of TiAl alloy start at the weld surface and then expand towards the weld and heat affected zone, resulting in fracture of the joint. The microcracks appearing on the weld surface, together with the stress concentration formed by the collapse of the weld, resulted in a poor joint strength. The fracture of the joint is a brittle fracture with a direction of approximating vertical to the tensile stress. The fracture surface has a metallic luster, no shrinkage, and the fracture elongation is almost zero. The fracture is characterized by cleavage fracture and transgranular fracture. As the post-weld cooling rate decreases, the (α2 + γ) biphasic lamellar structure in the weld increases, and the fracture may show delamination and penetration, and the fracture toughness increases compared with the single-phase sreucture.

4.3.4 Diffusion Welding of TiAl and Ti3 Al Alloys (1) Characteristics of TiAl alloy diffusion welding ➀ Direct diffusion Welding Process parameters (temperature, time, pressure, etc.) have a significant effect on the performance of diffusion welded joints of TiAl alloys. The process parameters and joint properties of direct diffusion welding are given in Table 4.19.

4.3 Welding of Ti–Al Intermetallic Compounds

215

During diffusion joining of Ti–48Al duplex casting alloy, the tensile strength of diffusion bonded joints gradually increases with increasing heating temperature, holding time and pressure. Diffusion bonded joints with no interfacial microscopic holes and good interfacial bonding were obtained at 1200 °C, 64 min and 15 MPa pressure, and the room temperature tensile strength of the joints reached 225 MPa, breaking at the base material. High-temperature tensile tests showed (Fig. 4.27) that the tensile strength of diffusion bonded joints at high temperature of 800 °C and 1000 °C decreased, and breaking at the bond interface, the tensile strength was about 180 MPa, which is about 40% lower than that of the base material. The reason for this is that the interface diffusion migration is less and the section is flat. Microstructure of the diffusion bonding interface has a significant effect on the joint properties. General, grain growth occurs after the diffusion bonded joint is treated with vacuum heating. For example, diffusion bonding of TiAl was performed at 1200 °C, 64 min and 10 MPa, and then the joints were vacuum heat treated at 1300 °C, 120 min and 1.3 MPa. Metallographic observation showed that the grain diameter increased from 65 μm in the diffusion bonded state to about 130 μm, and the tensile strength of the joint also decreased. In order to promote interfacial diffusion migration to improve the high temperature tensile strength at 1000 °C, the joint can be subjected to recrystallization heat treatment. The bonded joints obtained from the above vacuum diffusion bonding were subjected to recrystallization heat treatment at 1300 °C × 120 min and 1.3 × 10–3 Pa vacuum, and the grain diameter could be increased from 65 μm in the weld state to 130 μm. The tensile strength of the joints at 1000 °C at this time was 210 MPa, broken from the base material. The effect of vacuum degree on the joint tensile strength of TiAl alloy at 1000 °C during vacuum diffusion welding is shown in Fig. 4.28. It can be seen that increasing the vacuum degree is beneficial to improve the high temperature strength properties of diffusion bonded joints. The temperature and time required for diffusion bonding can be greatly reduced by using superplastic diffusion to join TiAl intermetallic compounds. For Ti–47Al– Cr–Mn–Nb–Si–B alloy, superplastic diffusion joining at a heating temperature of 923–1100 °C, a pressure of 20–40 MPa and a vacuum of 4.5 × 10–4 Pa can result in diffusion bonded joints with good performance and tensile tests broken in the base material. The tests showed that superplastic diffusion bonding of TiAl is easily achieved when the grain size of TiAl intermetallic compound is below 4 μm, the heating temperature is above 880 °C and the deformation rate is 10%. ➁ Diffusion welding with an intermediate layer To improve the performance of TiAl diffusion bonded joints, diffusion bonding can be carried out by adding a transition interlayer. The use of an interlayer can improve surface contact, promote plastic flow and diffusion process. The chemical composition, addition method and thickness of the interlayer have an important influence on the joint properties. The interlayer can be pure metal or an alloy containing active or reduced melting point elements. Table 4.20 gives the common interlayers and process parameters for TiAl diffusion bonding. As can be seen from the table, the

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60 35 64

1000

1000

1000

1200

1100

1250

1200–1350

Ti–52Al

Ti–48Al–2Cr–2Nb

Ti–48Al

Ti–48Al

Ti–47Al

Ti–47Al–2Cr

Ti–48Al–2Mn–Nb

15–45

60

60

60

Holding time/min

Process parameters

Heating temperature/°C

Material to be welded

15

30

30

15

10

10

10

Pressure/MPa

Vacuum

Vacuum

Ar

Ar

Ar

Vacuum

Vacuum

Atmosphere — — — 225 400 530 250

α2 TiO2 , Al2 TiO5 , γ γ + α2 γ, γ + α2 α2 /γ γ, α2

Tensile strength/MPa

γ, γ + α2

Interface product

Table 4.19 Process parameters, interfacial reaction products and tensile strength of joints for Ti–Al diffusion welding

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4.3 Welding of Ti–Al Intermetallic Compounds

217

Fig. 4.27 Tensile strength of Ti–Al diffusion welded joints at different temperatures

Fig. 4.28 Effect of vacuum level on tensile strength of TiAl alloys in joints at 1000 °C

use of an interlayer allows diffusion bonding of TiAl at relatively low temperature and pressure. Using Ti–18Al alloy and Ti–45Al alloy as the interlayer, the diffusion of elements will occur during diffusion bonding, but the strength of the joint is not high. If the heat treatment of 1150–1350°C is carried out after bonding to make sufficient diffusion, the interface structure of the joint tends to be consistent with the base material, and the strength and plasticity of the joint are improved and can reach the level of the base material.

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5–10

1150 900

Ti–48Al–2Cr–2Nb/Ti–15Cu–15Ni/Ti–48Al–2Cr–2Nb

Ti–52Al/Al/Ti–52Al 64

30

Holding time/min

1000

Heating temperature/°C

Process parameters

Ti–52Al/V/Ti–52Al

Material to be bonded (including intermediate layer)

Table 4.20 Intermediate layers and process parameters for TiAl diffusion bonding

10–30

—

15

Pressure/MPa

Vacuum

Vacuum

Vacuum

Atmosphere

— 200

β-Ti + α2 TiAl3 , TiAl2

200

Tensile strength/MPa Al3 V

Interface product

218 4 Joining of Nickel–Aluminium and Titanium–Aluminium Intermetallic …

4.3 Welding of Ti–Al Intermetallic Compounds

219

In addition, a lower melting point Ti–15Cu–15Ni was used as an interlayer for the transition liquid phase bonding of Ti–48Al–2Cr–2Nb alloy, which can well improve the interfacial contact and enhance the performance of diffusion welded joints. Microstructure of TiAl intermetallic compound is very sensitive to mechanical properties. When it contains more alloying elements, the coefficient of linear expansion is low. When welding with dissimilar materials, it is easy to generate large stresses. When using the fusion welding method, the joint composition is complex, and it is very easy to generate brittle intermetallic compounds, and the tendency of thermal cracking is serious. Therefore, TiAl intermetallic compound dissimilar materials are more often joined by diffusion bonding with an interlayer. (2) Diffusion bonding of Ti3 Al alloy Ti3 Al alloy can be diffusion bonded to achieve its connections. Figure 4.29a shows the effect of bonding temperature on the shear strength of diffusion bonded joints of Ti3 Al alloy at a bonding pressure of 9 MPa and a holding time of 30 min. In the bonding temperature range of 800 to 840 °C, the shear strength of the joint is low and changes slowly; when the bonding temperature exceeds 840 °C, the shear strength of the diffusion bonded joint increases rapidly, reaching 751 MPa at 940 °C. Figure 4.29b shows the effect of holding time on the shear strength of Ti3 Al alloy diffusion bonded joint at bonding temperature of 990 °C and a pressure of 12 MPa. It can be seen that as the holding time is extended from 15 to 30 min, the shear strength of the diffusion bonded joint increases rapidly. When the holding time exceeds 30 min, the shear strength of the joint increases at a slower rate. When the holding time is 70 min, the shear strength of the joint approaches that of the base material. When the holding time continues to increase, the shear strength of the joint decreases due to grain coarsening and growth. The heating temperature for diffusion bonding of Ti3 Al alloy is usually around 1000 °C. The required holding time depends on the heating temperature and pressure. The relationship curve between diffusion bonding temperature and holding time for

Fig. 4.29 Effect of diffusion bonding temperature and time on the shear strength of Ti3 Al alloy joints

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4 Joining of Nickel–Aluminium and Titanium–Aluminium Intermetallic …

Ti3 Al alloy is shown in Fig. 4.30. It can be seen that the holding time for diffusion welding can be reduced with the increase in bonding temperature at constant pressure. Figure 4.31 shows the relationship curve between diffusion bonding pressure and holding time for Ti3 Al alloy with a bonding temperature of 980 °C. The top right of the curve shown in the figure is the completely bonded area, and the diffusion welding parameters in the lower left interval do not result in a completely bonded joint. From the curves shown in the figure, it can be seen that increasing the diffusion welding pressure can accelerate the interface diffusion and shorten the diffusion joint time. However, too high pressure brings some other adverse effects on diffusion bonding, such as deformation, etc. Therefore, in practical applications, the reasonable coordination of process parameters should be taken into account, and generally do not use bonding parameters with very high pressure. Fig. 4.30 Ti3 Al diffusion holding time versus bonding temperature

Fig. 4.31 Ti3 Al diffusion holding time versus pressure

4.3 Welding of Ti–Al Intermetallic Compounds

221

4.3.5 Diffusion Bonding of TiAl Dissimilar Materials TiAl and structural steel or ceramic materials can be diffusion bonded with an alloy interlayer. The room temperature tensile strength of the joints can be more than 60% of that of the TiAl intermetallic base material. (1) Diffusion bonding of TiAl and 40Cr steel ➀ Welding process and parameters The different of chemical composition of TiAl intermetallic compound and 40Cr steel is large, and the compatibility is poor. While being diffusion bonded, pure Ti foil, V foil and Cu foil can be used as an intermediate layer. Before bonding, the oil and rust on the to-be-welded surface of TiAl intermetallic compound and 40Cr steel were removed by mechanical or chemical methods, and then assembled in the order of TiAl/Ti/V/Cu/40Cr and immediately put into the vacuum furnace. The thicknesses of the interlayer pure Ti foil, V foil and Cu foil are 30 μm, 100 μm and 20 μm respectively. The diffusion bonding process parameters are: bonding temperature of 950– 1000 °C, pressure of 20 MPa, and holding time of 20 min. ➁ Mechanical properties of diffusion bonded joints The effect of bonding temperature and alloy layer composition on the tensile strength of diffusion bonded joints of TiAl and 40Cr steel are shown in Fig. 4.32. Under the same diffusion bonding process parameters, the tensile strength of TiAl/40Cr steel diffusion bonded joints obtained by using Ti/V/Cu interlayer is higher than that of joints with V/Cu interlayer. And with the increase of bonding temperature, the tensile strength of diffusion bonded joints gradually increases. Because when the temperature is low, the strength of the bonded material matrix is still very high. Under the same pressure conditions, the contact surface plastic deformation is not enough, so the physical contact of the bonded interface is not sufficient. There may be a large number of defects at the diffusion welding interface, and a good metallurgical Fig. 4.32 Effect of bonding temperature on the tensile strength of TiAl/40Cr diffusion bonded joints

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4 Joining of Nickel–Aluminium and Titanium–Aluminium Intermetallic …

bond is not formed. With the increase of temperature, the yield strength of the bonded material decreases sharply. The area of physical contact between the bonded surfaces increases rapidly, and the bonding rate increases. As seen by the fracture composition analysis of TiAl/40Cr steel diffusion bonded joint (Table 4.21), the fracture of TiAl/40Cr steel diffusion bonded joint with Ti, V and Cu as the interlayer occurs at the interface between TiAl and the interlayer Ti foil. And the fracture of TiAl/40Cr steel diffusion bonded joint with V, Cu as an interlayer occurs at the TiAl and interlayer V foil interface. (iii) Microstructure near the diffusion interface The energy spectrum analysis of TiAl/40Cr steel diffusion bonded joints with Ti, V and Cu as interlayers is shown in Table 4.22. X-ray diffraction analysis shows that using Ti, V, and Cu as interlayers, Ti3 Al intermetallic compounds generates near the TiAl side and α-Ti solid solution on the Ti–rich side. These formed phases do not change with temperature, but the thickness of the diffusion reaction layer increases gradually as the bonding temperature increases and the elements diffuse more fully. At the contact interface between Cu foil and 40Cr steel, there is no obvious intermetallic compound forming transition layer and no stable transition plateau in element concentration. This is the main reason why the fracture of TiAl/40Cr steel diffusion bonded joint with Ti, V and Cu as the interlayer occurs at the interface between TiAl and Ti foil. While using V, Cu as an interlayer, the energy spectrum analysis of TiAl/40Cr steel diffusion bonded joint was found that Ti3 Al generate on the side of the joint near TiAl and Al3 V on the V side, which increased the brittleness at the interface of TiAl and V foil and easily caused the brittle fracture of TiAl/40Cr steel diffusion bonded joint. (2) Diffusion bonding of TiAl and SiC ceramics ➀ Welding process and parameters Before diffusion bonding of TiAl and SiC ceramics, the surfaces to be welded of TiAl alloy with 53% Al content and sintered SiC ceramics containing 2% to 3% Al2 O3 were scrubbed with acetone, then rinsed with water + alcohol and air dried. The weldments were then assembled in the order of SiC/TiAl/SiC from bottom to top, while a piece of mica was placed on each of the unbonded surfaces of the top and bottom SiC to prevent the SiC from being joined to the pressurized indenter. The diffusion bonding process is heated by resistance radiation heating. The process parameters for diffusion bonding of TiAl and SiC ceramics are: bonding Table 4.21 Composition analysis of fracture of TiAl/40Cr steel diffusion bonded joint/% Joint

Ti

Al

Cr

Nb

V

Cu

Fe

Ti, V and Cu are intermediate layers

50.19

45.96

2.02

1.83

–

–

Margin

67.90

25.31

3.19

3.60

–

–

–

39.25

38.97

–

2.07

19.71

Margin

–

V and Cu are intermediate layers

4.3 Welding of Ti–Al Intermetallic Compounds

223

temperature of 1300 °C, holding time of 30–45 min, welding pressure of 35 MPa, and vacuum of 6.6 × 10–3 Pa. ➁ Mechanical properties of diffusion bonded joint The chemical composition of three reaction layers of TiAl/SiC diffusion bonding joint are shown in Table 4.23. The chemical composition of the elements within the reaction layer varies greatly, making different structures between TiAl and SiC diffusion bonded joint. With the extension of the holding time, the thickness of the reaction layer in the diffusion bonded joint increases, and can reach a stable state in a certain period of time, so that the joint has a certain strength. The shear strength of the TiAl and SiC diffusion bonded joint at different holding time is shown in Fig. 4.33. Test results of TiAl and SiC diffusion bonded joints show that the shear strength of TiAl and SiC joints starts to decrease rapidly with increasing holding time when the bonding temperature is 1300 °C, and then decreases slowly and stabilizes after 4 h. The joint strength reaches 240 MPa when the holding time is 30 min. The chemical composition of the shear fracture of TiAl and SiC diffusion bonded joints analyzed by electron probe is shown in Table 4.24. The location of shear fracture in TiAl and SiC diffusion bonded joints changes with the holding time. At the holding time of 30 min, the formed TiC layer is very thin (0.58 μm) and the shear strength of the joint depends on the TiC + Ti5 Si3 Cx layer, with fracture occurring at the interface between the (TiAl2 + TiAl) and (TiC + Ti5 Si3 Cx ) layers. Table 4.22 Energy spectrum analysis of TiAl to 40Cr steel diffusion welded joints % Joint

Location

Ti

Al

Cr

Nb

V

Ti, V and Cu are intermediate layers

Near TiAl side

74.3

25.3

0.33

0.10

–

Near Ti side

95.5

0.21

0.09

0.17

–

V and Cu are intermediate layers

Near TiAl side

60.94

21.34

–

–

17.18

Near V side

16.62

68.89

0.54

–

14.49

Fig. 4.33 Shear strength of TiAl and SiC diffusion bonded joints at different holding time

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Table 4.23 Chemical composition of reaction layers of TiAl and SiC diffusion bonded joint % Reaction layer

Ti

Al

Si

C

Cr

1

33.5

62.4

0.8

2.1

1.2

2

54.2

4.4

28.8

12.3

0.3

3

44.3

10.2

5.3

40.1

0.1

Table 4.24 Results of electron probe analysis of shear fractures of TiAl and SiC diffusion bonded joints % Holding time/h

Ti

Al

C

Si

Surface phase

0.5

53.6

5.4

11.1

29.9

Ti5 Si3 Cx

53.1

5.8

10.8

30.3

Ti5 Si3 Cx

46.2

47.8

5.6

0.4

TiAl

54.1

6.2

10.2

29.5

Ti5 Si3 Cx

43.1

8.2

44.2

4.5

TiC

43.8

8.7

43.4

4.1

TiC

44.1

7.9

45.6

2.4

TiC

44.5

8.1

44.8

2.6

TiC

8

Although TiC is a high strength phase and is compatible with the SiC lattice, its strength decreases and becomes a fracture-prone layer when the thickness of the TiC layer is large and a certain number of Al atoms are dissolved. The TiC layer increased to a certain thickness (2.75 μm) and dissolved more Al atoms when the holding time was 8 h. The fracture strength of the joints depends on the thickness of the TiC layer and thus fracture occurs within the corresponding TiC single-phase layer. TiAl and SiC diffusion bonded joints are required to have a certain high temperature strength if they are in a high temperature working environment. As the test temperature increases, the shear strength of TiAl/SiC diffusion bonded joints decreases slightly, and the shear strength can still be maintained at 230 MPa at the test temperature of 700 °C. When the test temperature is higher than 700 °C, the sensitivity of the high temperature shear strength of TiAl and SiC diffusion bonded joints to the test temperature decreases. Therefore, as long as the TiAl/SiC diffusion bonded joint has sufficient shear strength at 700 °C, it can meet the requirements of guaranteed strength performance for use. ➂ Microstructure of diffusion bonded joints The strength of TiAl and SiC diffusion bonded joints and their destruction in service depend on the tissue structure formed in the joint area after diffusion bonding. The reaction layer of TiAl/SiC diffusion bonded joints near the TiAl side forms mainly (TiAl2 + TiAl), and the reaction layer near the SiC ceramic side forms a single phase TiC. The intermediate reaction layer forms mixed phases of TiC and Ti5 Si3 Cx . Therefore, the structure of TiAl/SiC diffusion bonded joint is (TiAl2 + TiAl), (TiC +

Bibliography

225

Ti5 Si3 Cx ) and then transition to TiC from TiAl to SiC ceramic. The above structure can be obtained by controlling the process parameters to meet the requirements of TiAl/SiC diffusion bonded joint.

Bibliography 1. Ren J, Wu A (2000) Advanced materials connection. Beijing: Machinery Industry Press 2. Feng J, Li Z, He P, et al (2003) Interfacial structure and phase growth of TiAl/40Cr diffusion connected joints. Chin J Nonferrous Metals 13(1):162–166 3. Yu Qi, Chunyuan S (2016) Welding of intermetallic compounds, Beijing: Mechanical Industry Press 4. Yonggang Z, Yafang H, Gguoliang C et al (2001) Structural materials of intermetallic compounds. National Defense Industry Press, Beijing 5. Zenglong Z, Hengqiang Y (1992) Intermetallic compounds (Proceedings of the First National Symposium on High Temperature Structured Intermetallic Compounds). Machinery Industry Press, Beijing 6. Dechun G, Wangyue Y, Min D et al (2000) Weldability of Fe-Al based intermetallic compounds. J Metals 36(1):87–92 7. Mckamey CG, Devan JH, Tortorelli PF et al (1991) A review of recent development in Fe3 Al-based alloy. J Mater Res 6(8):1779–1805 8. Jjianting G, Chao S, Minghui T et al (1990) Effect of alloying elements on the mechanical properties of Fe3 Al and FeAl alloys. J Metals 26A(1):20–25 9. Sun Z (2001) Advances in welding research on Fe3 Al-based intermetallic alloys. Materials Guide 15(2):10 10. David SA, Horton JA, Mckamey CG (1989) Welding of iron aluminides. Weld J 68(9):372–381 11. David SA, Zacharia T (1993) Weldability of Fe3 Al-Type Aluminide. Weld J 72(5):201–207 12. Yajiang L, Juan W, Yansheng Y, et al (2002) Phase constitution near the interface zone of diffusion bonding for Fe3 Al/Q235 dissimilar materials. Scr Mater 47(12):851–856 13. Yansheng Y, Zhongliang S, Junyou L (1996) Iron-aluminum intermetallic compounds— alloying and composition design. Shanghai Jiaotong University Press, Shanghai 14. Wang C, Zhu D, Lu L (2007) Progress in the study of intermetallic compounds Fe3 Al. Mater Guide 21(3):67–69 15. Xingquan Y, Yangshan S, Haibo H (1995) Effect of rolling processing on the organization and properties of Fe3 Al-based alloys. J Metals 31B(8):368–373 16. Haijun M, Li Y, Puchkov UA, et al (2008) Microstructural characterization of welded zone for Fe3 Al/Q235 fusion-bonded joint. Mater Chem Phys (112):810–815

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Chapter 5

Joining of Iron-Aluminium Intermetallic Compounds

The unique properties of Fe-Al intermetallic compounds make them very promising, however, welding is one of the main obstacles limiting the engineering applications of Fe-Al intermetallic compounds due to the brittleness of this material. It will promote the application of Fe-Al intermetallic compounds with its excellent resistance to oxidation and corrosion, as well as creep resistance to obtain a solid welded joints. The current welding methods used for Fe-Al intermetallic compounds are mainly fusion welding (such as electron beam welding, tungsten arc welding, electrode arc welding), solid phase welding (such as diffusion welding, friction welding) and brazing.

5.1 Iron-Aluminium Intermetallic Compounds and Its Weldability 5.1.1 Characteristics of Iron-Aluminium Intermetallic Compounds The commonly used Fe-Al intermetallic compounds are mainly Fe3 Al-based intermetallic compounds. The mechanical properties of Fe3 Al are mainly affected by the Al content. The room temperature mechanical properties of DO3 structure Fe3 Al with 23–29% atomic percentage of Al are shown in Fig. 5.1. The periodic fatigue properties of Fe-23.7Al and Fe-28.7Al are shown in Fig. 5.2. The yield strength of Fe3 Al is highest with 24–26% atomic percentage of Al (750 MPa) and then decreases rapidly to 350 MPa when the Al content is as high as 30% (atomic fraction). At 24–26% atomic percentage of Al, the Fe3 Al alloys have high yield strength due to age strengthening by precipitation of the disordered α phase from the ordered D03 phase. Higher Al content alloys have no age strengthening © Chemical Industry Press 2023 Y. Li, Joining Technology and Application of Advanced Materials, Advanced and Intelligent Manufacturing in China, https://doi.org/10.1007/978-981-19-9689-4_5

227

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5 Joining of Iron-Aluminium Intermetallic Compounds

Fig. 5.1 Effect of different aluminium content on yield strength and elongation of Fe3 Al alloy

Fig. 5.2 Comparison of the fatigue strength of Fe-23.7Al and Fe-28.7Al at 25 and 500 °C

because its composition at 500 °C is outside the α + D03 phase region. While the elongation of Fe3 Al alloys increases with the increasing of Al content, as seen in Fig. 5.1, the elongation of Fe3 Al increases from 1 to 5% when the atomic percentage of Al increases from 23 to 29%. The effect of different heat treatment temperature (500–1100 °C) on the room temperature mechanical properties of Fe3 Al alloys is shown in Fig. 5.3. The room temperature plasticity is significantly improved with stress relieving annealing at 700–750 °C, that is to suppresses the environmental hydrogen embrittlement to some extent. Figure 5.3 also shows that the plasticity and strength of the Fe3 Al alloy decreases continuously with the increasing of the heat treatment temperature, and the

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229

Fig. 5.3 Effect of heat treatment temperature on mechanical properties of Fe3 Al alloy (chemical composition: Fe-28Al -5Cr-0.1Zr-0.05B)

lowest plasticity and strength is obtained as the formation of completely recrystallized tissue when the annealing temperature above 1000 °C. Fe-23.7Al has a longer fatigue life than Fe-28.7Al due to different dislocation types with the same stress at room temperature, while Fe-23.7Al has better fatigue performance than Fe-28.7Al due to the second phase strengthening at 500 °C. The yield strength of metals and alloys usually decreases with increasing temperature, but the yield strength of Fe3 Al increases with increasing temperature above 300 °C, and reach the peaks at about 550 °C, and then decreases sharply with increasing temperature. This anomalous temperature relationship for the yield strength of Fe3 Al occurs in Fe3 Al alloys with 23–32% atomic percentage of Al. The elements that improve the room temperature plasticity of Fe3 Al are Cr and Nb. The room temperature yield strength of Fe-28Al alloys with Cr mass fraction of 2–6% decreases from 279 MPa to about 230 MPa, while the elongation increases from 4 to 8–10%; the yield strength increases slightly at 600 °C and the plasticity improves slightly. The corresponding fracture type changed from transgranular cleavage fracture to mixed crystal fracture. The solubility of Nb in Fe3 Al is low, only 2% (mass fraction) at 1300 °C; with decreasing temperature, the solubility decreases rapidly to 0.5% (mass fraction) at 700 °C. The Fe-25Al-2Nb alloy is quenched at 1300 °C and then age-treated at 700 °C for 8 h and air-cooled to obtain the L21 structured eutectic phase. By extending the aging time, the Fe2 Nb phase with C14 structure of solid solution Al was obtained. From room temperature to 600 °C, precipitation strengthening increased the yield strength by 50%. The thermal stability is significantly improved by added with 2% (mass fraction) Ti. B is effective for grain refinement of Fe3 Al, other elements such as Ce, S, Si, Zr and rare earth elements also have a refining effect, and Mo elements have

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5 Joining of Iron-Aluminium Intermetallic Compounds

a hindering effect on grain growth at high temperatures. The addition of 0.5% (mass fraction) TiB2 can control the grain size and improve the mechanical properties. Si, Ta and Mo can also significantly increase the yield strength of Fe3 Al, but will make Fe3 Al much less plastic. FeAl alloys have higher modulus of elasticity, higher melting point and higher specific strength. FeAl alloys with low Al content have severe environmental brittleness, while FeAl alloys with higher Al content exhibit extremely low plasticity and toughness under all test conditions due to the nature of weak grain boundary. It is difficult to increase its plasticity even by refining the grain size. FeAl mechanical properties are greatly influenced by the alloying elements, and the mechanical properties of FeAl alloys containing different alloying elements are shown in Fig. 5.4. FeAl yield strength and plasticity have a certain relationship with temperature. The strength of Fe-40Al alloy can be maintained above 270 MPa when it is raised from room temperature to 650 °C. The strength decreases rapidly when the temperature is higher than 650 °C, and the elongation increases from 8% at room temperature to more than 40% at 868 °C. The fracture form of FeAl alloy is intergranular fracture at room temperature, and transgranular cleavage fracture at high temperature. The yield strength of powder metallurgical pressed Fe-35Al and Fe-40Al alloys slowly decreased from room temperature to 600 °C with Fe-40Al alloy decreasing from 650 to 400 MPa and Fe-35Al alloy decreasing from 500 to 400 MPa, while the elongation increased from 7% at room temperature to 25% at 500 °C, but there was a decrease in plasticity while changing to intergranular fracture at 600 °C. The addition of Cr, Mn, Co, and Ti to B2 structurally ordered FeAl alloys can induce solid solution strengthening of FeAl alloys, while Nb, Ta, Hf, and Zr tend to form second-phase strengthening. Moreover, oxyphile elements such as Y, Hf, Ce, and La can restrain cavity formation to improve the denseness of FeAl alloys. The strengthening effect of Hf is larger, and the yield strength is maintained at 800 MPa, the room temperature plasticity is slightly reduced and the high temperature plasticity

Fig. 5.4 Effect of alloying elements on the mechanical properties of FeAl

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is greatly increased in the range of 27–427 °C. And the elongation of FeAl alloys is up to 50% at 827 °C. Appropriate hot processing processes (including forging, extrusion, hot rolling, warm rolling, etc.) can also improve the performance of Fe-Al intermetallic compounds. The use of intermediate processing technology of forging and extrusion before hot rolling and temperature-controlled rolling can refine grains to improve the processing properties of subsequent rolling processes. Hot rolling above recrystallization temperature leads to further grain refinement of Fe3 Al intermetallic compounds, and warm rolling below recrystallization temperature allows the grains to become bar-like morphology, which is conducive to reducing the diffusion channels of hydrogen atoms and improving the room temperature plasticity of Fe3 Al. The mechanical properties of Fe3 Al obtained by different thermal processing and heat treatment processes are shown in Table 5.1. The heat treatment process has a significant effect on the mechanical properties of Fe3 Al. The heat treatment process of repeated controlled temperature rolling followed by annealing under below recrystallization temperature conditions followed by quenching resulted in a significant improvement in the mechanical properties of Fe3 Al, with yield strength reaching about 700 MPa and room temperature elongation increasing from 2–3 to 12%. Mechanical alloying is a new process for the production of Fe3 Al, which involves ball milling in a high energy ball mill to form finely organized alloys that are alloyed in Table 5.1 Mechanical properties of Fe3 Al obtained by different thermal processing and heat treatment Alloys

Process

Heating treatment

Tensile strength/MPa

Yield strength /MPa

Fe3 Al (5.1%Cr, 0.01%Zr, 0.05%B)

Forging, rolling

Anealing above recrystallization temperature

461

260

6.3

Forging, rolling

Anealing below recrystallization temperature

590

310

10.1

Forging, rolling

Anealing below recrystallization temperature

639

340

12.3

Casting rolling

Anealing below recrystallization temperature

671

380

7.1

Forging, rolling

Anealing below recrystallization temperature火

690

420

12.5

Forging, rolling

Anealing below recrystallization temperature

705

470

10.3

Fe3 Al (4.5%Cr, 0.05%Zr)

Fe3 Al (2.35%Cr, 0.01%Ce)

Elongation/%

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5 Joining of Iron-Aluminium Intermetallic Compounds

the solid phase state. The Fe3 Al-based alloys synthesized using mechanical alloying technique achieved a tensile strength of 690 MPa and room temperature elongation of 10%.

5.1.2 Weldability Characteristics of Iron-Aluminium Intermetallic Compounds When Fe3 Al are welded by fusion welding methods (such as tungsten argon arc welding, arc welding, etc.), the rapid solidification and cooling of the weld causes large stresses, and the weld is sensitive to the alloy composition and process parameters to induce cracks. The addition of Zr and B elements to Fe3 Al is difficult to prevent weld cold cracking, although it can refine the organization of the Fe3 Al base material. The plate thickness of 0.5 mm can avoid weld cracking with Fe3 Al-based alloy wire containing 5.45% Cr, 0.97% Nb, and 0.05% C under strict control of welding speed and heat input. For Fe3 Al plates thicker than 1 mm, strict control of heat input or preheating before welding and slow cooling after welding is required to avoid delayed cracking. Preheating temperature is usually 300–350 °C, and heat treatment 600–700 °C × 1 h after welding. Fe3 AlCr alloy, low and medium carbon CrMo steel, Cr25Ni13 stainless steel and Ni-based alloys can be used as filler materials for tungsten argon arc welding (GTAW) of Fe3 Al homogeneous and heterogeneous materials. When low and medium carbon CrMo steel wire are used as filler material, Fe3 Al showed better weldability as the weld composition varied continuously to be in favour of stable properties. Although the Ni-based alloy itself has high toughness, the tendency of cracking in the Fe3 Al joint area after welding is still serious, which is due to the large coefficient of thermal expansion of the Ni-based wire resulting in larger stresses during solidification. In addition, the addition of Ni makes the composition, organization and phase structure of the fusion zone complex, the melting pool metal can not connect the semi-melting grains of the base material to form epitaxial crystallization, and the formation of tissue separation zone at the fusion junction occurred. Controlling the welding current and heat input is conducive to improving the crack resistance of Fe3 Al when using the same base material with dissimilar wire, under the condition of ensuring weld penetration. The Fe3 Al alloys were rolled into sheets by a hot-rolling-controlled temperature rolling process after vacuum melting into ingots, with a vacuum of 1.33 × 10–2 Pa during the melting process. The main chemical composition of the Fe3 Albased alloys used for the tests was: Al 16.0–17.0%, Cr 2.40–2.55%, Nb 0.95–0.98%, Zr 0.05–0.15%, Fe 81.0–82.5%. The main problem of welding for Fe3 Al-based alloy is microcracking due to the brittleness and poor weldability. The filler alloy is required to contain alloying elements that can improve the plasticity and toughness of Fe3 Al, and improve the crack resistance of the fusion zone during welding to avoid the generation of weld cracks. Cr is the most effective alloying element to improve the plasticity of Fe3 Al,

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and Ni is a commonly used element to improve alloy toughness. Therefore, the FeCr-Ni alloy system can be used as a filler material for Fe3 Al welding, with a filler wire diameter of 2.5–3.0 mm. The excellent high-temperature performance of Fe3 Al intermetallic makes it a potential structure material. However, Fe3 Al is difficult to fusion welded due to the brittleness of this material, which is the main obstacle limiting their engineering applications. The performance of the joint is determined by the microstructure of the weld zone. The relationship between the microstructure and the performance of the welding joint can be established by analyzing the microstructure and alloying element distribution in the Fe3 Al TIG joint zone, which can provide an experimental basis for the optimum welding parameters.

5.1.3 Cracking in the Fe3 Al Welded Joint Area (1) Origin and expansion of cracks Fe3 Al joints welding cracks are easily induced when the welding heat input is too little or too much. The longitudinal cracking usually occurred at the ends of the welded joint and transverse cracking inside the joint, which is related to the distribution of weld stresses. Cracks in Fe3 Al/18–8 steel welded joints all originate in the partial fusion zone of the Fe3 Al side, which is mainly caused by: (a) The brittle phases is easily formed in the partial fusion zone where the most complex interactions of alloying elements such as Al, Fe, Cr and Ni occur. (b) The diffusion of atomic hydrogen into defects such as anti-phase domain boundaries induce possible combination into hydrogen molecules, which increases the microstress at the defects and provides conditions for the origin of cracks. (c) The partial fusion zone is the area of maximum welding stress, which is conducive to crack initiation and expansion. The cracks originated in the partial fusion zone can expand both longitudinally along the partial fusion zone and laterally into the Fe3 Al side heat affected zone. It is difficult for cracks to expand through inhomogeneous mixing zone as the organization of the γ + δ in this region, and a small amount of cracks that expand in the direction of the weld are controlled at the inhomogeneous mixing zone. When the welding heat input is small, the cooling rate of the weld is faster, resulting in higher stresses in the joint, cracks in the Fe3 Al side heat-affected zone can extend over a longer distance, and the number of cracks and the extension distance are greater than in the case of larger welding heat input; even new sources of cracks are formed in the heat-affected zone, and the extension direction is more chaotic, resulting in the Fe3 Al heat-affected zone becoming a weak region of the joint. When the welding heat input is larger, the joint cooling rate is slower, which is conducive to the transition of Cr elements into Fe3 Al and improves the plasticity and toughness of the Fe3 Al fusion zone. In addition, Al, Fe, Cr, Ni and other alloying

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5 Joining of Iron-Aluminium Intermetallic Compounds

elements can be fully fused and diffused, reducing the generation of elemental segregation and brittle phases. Meanwhile, the high temperature diffusion time of Al elements grows, the A2 disorder structure in the Fe3 Al heat affected zone increases, reducing its brittleness and hardness, which is beneficial to stop the crack expansion. At welding heat input E = 10.8 kJ/cm, the crack extension distance of Fe3 Al/18–8 joints is small, and the crack length is between 50 and 150 μm. (2) Influencing factors on the cracks The weldability of Fe3 Al intermetallic compounds is poor, mainly in the following two aspects. Firstly, Fe3 Al intermetallic compounds have high stress concentrations due to difficulties in cross-slip migration, resulting in high room temperature brittleness, low plasticity and susceptibility to cold cracking during welding. Secondly, the low thermal conductivity of Fe3 Al leads to a large temperature gradient among the heat affected zone, the fusion zone and the weld seam, coupled with a large coefficient of linear expansion, which tends to produce large residual stresses during cooling, leading to thermal cracking. Fe3 Al weld cracks originate in the partial fusion zone of the Fe3 Al side fusion zone and extend in the partially melted zone and the heat affected zone, with only a small amount of cracks extending into the weld. Fe3 Al cracks are mainly caused by its brittle nature, the brittle phases in the fusion zone and the welding stresses, which include the following: (1) Grain state of Fe3 Al base material. The finer the grain size of Fe3 Al, the better it is to prevent cracking. (2) Welding heat input. Too small or large welding heat input can easily lead to weld cracks. (3) The degree of segregation of alloying elements and the number of brittle phases in the partial fusion zone. (4) Microstructure of Fe3 Al heat affected zone. The more A2 disordered structure and B2 partially ordered structure in the Fe3 Al heat affected zone, the better it is to prevent the crack generation and extension. (5) The amount of diffused hydrogen in the welding joint. The lower the content of diffused hydrogen in the joint, the better its resistance to cracking. The following measures can prevent or reduce the development of Fe3 Al weld cracks. (1) (2) (3) (4)

Fine-grained Fe3 Al base material was used. Increase the welding heat input appropriately. Use of suitable filling materials. Enhance gas protection for the welding process.

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5.2 Wire-Filled Tungsten Arc Welding of Fe3 Al and Steel (Q235, 18–8 Steel)

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5.2 Wire-Filled Tungsten Arc Welding of Fe3 Al and Steel (Q235, 18–8 Steel) 5.2.1 Characteristics of the Tungsten Arc Welding Process of Fe3 Al and Steel (1) Welding method In the tungsten arc welding (GTAW) of Fe3 Al with steel, the grain size of the heat affected zone is very coarse affected by the welding thermal cycle, and its high temperature oxidation resistance is also slightly lower than that of the Fe3 Al base material due to the burnout of Al elements during the welding process. The tensile strength of the welding joint area is lower than that of the base material and breaks in the overheated zone of heat-affected zone. The overheated zone undergoes a welding heating and subsequent cooling process during welding and the originally higher degree of ordering is significantly reduced. The orderliness of the overheated zone is difficult to recover even after post-weld heat treatment. Therefore, the strength and hardness of the overheated zone is reduced and becomes the weak point of the joint. The tensile strength and elongation of the overheated zone in the heat affected zone are reduced compared to the Fe3 Al base material. The test base materials were Fe3 Al intermetallic compound, Q235 steel and 1Cr18Ni9Ti austenitic stainless steel (18-8 steel). Fe3 Al intermetallic compound was vacuum melted into ingots and then rolled into sheets using a hot rolling-controlled temperature process and annealed at 1000 °C for homogenization. The raw material Fe was de-rusted by rolling before melting, and the raw material Al was cleaned with NaOH solution and dried to obtain Fe3 Al intermetallic compound with dense organization and good properties. Fe3 Al, Q235 steel and 18–8 steel were cut into plates of 8, 5 and 2.5 mm thickness, respectively, using wire cutting and mechanical machining methods. The chemical composition and thermophysical properties of the Fe3 Al intermetallic compound used for the tests are shown in Table 5.2. The Fe3 Al intermetallic compound contains alloying elements of Cr, Nb, and Zr. The microstructure consists of coarse massive grains with Cr- and Nb-rich secondphase particles distributed inside the grains and at the boundaries, as shown in Fig. 5.5a. These second-phase particles hinder the movement of dislocations along the grain boundaries, increase the rate of compressive deformation of Fe3 Al, and improve its strength, plasticity and toughness. The microstructure of the 18–8 steel is γ -austenite +a small amount of δ-ferrite, as shown in Fig. 5.5b. Fe3 Al and Q235 steel (or 18-8 steel) was welded via wire-filled tungsten arc welding (GTAW) without preheating and post-weld heat treatment. Tungsten tungsten arc welding (GTAW) was performed with a ZX69-150 AC-DC silicon rectifier argon arc welder.

16.82

81.02

Ordered critical temperature/°C

480–570

Structure

DO3

Thermo-physical properties

Al

Fe

Chemical compositions/wt.%

140

Yang’s modulus/GPa

0.78

Cr

1540

Melting point/°C

0.63

Nb

11.5

Coefficient of heat expansion/10–6 ·K−1

0.28

Zr

6720

Density /kg·m−3

0.01

B

Table 5.2 Chemical composition and thermophysical properties of Fe3 Al intermetallic compounds

455

Tensile strength/MPa

0.1

Mn

3

Elongation /%

≥29

Hardness HRC

0.15

Ce

236 5 Joining of Iron-Aluminium Intermetallic Compounds

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5.2 Wire-Filled Tungsten Arc Welding of Fe3 Al and Steel (Q235, 18–8 Steel)

237

Fig. 5.5 Microstructure of Fe3 Al intermetallic compound and 18-8 steel

(2) Selection of welding materials Fe-Cr-Ni alloy system was used as filler material to study the welding behavior of Fe3 Al. GTAW filler alloys were Cr19-Ni10, Cr18-Ni12Mo2, Cr23-Ni13 and Cr26Ni21 with a filler wire diameter of 2.5 mm. The surfaces of the specimens to be welded (Fe3 Al intermetallic compound, Q235 steel and 18-8 stainless steel) are machined before welding to ensure that the upper and lower surfaces are parallel and a surface finish of grade 5. Oxide film, oil and rust of the specimen plate and filler material are removed by chemical means. The treatment steps are: sandpaper grinding → acetone cleaning → water rinsing → alcohol cleaning → blowing dry. (3) Technological parameters The butt welding tests of series Fe3 Al/Fe3 Al, Fe3 Al/Q235 steel and Fe3 Al/18-8 steel with no-preheating was carried out via filler tungsten arc welding (GTAW). The parameters used are shown in Table 5.3. The tests show that welding cracks were easily occurred when the heat input too much or too little, and the effect of welding heat input on the susceptibility of Fe3 Al joints to cracking exceeds that of the filler material. When the welding heat input is too small, the weld cools quickly and produces significant surface cracks. In tungsten arc welding (GTAW), under the action of flowing argon gas, the cooling rate of the weld is faster than that of electrode arc welding (SMAW), so the cracking tendency is more serious. Excessive heat input and long overheating time of the molten pool leads to coarsening of the weld tissue Table 5.3 Process parameters used for Fe3 Al filled wire GTAW Method

GTAW

Process parameters Current I /A

Voltage U /V

Weling speed v /cm·s−1

Ar flow rate L/min

Heat input E /kJ·cm−1

100–120

11–12

0.15–0.25

8–12

4.5–9.6

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5 Joining of Iron-Aluminium Intermetallic Compounds 90mm Fe3Al

Fig. 5.6 Schematic diagram of Fe3 Al butt joint specimen

Filler

50mm 2.5mm

Welding direction 18-8

to induce cracking. Fe3 Al joints are prone to cracking regardless of which kind of filler material used. Tests have shown that a crack-free Fe3 Al joint can be obtained by tungsten arc welding (GTAW) with Cr23-Ni13 filler alloy using a suitable welding heat input. Cr content has an important influence on the crack sensitivity of Fe3 Al. The Cr content of the welding material is 23–26%, which ensures an appropriate amount of Cr transition into the Fe3 Al fusion zone to improve the crack resistance of the joint. Welding for Fe3 Al-based alloy is advisable to use a low-current, low-speed welding process influenced by the thermophysical properties and weld formation of Fe3 Al. Depending on the plate thickness, the welding heat input should be controlled. The heat input can be adjusted appropriately for filler tungsten arc welding (GTAW), because the cooling rate of the weld is faster than that of electrode arc welding (SMAW) under the action of flowing argon gas. (4) Preparation of Fe3 Al butt joint specimens The series of Fe3 Al/Fe3 Al, Fe3 Al/Q235 steel and Fe3 Al/18–8 steel joint specimens were cut by wire cutting method. The schematic diagram of Fe3 Al butt joint specimens is shown in Fig. 5.6. The specimens were sandpapered, mechanically polished and etched, using Cr2 O3 aqueous solution as polishing agent. For Fe3 Al/Q235 joints, due to the large difference in corrosion resistance between Fe3 Al and Q235 steel, the specimens were etched with 3% nitric acid alcohol solution on the Q235 steel side first, and then sealed with paraffin; then the Fe3 Al side was etched with aqua regia solution (HNO3 :HCl = 1:3), and finally the paraffin wax on the Q235 steel side was polished and removed. As the heat-affected zone on the Fe3 Al side was difficult to corrode, a mixture of aqua regia and hydrochloric acid + acetic acid + nitric acid (HCl:HNO3 :CH3 COOH = 1:3:4) was used. The Fe3 Al/18–8 joint was corroded directly with the aqua regia solution. Thin-film specimens for TEM analysis were cut from the welds, Fe3 Al side fusion zone and Fe3 Al heat-affected zone of Fe3 Al/Fe3 Al, Fe3 Al/Q235 and Fe3 Al/18-8 joints, the, respectively. The thin film specimens were then mechanically ground to a thickness of about 50 μm and then chemically and electrolytically double-sprayed to form thin film specimens suitable for TEM testing. A series of thin film specimens were then subjected to TEM and selected area electron diffraction analysis.

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5.2.2 Microstructure Characteristics of the Fe3 Al/Steel Joint Zone of Filled Wire GTAW (1) Crystallization process of wire-filled GTAW weld Factors affecting the interaction between elements are related to the solid solution degree of the elements, including infinite solid solution and limited solid solution. The weldability between metals with infinite solid solution is better than that of limited solid solution. GTAW for Fe3 Al mainly involves the interaction of Fe, Al, Cr, and Ni elements. And the interaction characteristics of the four elements of Fe, Al, Cr, and Ni are shown in Table 5.4. Solid solution or compounds can be generated due to the diffusion and interaction of Fe, Al, Cr and Ni elements in the process of solidification and crystallization of the molten pool. Instantaneous local melting of Fe3 Al and 18-8 steel occurred and mixed molten Cr23-Ni13 filler metal to form molten pool during Fe3 Al/18-8 steel filler wire GTAW. And Fe3 Al can get into the molten pool in the way of mechanical mixing, metal mass and diffusion mixing. A small amount of Al entering the melt pool can improve the oxidation and corrosion resistance of Cr-Ni welds.Al is ferritisey that limits γ austenite formation and promotes δ ferrite formation, which will expand the δ and δ + γ phase zones and reduce the γ phase zone. More Ni is required to form austenite in Cr-Ni welds containing higher Al. Al also increases the activity of carbon in the steel and promotes carbide precipitation. The Cr and Ni content has an important influence on the weld microstructure. The ratio of Cr and Ni in the Fe3 Al/18-8 steel is about 2:1 when Cr23-Ni13 alloy is used as the filler metal. The δ ferrite is first precipitated from the liquid phase L when the molten pool solidifies. As the temperature decreases, δ phase continued to precipitate from the (L + δ + γ) three-phase region, and the γ phase also begins to precipitate. When the molten pool solidifies completely, the δ phase transforms Table 5.4 Interaction characteristics of Fe, Al, Cr and Ni elements Element Melting Crystal Lattice Atomic Solid solution Compounds point transition type radius/nm Unlimited Limited temperature/°C /°C 910

α-Fe BCC γ-Fe FCC

0.1241

α-Cr, γ-Ni Al, γ-Cr, α-Ni

Cr, Ni, Al

660

–

FCC

0.1431

–

Ni, Cr, Fe

Cr, Fe, Ni

Cr

1875

–

BCC

0.1249

α-Fe

γ-Fe, Ni, Al

Fe, Ni, Al

Ni

1453

–

FCC

0.1245

γ-Fe

Cr, Al, α-Fe

Cr, Fe, Al

Fe

1536

Al

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5 Joining of Iron-Aluminium Intermetallic Compounds

into the γ phase, and part of the δ phase remains in the weld. As Al is ferritisey, the range of the γ-phase zone is reduced, and the γ → α transformation occurred at a lower temperature. The α-phase combined the carbide to form ferrite, bainite and martensite and other organizations. (2) Division of GTAW joint zone The welding joint consists of three zones: the weld, the fusion zone and the heat affected zone. In order to analyze the characteristics of different zones of Fe3 Al filled wire tungsten arc welding joints, the Fe3 Al side welding zone can be divided into four characteristic zones: homogeneous mixture zone (HMZ), partial mixture zone (PMZ), partially fused zone (PFZ) and heat-affected zone (HAZ). The weld is made up of HMZ and PMZ. PFZ and PMZ adjacent to PFZ are collectively referred to as the fusion zone. PMZ is the transition region between the weld and the fusion zone, which has the most complex microstructure in the Fe3 Al joint. Under the effect of convection and stirring of the molten pool, a small amount of unmelted base material particles enter the molten pool and are melted and mixed homogeneously to form a uniform mixing zone in the upper part of the Fe3 Al joint after solidification. The morphology of HMZ of the Fe3 Al filled wire GTAW joint is shown in Fig. 5.7. The organization of HMZ consists of columnar crystals and equiaxed crystals. In the middle of HMZ of Fe3 Al joints, equiaxed austenite grains are formed as small temperature gradient, and pro-eutectoid ferrite grows from the austenite grain boundaries towards inside, as shown in Fig. 5.7a. At the bottom of the HMZ, austenitic columnar crystals grow along the maximum temperature gradient, pro-eutectoid ferrite grows parallel along the austenite grain boundaries, and eutectic structure exists inside the austenite grains, as shown in Fig. 5.7b. The action of convection and scouring is different at the bottom of the molten pool, a small amount of incompletely melted Fe3 Al, and some completely melted but not disperse Fe3 Al form the partial mixture zone (PMZ) of the joint when solidification.

(b)

(a)

γ

γ

PF PF 100μm

100μm Equiaxed grains in HMZ

Columnar grains in HMZ

Fig. 5.7 Morphology of the HMZ of the Fe3 Al filled wire GTAW joint

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5.2 Wire-Filled Tungsten Arc Welding of Fe3 Al and Steel (Q235, 18–8 Steel)

(a)

(b)

AF

γ

241

Carbide

γ UB PF PF LM 100μm

25μm

Bu

OM, 100

SEM, 400

Fig. 5.8 Microstructure of the homogeneous mixture zone of the wire-filled GTAW Fe3 Al/18-8 steel welds

And different morphology appear under the combined action of alloying elements in partial mixture zone (PMZ). Grains of Fe3 Al is coarse, under the action of welding heat, Fe3 Al melts from the grain boundaries and react with with the weld metal to forms a partially fused zone (PFZ). A sandwich structure is prone to form in partially fused zone (PFZ) due to the large difference in composition of Fe3 Al and Cr23-Ni13 filler material. Fe3 Al heat-affected zone is formed after welding, in which the grain size change little with suitable heat input. (3) Microstructure characteristics of Fe3 Al/18-8 joint The microstructure of the Fe3 Al/18-8 steel filled wire tungsten arc weld is shown in Fig. 5.8. The filled wire was austenitic steel selected from Cr25-Ni13 series, and the weld organization consisted mainly of austenite and a small amount of lath martensite, with a little ferrite and ferrite side plate at the austenite grain boundaries. The homogeneous mixture zone of Fe3 Al/18-8 steel joints is based on massive γphase with lamellar pro-eutectoid ferrite (PF) precipitation at the γ-grain boundaries to form a network of pro-eutectoid ferrite. The upper bainite (Bu ) is nucleated at the grain boundaries and grows parallel to the inside. There distributed small acicular ferrite (AF) and lath martensite (LM) within γ crystals. The strength and toughness can be guaranteed by this γ + α organization, which enhances the crack resistance of the joint. The Al element drives the bainite transformation, so bainite is easily formed in alloys containing Al element, thus more bainite organization is found in Fe3 Al/18-8 welds. Bainite is a mechanical mixture of α-Fe and carbide, and its morphology is related to the formation temperature. The morphology of upper bainite in the Fe3 Al/18-8 homogeneous mixture zone is shown in Fig. 5.9a; The austenite between the ferrite laths is carbon rich and stabilized due to the role of Al in retarding the precipitation of cementite. Therefore, upper bainite with residual austenite between lath ferrite is formed, as shown in Fig. 5.9b.

242

5 Joining of Iron-Aluminium Intermetallic Compounds

(a)

(b)

Bu PF γ

Bu 10μm

100μm Upper bainite by OM

Upper bainite (SEM)

(c)

(d) γ Bl Bl

PF

10μm

100μm Lower bainite by OM

Lower bainite by SEM

Fig. 5.9 Bainite in the homogeneous mixture zone of Fe3 Al/18-8 steel joints

Lower bainite mainly distributed inside the ferrite lath. The carbides in lower bainite can be either cementite or ε-carbide. lower bainite in the Fe3 Al/18-8 homogeneous mixture zone is black needle-like or lamellar with a certain angle of intersection under light microscopy, as shown in Fig. 5.9c; There distributed lamellar or granular carbides in the lower bainite, arranged at an angle of 55–60° with the long axis of ferrite when observed under SEM. There is no obvious carbide precipitation in the lower bainite ferrite due to the influence of Al elements, as shown in Fig. 5.9d. With the decreasing of bainite formation temperature, the carbon content of ferrite in bainite gradually increases. The microstructure of partial mixture zone of Fe3 Al/18-8 steel is dominated by γ austenite and a little δ ferrite. γ austenite is mainly coarse cellular dendritic crystals. Fine γ austenite is formed in partial mixture zone near partially fused zone, where there is a large subcooling to form more nuclei when solidification. The δ-phase zone is expanded because of the high content of Al elements at this location, the primary δ ferrite can grow upto the solid-phase line without passing through the (L + δ + γ) three-phase zone and not affected by primary γ austenite in the (L + δ + γ) three-phase zone. The δ → γ transformation is suppressed with the increase of the cooling rate, leading to a significant increasing of the residual δ ferrite. The morphology of Fe3 Al/18-8 steel weld is different depending on the distance from the fusion zone. The nucleation and growth of crystals need a certain degree of

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243

subcooling, which is divided into positive temperature gradient (G > 0) and negative temperature gradient (G < 0), as shown in Fig. 5.10. Positive temperature gradient conditions tend to form equiaxed crystals and negative temperature gradient tend to form dendrites. When the weld metal solidifies, there is constitutional supercooling in the solid– liquid interface except temperature subcooling, which is the reason why the dendrites are easy to appear in the weld without big subcooling. The morphology of Fe3 Al/18-8 steel weld is different in different regions due to different degrees of subcooling. The composition in both sides of partial mixture zone (PMZ) of Fe3 Al/18-8 steel joint is in big difference, one side is rich in Al, the other side is rich in Cr, Ni, resulting in a large constitutional supercooling. The protrusion part can penetrate deeper into the liquid weld inside a longer distance, while the protrusion part also discharge solute to the surrounding inducing the constitutional supercooling, the short secondary cross branch grow from the main stem, as shown in Fig. 5.11. Fe3 Al side partially fused zone consists of white bright and dark laminar tissue with a clear boundary with the Fe3 Al heat affected zone. Some black phase is present

Fig. 5.10 Subcooling of liquid metal during crystallization. TM —metal freezing point; ΔT— subcooling

(a)

(b) γ

δ

γ

ML

50μm

δ

Fig. 5.11 Cellular dendrites in the partial mixture zone of Fe3 Al/18-8 steel joints

50μm

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5 Joining of Iron-Aluminium Intermetallic Compounds

(b)

(a) γ

PFZ HAZ

MZ

δ

γ-rich band

PFZ PMZ 50μm

10μm

Fig. 5.12 Morphology of the Fe3 Al side fusion zone of Fe3 Al/18-8 steel joint

in the laminar tissue near partial mixture zone, which is caused probably by the oxidation of alloying elements during welding, as shown in Fig. 5.12a. This laminated structure adversely affects the properties of the joint, but no cracks were found. There is a about 30 μm width “austenite rich band” with relatively few δ ferrite along partially fused zone. Austenite is arranged parallelly in slate-like shape, slate width is about 5 μm, and make a 50–70° angle with partially fused zone, as shown in Fig. 5.12b. The existence of “austenite rich band” has two roles: first, to reduce the hazards of the brittle phase to the partial mixture zone; second, to restrict diffusion of the hydrogen in the weld to partially fused zone and heat-affected zone due to the greater solubility of hydrogen in the austenite, to prevent the generation of hydrogen cracking. (4) Microstructure characteristics of Fe3 Al/Q235 steel joint The microstructure of homogeneous mixture zone of Fe3 Al/Q235 steel filled wire GTAW joints are based on γ austenite, with lamellar pro-eutectoid ferrite (PF) precipitation at the austenite grain boundaries, forming a pro-eutectoid ferrite network. Upper bainite (Bu ) is nucleated at γ grain boundaries and grows parallel to the grain inside. Compared to Fe3 Al/18-8 steel joints, the amount of pro-eutectoid ferrite (PF) and upper bainite (Bu ) is reduced, as shown in Fig. 5.13. In the partial mixture zone of the Fe3 Al/Q235 steel joint, γ columnar crystals are formed along the direction of the maximum temperature gradient due to the faster cooling rate, the growth direction of the columnar crystal is basically perpendicular to the fusion line, and inside the austenite grains, more upper bainite (Bu ) precipitates along the direction parallel to the fusion zone, as shown in Fig. 5.14a, which may be caused by the segregation of Al elements. A large number of cellular crystals appear in the partial mixture zone near partial fusion zone. Under subcooling conditions, the crystalline surface in an unstable state, many parallel bundles of buds grow into the subcooled liquid weld from the solidification interface to cellular crystals as shown in Fig. 5.14b. These cellular austenite constitutes an “austenite-rich zone”, and the width of the austenite lath significantly reduced to about 2 μm compared with Fe3 Al/18-8 steel fusion zone, and basically perpendicular to the partial fusion zone.

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5.2 Wire-Filled Tungsten Arc Welding of Fe3 Al and Steel (Q235, 18–8 Steel) (a)

245

(b) γ PF

γ

Bu

PF 100μm

100μm OM

SEM

Fig. 5.13 Morphology of the homogeneous mixture zone of Fe3 Al/Q235 steel joints

(b)

(a)

PFZ γ

Bu

γ

PMZ

PF Columnar dendrites

Cellular dendrites

Fig. 5.14 Morphology of the partial mixture zone in Fe3 Al/Q235 steel joints

The comparison of the alloying element contents in the homogeneous mixture zone and partial mixture zone of Fe3 Al/Q235 steel joints is shown in Table 5.5. Compared with the homogeneous mixture zone, the content of Cr and Ni in the partial mixture zone decreases significantly, while the content of Al elements doubles. The ratio of Cr to Ni content is about 3:1. Primary δ ferrite first precipitates in partial mixture zone during solidification, transformation of δ ferrite to γ austenite and γ austenite to α ferrite further occur during cooling, and finally γ + α mixed organization is formed. Table 5.5 Comparison of the composition of HMZ and PMZ of Fe3 Al/Q235 joint Location

Composition wt.% Al

Si

Cr

Mn

Fe

Ni

(HMZ)

1.65

0.43

16.77

1.22

70.01

9.92

(PMZ)

3.66

0.77

13.00

1.57

75.23

3.77

246

5 Joining of Iron-Aluminium Intermetallic Compounds

There is no obvious interface between partially fused zone and heat affected zone of Fe3 Al/Q235 steel GTAW joint, and no laminar structure exists. The organization of the partial mixture zone near the partially fused zone is dominated by γ + α, and the austenite and ferrite show different morphological characteristics depending on the local heat dissipation conditions and the distribution of alloying elements. In the region with relatively slow cooling rate, the γ → α transition is more adequate, and the γ phase content is relatively small and mostly worm-like. (5) Microstructure characteristics of Fe3 Al/Fe3 Al joint When Fe3 Al is butt welded to Fe3 Al, the base material has a large dilution effect on the weld alloying elements, with an average content of 7.2% Al(mass fraction), which is about half of that in the Fe3 Al base material. According to the Scheffler organization chart, the weld has 25.43% chromium equivalent and 6.85% nickel equivalent, and the weld is based on α-Fe(Al) solid solution. The morphology of the homogeneous mixture zone has some of the “genetic” characteristics of the Fe3 Al parent material, i.e., the entire homogeneous mixture zone is organized as a matrix of coarse α-Fe(Al) with grain sizes similar to those of the Fe3 Al base material. These coarse α-Fe(Al) are composed of many smaller massive subgrains, as shown in Fig. 5.15a. The Fe3 Al parent material has coarse grains, and the morphology is also characterized by coarsening in the partial fusion zone near Fe3 Al side of Fe3 Al/18-8 steel and Fe3 Al/Q235 steel joints, such as coarse austenitic cellular dendrites in the partial fusion zone of Fe3 Al/18-8 joints, and coarse columnar crystals in the partial fusion zone of Fe3 Al/Q235 joints. For Fe3 Al/Fe3 Al welding joints, the partial fusion zone is coarsened by the influence of co-generation crystallization, and the organization of the homogeneous mixed zone is also coarsened and consists of small fragmented subgrain, which can be regarded as a continuation of co-generation crystallization in the weld. A small amount of equiaxed dendrites are also present in localized areas of the homogeneous mixture zone, as shown in Fig. 5.15b. This is associated with the

(b)

(a)

2

1

50μm OM

SEM

Fig. 5.15 Characteristics of austenite in the HMZ of Fe3 Al/Fe3 Al joints

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5.2 Wire-Filled Tungsten Arc Welding of Fe3 Al and Steel (Q235, 18–8 Steel)

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Table 5.6 Composition of medium axis dendrites in the homogeneous mixture zone of Fe3 Al/Fe3 Al joint Composition wt.%

Location

Al

Cr

Mn

Fe

Ni

Point 1

1.55

20.61

1.48

62.82

13.23

Point 2

1.84

20.67

1.53

63.21

12.53

Average

1.70

20.64

1.51

63.02

12.88

Weld

7.83

12.93

0.80

71.49

6.45

segregation of the alloying elements, causing a larger composition of subcooling in this region, resulting in a secondary transverse dendrite protruding in all directions in addition to producing a very long main stem in the weld. To determine the type of element in which the segregation occurred, the composition of the dendrite was determined and the location and results are shown in Fig. 5.15b and Table 5.6. Compared with the average composition of the weld, the Cr and Ni content in the dendrites increased significantly, by 45 and 102%, respectively; the Al content was less than 1/4 of the weld matrix. The formation of equiaxed dendrites was caused by the segregation of Cr and Ni elements, and these equiaxed dendrites were determined to be austenitic according to the composition characteristics. The morphology of the Fe3 Al/Fe3 Al fusion zone is shown in Fig. 5.16. The αFe(Al) grain morphology in this region is irregular, with second phase precipitates inside the grains. α-Fe(Al) grain boundaries extend along the partially fused zone and can even extend inside the partially fused zone, forming a semi-grain morphology due to the restriction of the Fe3 Al parent material. The width of the partial fusion zone is small and there is no obvious boundary with the heat affected zone. Since the grain boundary is the region of weaker bonding, the extension of the grain boundary is bound to affect the bond strength of the joint. Fe3 Al is the alloy with narrow crystallization temperature range, with the increase of Al content, the crystallization temperature range of the alloy widens, so that the

(b)

(a) PFZ

PFZ PMZ

PMZ

100μm PMZ

50μm PMZ and HAZ

Fig. 5.16 Morphology of the fusion zone and heat affected zone of Fe3 Al/Fe3 Al joints

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5 Joining of Iron-Aluminium Intermetallic Compounds

solid–liquid two-phase area is expanded, and the solidification process is more likely to form dendrites, so the fluidity of liquid Fe3 Al alloy becomes poor. In addition, with the increase of Al content, the tendency of oxidation to generate high melting point Al2 O3 intensifies. Al2 O3 is in solid state and is caught in the molten pool under the action of arc blowing force, etc., which deteriorates the fluidity of the molten pool. Cr reduces the fluidity of liquid Fe3 Al alloy because it increases the Fe3 Al liquid phase line temperature, which is equivalent to increasing the crystallization temperature range of Fe3 Al, but its influence is less than that of Al element. Due to the Al content greater than 28% (atom fraction) and the high Cr content of the filler material, the liquid Fe3 Al is less mobile, i.e., the surface tension of the liquid Fe3 Al is higher, and the weaker scouring force is difficult to separate the partially fused Fe3 Al from the matrix, so a melt retention layer tends to form in the Fe3 Al joint. The Cr and Ni content in the melt retention layer is significantly higher than that of the Fe3 Al base material.

5.2.3 Microhardness of the Fe3 Al/Steel Filled Wire GTAW Joint In order to determine the changes of the microstructure of the Fe3 Al welding joint, the microhardness near the fusion zone was tested using a microhardness tester with a loading load of 50 g and a loading time of 10 s. The microhardness near the fusion zone of Fe3 Al/18-8 steel and Fe3 Al/Q235 steel joints was determined and the corresponding measurement positions were given, respectively. (1) Microhardness of Fe3 Al/18-8 steel near the fusion zone The microhardness distribution of the Fe3 Al/18-8 steel GTAW joint is shown in Figs. 5.17 and 5.18. The microhardness on both sides of the fusion zone of Fe3 Al/188 steel is very different, and the microhardness of the fusion zone near 18-8 steel side is reduced compared to that of Fe3 Al side, which is due to the higher Al content near the fusion zone of Fe3 Al side, which tends to form a high hardness brittle Fe-Al phase. The microhardness near Fe3 Al side fusion zone of GTAW welding joints is higher than the heat-affected zone and weld, up to 580 HM; the microhardness of Fe3 Al heat-affected zone is 330–400 HM. Despite the high microhardness of Fe3 Al side fusion zone, there is no FeAl2 , Fe2 Al5 and other high hardness brittle phase, the Fe-Al phase generated in the welding may be Fe3 Al and FeAl mixed organization. (2) Microhardness near the Fe3 Al/Q235 fusion zone The microhardness distribution and the location of the measurements near the fusion zone of Fe3 Al/Q235 joint are shown in Figs. 5.19 and 5.20. The hardness of the fusion zone and welds on the Fe3 Al side is slightly higher than that of Q235 steel side, which is mainly influenced by the Fe-Al alloy phase.

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5.2 Wire-Filled Tungsten Arc Welding of Fe3 Al and Steel (Q235, 18–8 Steel)

249

600

(b)

(a)

FZ WM

Fe3Al

Microhardness (HM)

550

500

Fe3Al

450

WM

FZ

400

350

50μm

300

0.0

0.2

0.4

0.6

0.8

Distance /mm Microhardness distribution

Microstructure

Fig. 5.17 Microhardness of Fe3 Al/18-8 steel joints near the Fe3 Al side fusion zone 500

FZ

18-8

WM

Microhardness (HM)

(a)

(b)

FZ

450 400 350

18-8

WM

300 250 200

50μm

0.0

0.2

0.4

0.6

0.8

Distance /mm Microhardness distribution

Microstructure

Fig. 5.18 Microhardness of Fe3 Al/18-8 steel joint near the 18-8 steel side fusion zone 460

(a)

e3Al

FZ

WM

Microhardness (HM)

(b) 440

420

400

FZ

380

WM

Fe3Al

360

340

Microstructure

0.0

0.2

0.4

0.6

Distance /mm Microhardness distribution

Fig. 5.19 Microhardness of Fe3 Al/Q235 joint near the Fe3 Al side fusion zone

0.8

250

5 Joining of Iron-Aluminium Intermetallic Compounds

Fig. 5.20 Microhardness of Fe3 Al/Q235 steel joints near the Q235 side fusion zone

Compared to Fe3 Al/18-8 steel joints, the microhardness of the Fe3 Al side fusion zone is reduced, indicating that in addition to being influenced by the Fe-Al phase, alloying elements such as Cr and Ni in the weld also have an effect on the microhardness. The Cr and Ni content in the Fe3 Al/18-8 steel weld is higher than that in the Fe3 Al/Q235 weld, resulting in a higher microhardness of the organization. There is low microhardness zone about 350 HM in the Fe3 Al heat affected zone. The Al content in Fe3 Al welds is higher than that in Fe3 Al/steel joints, and more Al elements are solidly soluble in the α-Fe(Al) phase, resulting in high microhardness of the welds, with microhardness up to about 480 HM. The microhardness of the fusion zone is slightly higher than that of Fe3 Al side fusion zone of the Fe3 Al/Q235 steel joint and slightly lower than that of the Fe3 Al/18-8 steel joint. Similar to the Fe3 Al/steel joint, a low microhardness zone exists in the Fe3 Al heat affected zone with a microhardness of about 325 HM. (3) Microhardness of Fe3 Al/Fe3 Al joints The Al content in the Fe3 Al/Fe3 Al weld is higher than that in the Fe3 Al/steel joint, and more Al elements are solidly solved in the α-Fe phase, resulting in a high mcirohardness of the weld up to 480 HM, see Fig. 5.21. The mcirohardness in the Fe3 Al/Fe3 Al fusion zone is slightly higher than that in Fe3 Al side fusion zone of the Fe3 Al/Q235 steel joint and slightly lower than that of the Fe3 Al/18-8 steel joint. Similar to the Fe3 Al/steel joint, a low microhardness zone exists in the Fe3 Al heat affected zone with a microhardness of about 325 HM. (4) Softening of Fe3 Al heat affected zone and influencing factors There is a zone of low microhardness in the Fe3 Al heat affected zone, i.e. a zone of local softening. The diffusion of Al from the Fe3 Al side to the weld at high temperatures leads to a change in the organization of the Fe3 Al heat affected zone. Due to the absence of Al element, the organization of part of the heat affected zone is no longer D03 ordered structure, but disordered structure. Compared with the

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5.2 Wire-Filled Tungsten Arc Welding of Fe3 Al and Steel (Q235, 18–8 Steel) 560

(a) Weld

Fe3Al

(b)

520

Microhardness (HM)

FZ

251

0μm

Weld

480

Fe3Al

440 400 360 320 0.0

0.2

0.4

0.6

0.8

Distance (mm) Microstructure

Microhardness distribution

Fig. 5.21 Microhardness distribution near the fusion zone of Fe3 Al/Fe3 Al joints

D03 structure, the disordered structure has good plasticity, but lower strength and hardness. During the cooling process of welding, the Fe3 Al heat affected zone undergoes an ordered structural transformation, i.e., the transformation of the partially ordered B2 structure to the fully ordered D03 structure. This transformation process is an exothermic process, and the release of latent heat of phase change can eliminate defects such as excess vacancies in Fe3 Al and reduce the hardness of the Fe3 Al heat affected zone. Welding heat input affected the hardness of the softening zone of the Fe3 Al heat zone, with the increase of welding heat input, the minimum hardness gradually decreased, see Table 5.7. With the increase in welding heat input, high temperature residence time in the heat affected zone grow, the amount of diffusion of Al elements increased, the disordered structured also increased in the heat affected zone; in addition, during the welding cooling process, the ordered structure transformation will occur in the Fe3 Al heat affected zone. That is, the partially ordered B2 structure transform to the fully ordered D03 structure. This transformation process is an exothermic process, the release of the latent heat of phase change correspondingly increased, to eliminate defects such as excess vacancies in Fe3 Al. Under the combined effect of the above factors, the hardness of the heat affected zone of Fe3 Al is reduced. Table 5.7 Effect of welding heat input on the minimum microhardness in the heat affected zone of Fe3 Al Joint

Current I/A

Voltage U/V

Speed v/cm·s−1

Heat input E/ kJ·cm−1

Minimum hardness/HM

Fe3 Al/18-8

90

10

0.18

5.00

342

Fe3 Al/18-8

95

11

0.18

5.81

330

Fe3 Al/18-8

100

12

0.18

6.67

318

Note The effective heating coefficient of the arc η is taken as 0.75

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5 Joining of Iron-Aluminium Intermetallic Compounds

5.2.4 Shear Strength and Fracture Morphology of Fe3 Al/Steel GTAW Joints (1) Shear strength of Fe3 Al/steel GTAW joints The process parameters affect the organization of Fe3 Al/steel filled wire GTAW joints and determine the bond strength and fracture morphology of the joints. In order to study the mechanical properties of Fe3 Al filled wire GTAW joints, the shear strength of Fe3 Al/steel joints obtained under different welding parameters was measured using a CMT5150 micro-controlled electronic universal testing machine, the test results are shown in Table 5.8. Under the same process parameters and filler alloy (Cr23-Ni13), the shear strength of Fe3 Al/18-8 steel GTAW joint was the largest with 591 MPa; FeAl/18-8 steel GTAW joint was the second largest with 497 MPa; Fe3 Al/Fe3 Al GTAW joint had the smallest shear strength of 127 MPa. Microstructure analysis shows that Fe3 Al/Q235 steel and Fe3 Al/18-8 steel filled wire GTAW welds have similar microstructure (filled CR25-Ni13 wire), both with γ austenite as the matrix and containing a certain content of δ ferrite, and the shear strength of the welded joints is different from the proportion difference of γ phase of Fe3 Al/Fe3 Al joints. There is a higher content of Al elements solidly dissolved in the weld, forming a brittle phase, resulting in a high hardness and brittleness, and even local crackings occur along the grain boundary in the weld, causing lower shear strength. The welding heat input has an important effect on the shear strength of GTAW joints, Table 5.9 shows the variation of the shear strength of Fe3 Al/18-8 steel GTAW joints with the welding heat input. With the increase of welding heat input, the shear strength of Fe3 Al/18-8 steel joints gradually increases, and when the welding heat input is about 5.78 kJ/cm, the shear strength reaches a maximum value of 497 MPa, but when the welding heat input continue to increase, the shear strength begins to decrease. When the welding heat input is small, the joint cooling rate is faster, resulting in large welding stress and brittle phase; as the welding heat input increases, the joint cooling rate becomes slower, the welding stress is released, and the microstructure of the weld tends to be uniform, so the shear strength of Fe3 Al/18-8 steel GTAW joints Table 5.8 Test results for the shear strength of Fe3 Al filled wire GTAW joints Butt joint

Parameters/I × U Heat input Shear area Max loading Mean shear strength E/kJ·cm−1 A/mm2 F m /kN σ τ /MPa

Fe3 Al/Q235

105A × 11 V

5.78

105A × 11 V

5.78

Fe3 Al/Fe3 Al 105A × 11 V

5.78

Fe3 Al/18–8

26.5

15.5

26.1

15.3

26.4

14.1

25.5

12.7

25.4

4.1

26.4

4.5

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591 497 127

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Table 5.9 Effect of welding heat input on the shear strength of Fe3 Al/18-8 steel joints Current/A

Voltage /V

Heat input E/kJ·cm−1

90

10

4.50

Shear area A/mm2

Max loading F m /kN

Mean shear strength σ τ /MPa 469

24.3

11.3

24.5

11.6 14.1

105

11

5.78

26.4 25.5

12.7

120

12

7.20

28.9

14.9

29.5

14.2

497 481

gradually increases. However, when the welding heat input is too large, the joint superheats for a long time and the the microstructure of the weld coarsens, resulting in a decrease in the shear strength of the joint. (2) Fracture morphology of Fe3 Al/steel GTAW joints (1)

Fracture morphology of Fe3 Al/18-8 steel joints

The fracture morphology of the joints reflects the process of the sprouting, extension and fracture of the crack. Fe3 Al/18-8 steel joints have uneven fractures, and the fracture morphology of the partially fused zone of the joint is shown in Fig. 5.22. The fracture of the partially fused zone is mainly transgranular cleavage fracture (see Fig. 5.22a), with large cleavage plane, indicates that the Fe3 Al side fusion zone has coarse grains and the cleavage cracks are easy to destabilize and expand. There is a clear river pattern on the cleavage surface, and the river consists of many cleavage steps (see Fig. 5.22b). The presence of crystal defects such as dislocations, second phase particles, and inclusions in the fusion zone leads to the initiation of microcrack on the cleavage surfaces and between the grains, which reduces the shear strength of the joint and becomes the origin of cracks. A small number of tear ribs are also present in the brittle fracture zone in the partially fused zone of Fe3 Al/18-8 joints. (b)

(a)

Crack

Cleavage fractures

River pattern

Fig. 5.22 Shear fracture morphology in the partially fused zone of Fe3 Al/18-8 steel joints

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5 Joining of Iron-Aluminium Intermetallic Compounds

σ3 σ1

Micro-void

Fig. 5.23 Schematic diagram of the formation of parabola-type dimples in the Fe3Al/18-8 PFZ

The fracture in the GTAW weld zone of Fe3 Al/18-8 steel shows ductile fracture. The homogeneous mixture zone (HMZ) has obvious dimple tearing characteristics with a large depth, indicating that the toughness of this zone is good, and the dimples are elongated under shear stress, and some of them even penetrate each other, and a small amount of gray second phase particles are present in the dimple. The shear dimple in partially fused zone (PFZ) have obvious parabolic features and the dimple tearing is not obvious, which indicates that the toughness of this zone is lower than that of the homogeneous mixture zone. Dimple in fractures can be classified into two categories: equiaxed dimple and parabolic tough dimple. Parabolic shear dimples are easily formed under shear stress in opposite directions on the fracture surface, and the formation mechanism is shown in Fig. 5.23. It can be found that there are dark gray second-phase particles inside the dimple in partially fused zone, and the particle diameter is less than 5 μm. The analysis shows that the second phase particles in the shear dimple have an Al content of 64.54% (atom fraction) and high N and Fe content, and also contain small amounts of Cr, Mn, and Ni. Based on the composition, it is determined that these gray spherical precipitates are probably AlN and Fe-Al compounds with high Al content. AlN is a covalent bond compound with good chemical stability and remains stable in air at temperatures of 1000 °C and up to 1400 °C in vacuum. (2)

Fracture morphology of Fe3 Al/Q235 steel joints

The partially fused zone of Fe3 Al/Q235 steel GTAW welding joints have a shear fracture that is torn and uneven, with many small shiny planes present on the fracture, an obvious brittle fracture, as shown in Fig. 5.24. The fracture morphology is consistent with the coarse grain characteristics of the Fe3 Al parent material, with a coarse cleavage surface of a clear river pattern. The cleavage surface a consist of series of cleavage steps, almost perpendicular to the cleavage surface, and a few cracks along the crystal between the coarse grains. Further analysis of the area enclosed by the white box in Fig. 5.24a reveals that the cleavage steps actually consist of a series of smaller sized microsteps (Fig. 5.24b),

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5.2 Wire-Filled Tungsten Arc Welding of Fe3 Al and Steel (Q235, 18–8 Steel) (a)

255

(b) Cracks

River patterns Micro-steps Cleavage steps

Cleavage steps

Secondary cleavage steps

Fig. 5.24 Fracture morphology of the partially fused zone of Fe3 Al/Q235 joints

making the grain boundaries of the partially fused zone difficult to slip, leading to its higher brittleness, which is the reason for the origin of weld cracks from the partially fused zone in Fe3 Al joints. The two parallel primary cleavage surfaces are continuously extended forward under the action of shear stress, and with the increase of stress, the primary cleavage surface is expanded laterally to produce secondary cleavage, and the secondary cleavage surface connects the primary decomposition surface to form a right-angle cleavage step. The cleavage step is generally parallel to the crack expansion direction and perpendicular to the cracks, because the energy required to form a new free surface is minimal in this way. The shear fracture morphology of the Fe3 Al/Q235 welding joint is shown in Fig. 5.25, which is a mixture of brittle fracture and ductile fracture, and this brittletough combination can obtain a relatively high shear strength. The proportion of brittle fracture zone is larger in the partially fused zone, and the proportion of ductile fracture zone is larger in the homogeneous mixture zone. The transgranular cleavage fracture zone contains both river patterns and tongue patterns, see Fig. 5.25a. The ductile fracture zone consists of shearing dimple, which is surrounded by a cleavage zone, and the transition zone between the dimple zone and the cleavage zone consists of a number of tearing ribs and small cleavage facets, as shown in Fig. 5.25b. White second-phase particles and a few microholes present in the dimple, and the particles are high in Al, O and Fe, so these particles may be a mixture of Al2 O3 and a small amount of Fe-Al compounds. Under the action of shear stress, the weld metal is plastically deformed and microholes are formed with the second phase particles as nuclei. As the stress increases, the holes grow and interconnect and plastic fracture occurs, forming a shear dimple. (3) Fracture morphology of Fe3 Al/Fe3 Al joints Fe3 Al/Fe3 Al welding joints have relatively flush, slightly torn fractures with a brighter metallic luster. The experimental results show that the shear strength of Fe3 Al/Fe3 Al welding joints is much lower than that of Fe3 Al/steel welding joints, and there are cracks along the grain boundary in the weld. Matrix of the Fe3 Al/Fe3 Al

256

5 Joining of Iron-Aluminium Intermetallic Compounds (b) Dimples zone

(a)

River patterns

Dimples zone

Cleavage zone Tongue patterns

Partial (PMZ)

Homogeneous mixing zone (HMZ)

Fig. 5.25 Shear fracture morphology of the Fe3 Al/Q235 welding joints

weld is coarse α-Fe(Al) solid solution, the intergranular strength is relatively low, and is prone to fracture along the grain boundary when subjected to external forces. Figure 5.26 shows the fracture morphology characteristics of the Fe3 Al/Fe3 Al welding joint. The bonding state between subgrains in the weld can be seen, and microcracks exist between primary grains (Fig. 5.26a), which is consistent with the results of cracking analysis, and few microcracks occur between subgrains, indicating that the bond strength between subgrains is higher than that of primary grains. There is only a small number of transgranular cleavage fracture in Fe3 Al/Fe3 Al welding joints, and river patterns and secondary cleavage steps is present on the grain boundary, see Fig. 5.26b. Fracture of partially fused zone of Fe3 Al/Fe3 Al welding joint shows a significantly coarsened fracture morphology along the crystal, with a smooth surface and few precipitation phases and tearing marks. The presence of coarse grains and some smaller size sub-grains in the shear fracture plane is also consistent with its microstructure characteristics. The partial mixture zone shows the fracture morphology of columnar crystals, with long and oriented grains and a few tearing ribs on the grain boundaries.

(a)

(b) Crack

Steps Sub-grain River pattern

Intergranular cracks and sub-grains

Fig. 5.26 Shear fracture morphology of Fe3 Al/Fe3 Al joints

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Cleavage fracture

5.2 Wire-Filled Tungsten Arc Welding of Fe3 Al and Steel (Q235, 18–8 Steel)

257

(3) Fracture mechanism of in Fe3 Al/steel GTAW joints The Fe3 Al filled wire GTAW joints are brittle fracture, in which Fe3 Al/18-8 steel and Fe3 Al/Q235 steel joints are dominated by transgranular cleavage fracture, and Fe3 Al/Fe3 Al welding joints are dominated by intergranular fracture. The cleavage fracture process of Fe3 Al/18-8 steel joints can be divided into two stages: crack initiation and unstable propagation. The initiation and unstable propagation of cleavage fracture are subject to certain conditions, and Griffith proposed an energy theory of brittle fracture for brittle materials, stating that the condition for cleavage crack propagation is that the elastic energy released by cleavage is greater than the energy required to form a new surface for the crack. The conditions for crack propagation in a plane stress state are: σ ≥ σc ( σc =

2γ E πa

) 21 (5.1)

where γ —the surface energy of the material. A—half-length of the crack. E—modulus of elasticity. σ c —critical stress. The occurrence of cleavage fracture in Fe3 Al joints is also accompanied by a degree of ductile fracture, and thus the plastic deformation work at the crack front must also be overcome γ P . Orowen modified the Griffith energy condition to obtain the expressions for cleavage crack propagation in the plane stress state and in the plane strain state as follows. ( Planar stress state σ c = ( Plane strain state σ c =

2EγP πa

) 21

2EγP π a(1 − v 2

(5.2) )1/2 (5.3)

where γ P —work of plastic deformation. υ—Poisson’s ratio. On the basis of satisfying the energetic conditions described above, cleavage fracture can be divided into two steps, firstly microcrack nucleation and secondly microcrack propagation in the matrix. Microcracks in Fe3 Al joints tend to form at. (1) At the brittle second phase particles. When the Fe3 Al joint is deformed by shear stress, the brittle second-phase particles are not easily deformed, and under the action of additional forces and dislocation forces caused by the uncoordinated deformation, the second-phase particles are detached from the matrix or crack themselves to form microcracks. There is dimple in both Fe3 Al/18-8

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steel and Fe3 Al/Q235 steel joint fractures, formed with AlN or Al2 O3 second phase particles as cores, and these brittle particles become the origin of shear microcracking. (2) Where the slip zone is obstructed. There is a large amount of dislocation plugging in the fusion zone of Fe3 Al joints, which can easily become a nucleation for microcracking. (3) Where grain boundaries are weakened. Grain boundary embrittlement due to segregation of microelements at grain boundaries may also cause microcracks to develop along grain boundaries and then extend into the grain. This is the case for shear fracture in Fe3 Al/Fe3 Al joints. (4) At the intersection of twins. Microcracks can form at the intersection of deformation twins and tissue twins formed in Fe3 Al joints under shear, and microcracks can also form at the intersection of twins and the parent phase. After microcrack nucleation, the crack can only extend in the matrix when the local stress exceeds the critical stress, which implies a continuous loading of shear force in shear tests. In addition, since cleavage is the separation of atomic bonds fracture along a certain crystal plane, the tip of the crack nucleus that initiates the cleavage fracture should have the sharpness of the atomic spacing, and under the shear stress, the microcrack is connected to the main crack, causing the cleavage fracture of the Fe3 Al joint. Process of Fe3 Al/18-8 steel joint shear fracture: before the application of shear stress, crystal defects such as twins and dislocations exist in the Fe3 Al side fusion zone, and there is a high stress–strain zone around the defects, which becomes a potential crack origin. The shear stress near the shear surface is the largest, and the slip of grain boundaries and cleavage surfaces occurs first, forming slip steps, which in turn leads to the formation of the main crack. As the shear force increases, the microcracks already present at defects such as twin crystal substructures and dislocations in the Fe3 Al side fusion zone start to initiate and extend to the Fe3 Al heat affected zone and the shear surface. The propagation of microcracks to the heat-affected zone is due to the greater brittleness of the Fe3 Al heat-affected zone, the small energy required for crack propagation, and the small resistance to propagation; the propagation to the shear surface is due to the greater the shear stress, the greater the energy provided for microcrack propagation, and the faster the crack propagation rate. With the further increase of shear, these microcracks propagate and grow, and when converging with the main crack directly caused by shear, shear fracture of Fe3 Al/18-8 steel joints occurs. As the microcracks propagate into the Fe3 Al heat affected zone, leading to partial fracture in the Fe3 Al fusion and heat affected zones.

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5.3 Vacuum Diffusion Welding of Fe3 Al to Steel (Q235, 18-8 Steel)

259

5.3 Vacuum Diffusion Welding of Fe3 Al to Steel (Q235, 18-8 Steel) Fe3 Al intermetallic compound is brittle and hard, low plasticity and toughness, when using conventional fusion welding method to weld Fe3 Al intermetallic compound, the joint composition is complicated, there is a tendency to crack and brittle phase is easily generated, it is difficult to get welded joints that meet the requirements of use. The advanced vacuum diffusion welding process can suppress the generation of brittle phases near the Fe3 Al/steel joints.

5.3.1 Process Characteristics of Fe3 Al/Steel Vacuum Diffusion Welding (1) Test materials The test base materials were Fe3 Al intermetallic compound, Q235 steel and 18-8 austenitic steel. Fe3 Al intermetallic compound was prepared by vacuum induction melting method and annealed by homogenization at 1000 °C. The raw Fe was derusted by rolling in a ball mill before melting and the raw Al was cleaned with NaOH solution and dried. Vacuum was drawn to 10–2 Pa during the melting process. The Fe3 Al intermetallic compound was machined into sheets of 20 mm thickness using wire cutting method. The test 18-8 steel was a 1Cr18Ni9Ti austenitic stainless steel with a thickness of 8 mm and a microstructure of austenite + a small amount of δ-ferrite. Fe3 Al intermetallic compounds have a strong hydrogen embrittlement sensitivity, and the fusion welding process generates a large thermal stress at the joint, which can easily lead to the generation of weld cracks, which is the main obstacle to the application of Fe3 Al as a structural material and a challenge to be solved in the welding application of wear and corrosion resistance. (2) Diffusion welding equipment When Fe3 Al intermetallic compounds are fusion welded to steel, brittle intermetallic compounds with high aluminum content tend to form at the joint due to differences in thermophysical and chemical properties, which degrade the toughness of the welded joint. Using diffusion technology, the joining of Fe3 Al/Q235 steel as well as Fe3 Al/18-8 steel can be achieved by controlling the influence of process parameters on the microstructure of the Fe3 Al/steel diffusion welding interface. The test was conducted using Workhorse II vacuum diffusion welding equipment imported from C/VI, USA, with 45 kW heating power and 30 T double-acting hydraulic pressurization. The main parameters of the vacuum diffusion welding equipment are shown in Table 5.10.

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5 Joining of Iron-Aluminium Intermetallic Compounds

(3) Diffusion welding process and parameters (1) Pre-welding Preparation The surface of the specimen to be welded (Fe3 Al intermetallic compound, Q235 steel and 18-8 steel) is first machined on a grinder to ensure that the upper and lower surfaces of the specimen are parallel and that the surface finish is grade 6. Chemical methods are used to remove oxide film, oil and rust from the surface of the specimen to be welded before welding. The steps for mechanical and chemical treatment of the specimen surface are: sandpaper grinding → acetone cleaning → water rinsing → alcohol cleaning → blowing dry. The surface cleaned specimens to be welded (Fe3 Al intermetallic compound with Q235 and Fe3 Al with 18-8 steel) were stacked together in a diffusion welding vacuum chamber and mica sheets were placed on the contact area between the surface of the specimen to be welded and the indenter to prevent diffusion connection between the surface of the specimen and the indenter. The dimensions of the diffusion welded specimens are: 100 × 20 × 20 mm for Fe3 Al material; 100 × 20 × 20 mm for Q235 steel; and 100 × 20 × 8 mm for 18-8 steel. (2) Process route and parameters In order to improve the uniformity of heating of the welded parts in the diffusion welding process, graded heating is used and several holding time platforms are set; the cooling process uses circulating water cooling to 100 °C and then cools with the furnace. The process parameter curves of the diffusion welding process are shown in Fig. 5.27. The process parameters for diffusion welding of Fe3 Al to Q235 steel are shown in Table 5.11. The bond strength, fracture location and fracture morphology of Fe3 Al/Q235 steel diffusion welded joints depend on the process parameters, such as heating temperature, holding time, welding pressure and cooling rate. The heating temperature determines the diffusion activity of the elements; holding time determines the degree of homogenization of element diffusion at the Fe3 Al/Q235 steel interface; the pressure is to make the Fe3 Al/Q235 contact interface microplastic deformation, promote close contact between the base materials, prevent interface voids and control the deformation of the welded parts, otherwise it can promote element diffusion; the main role of cooling rate is to maintain the interface’s organizational properties stability. (3) Test specimen preparation During Fe3 Al/Q235 steel and Fe3 Al/18-8 steel diffusion weld interface structure and performance analysis, specimens are cut and etched to meet different test requirements. Fe3 Al intermetallic compound is hard, and the weldments are cut and machined into diffusion weld joint specimens using wire cutting method.

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Type

3033–1350 – 30 T

Company

C/VI

1.33

Final vacuum/10−5 Pa

Main parameters

1623

Temperature /K 30

Pressure /T

Table 5.10 Main parameters of Workhorse II vacuum diffusion welding equipment Power/kW 45

Box size/mm 304.5 × 304.5 × 457

380

Voltage/V

N2 , Ar

Shield gas

5.3 Vacuum Diffusion Welding of Fe3 Al to Steel (Q235, 18-8 Steel) 261

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5 Joining of Iron-Aluminium Intermetallic Compounds

Fig. 5.27 Process parameter curves for diffusion bonding of Fe3 Al/steel

Table 5.11 Process parameters for diffusion welding of Fe3 Al to Q235 steel Heat temperature/°C

Holding time/min

Heating speed /°C·min−1

Cooling speed /°C·min−1

Pressure/MPa

Vacuum/Pa

980–1080

30–60

15

30

12–18.5

1.33 × 10–4

Due to the different corrosion resistance between the Fe3 Al intermetallic compound and Q235 steel, when micro-etching is acted on Fe3 Al/Q235 steel diffusion welding joint, firstly 3% nitric acid alcohol solution is adopted to etch Q235 steel side, and then sealed with paraffin; then the aqua regia solution (HNOpp3 : HCl = 1:3) is adopted to etch Fe3 Al side, finally the paraffin on Q235 steel side is polished to remove for Fe3 Al/18-8 diffusion welded joint, the aqua regia solution is adopted to micro-etchingetch. During diffusion welding Fe3 Al with steel, due to the very differences of chemical composition and thermophysical properties between these two kinds of the materials, the elements diffusion in the Fe3 Al/steel contact interface occurs, when a certain concentration arrived, diffusion reactions will happen, a series of intermediate phase structure could be formed with different organizational properties from the base material, which can affect the organization and properties of Fe3 Al/steel diffusion welded joints. The formation of these phase structures is related to the elements contained in the base material, and the formation conditions mainly depend on the diffusion welding process parameters.

5.3.2 Shear Strength of the Fe3 Al/Steel Diffusion Weld Interface The welding process parameters directly affect the bonding characteristics of the diffusion welded interface, then determine the bond strength, joint fracture location

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263

8mm

8mm

5.3 Vacuum Diffusion Welding of Fe3 Al to Steel (Q235, 18-8 Steel)

Fe3Al 20mm (a) Fe3Al/Q235

Q235

Fe3Al 20mm

18-8

(b) Fe3Al/18-8

Fig. 5.28 Dimensions of shear specimens for Fe3 Al/Q235 steel and Fe3 Al/18-8 steel diffusion welded interface

and fracture morphology of the diffusion welded interface. In order To investigate the mechanical properties of the Fe3 Al/steel diffusion weld interface, the shear strength of the Fe3 Al/Q235 steel and Fe3 Al/18-8 steel diffusion weld interfaces with different process parameters was obtained by using a digital display pressure testing machine, the dimensions of the shear specimens are shown in Fig. 5.28. (1) Shear strength of Fe3 Al/Q235 steel diffusion weld interface For Fe3 Al/Q235 steel diffusion welded joints obtained under different process parameters, 20 × 8 × 8 mm specimens (2 specimens for each process parameter) were cut from the diffusion welded joint location using wire cutting method. The specimen surfaces were ground and then subjected to shear tests on a digital pressure tester, the test and calculated results are shown in Table 5.12. The test results showed that the shear strength of the Fe3 Al/Q235 steel diffusion welded interface increased from 39.9 to 112.3 MPa when the holding time was 60 min and the welding pressure was reduced from 17.5 to 12 MPa (keeping the joint free from macroscopic deformation), and the heating temperature was increased from 1000 to 1060 °C (Fig. 5.29). This is due to the higher energy obtained for atomic diffusion near the Fe3 Al/Q235 interface as the heating temperature increases, which diffuses more fully at the interface and a good metallurgical bond is formed near the interface. However, when the heating temperature was increased to 1080 °C, the shear strength of the Fe3 Al/Q235 steel diffusion welded interface decreased to 82.1 MPa. This is due to the fact that the microstructure near the Fe3 Al/Q235 steel diffusion welded interface was gradually coarsened at too high a heating temperature under the condition of keeping the diffusion welded joint free from macroscopic deformation, resulting in some reduction in the shear strength of the diffusion welded interface. At the heating temperature of 1060 °C, with the increase of holding time, the atoms near the diffusion interface get sufficient mutual diffusion and interfacial reaction occurs, a dense intermediate diffusion reaction layer was formed, therefor the shear strength of the diffusion interfac’xce of Fe3 Al/Q235 steel is significantly increased. Because the extended holding time can induce the atoms near the diffusion welding

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5 Joining of Iron-Aluminium Intermetallic Compounds

Table 5.12 Experimental and calculated results for the shear strength of the diffusion welded interface of Fe3 Al/Q235 steel No.

Parameters (T × t, P)

Shear area/mm2

Max load F m /N

01

1000 °C × 60 min, 17.5 MPa

9.98 × 8.02

3346

40.8

02

1000 °C × 60 min, 17.5 MPa

9.97 × 7.99

3115

39.1

03

1020 °C × 60 min, 17.5 MPa

9.97 × 7.98

5370

67.5

04

1020 °C × 60 min, 17.5 MPa

9.98 × 8.00

5395

67.6

05

1040 °C × 60 min, 15.0 MPa

9.95 × 7.96

7072

89.2

06

1040 °C × 60 min, 15.0 MPa

9.98 × 8.02

7408

92.5

07

1060 °C × 30 min, 15.0 MPa

9.98 × 8.00

3377

42.3

08

1060 °C × 30 min, 15.0 MPa

10.00 × 7.98

3551

44.5

09

1060 °C × 45 min, 15.0 MPa

9.97 × 7.96

7901

96.8

10

1060 °C × 45 min, 15.0 MPa

9.96 × 7.98

7941

97.2

11

1060 °C × 60 min, 12.0 MPa

9.93 × 7.96

8960

113.3

12

1060 °C × 60 min, 12.0 MPa

10.00 × 7.93

8834

101.4

13

1080 °C × 60 min, 12.0 MPa

9.95 × 7.98

6392

80.5

14

1080 °C × 60 min, 12.0 MPa

10.02 × 7.96

6668

83.6

Shear strength σ τ /MPa

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Mean shear strength σ τ /MPa 39.9

67.5

90.8

43.4

97.0

112.3

82.1

5.3 Vacuum Diffusion Welding of Fe3 Al to Steel (Q235, 18-8 Steel)

265

Fig. 5.29 Variation of shear strength of Fe3 Al/Q235 steel diffusion weld interface with heating temperature

interface to get sufficient mutual diffusion and a certain diffusion reaction occurs, a dense diffusion welding interface transition zone was formed. However, when the holding time was increased to 80 min with aheating temperature of 1080 °C and welding pressure of 12 MPa, a significant macroscopic plastic deformation was found in the Fe3 Al/Q235 joint. Therefore, the heating temperature should not be too high while maintaining the diffusion welded joint without deformation, as the organization of the Fe3 Al/Q235 steel diffusion welded joint will grow up at high temperature, which is not conducive to ensuring the shear strength of the joint. The analysis of the shear strength of Fe3 Al/Q235 steel diffusion welded interface shows that the heating temperature is controlled at about 1060 °C, holding time is 45–60 min and keeping the diffusion welded joint without macroscopic deformation (p = 12–15 MPa), a tight bonding and high shear strength Fe3 Al/Q235 steel diffusion welded joint without microscopic voids can be obtained. (2) Shear strength of the Fe3 Al/18–8 diffusion weld interface. (2) Shear strength of the Fe3 Al/18-8 diffusion weld interface A 20 × 8 × 8 mm specimen (2 specimens for each process parameter) was cut from the Fe3 Al/18-8 steel diffusion welded joint by using wire cutting method. The specimen surfaces were ground and then subjected to shear tests on a digital pressure tester, the test and calculated results are shown in in Table 5.13. The effect of heating temperature and holding time on the shear strength is shown in Fig. 5.30. The shear strength of the Fe3 Al/18-8 steel diffusion weld interface increased from 149.9 to 226.2 MPa when the heating temperature was increased from 980 to 1040 °C, see Fig. 5.30a. However, the shear strength of the Fe3 Al/18-8 steel diffusion welded interface increased rapidly with increasing heating temperature below 1000 °C, the

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Table 5.13 Experimental and calculated results for the shear strength of the diffusion welded interface of Fe3 Al/18-8 steel No.

Parameters (T × t, P)

Shear area/mm2

Max load F m /N

Shear strength σ τ /MPa

Mean shear strength σ τ /MPa

01

980 °C × 60 min, 17.5 MPa

9.98 × 8.01

11,951

149.5

149.9

02

980 °C × 60 min, 17.5 MPa

9.97 × 7.99

11,987

150.4

03

1000 °C × 60 min, 17.5 MPa

9.98 × 8.02

17,023

201.7

04

1000 °C × 60 min, 17.5 MPa

9.97 × 7.99

18,834

196.4

05

1020 °C × 60 min, 17.5 MPa

9.97 × 7.98

18,956

218.8

06

1020 °C × 60 min, 17.5 MPa

9.98 × 8.00

16,302

194.4

07

1040 °C × 60 min, 15.0 MPa

9.95 × 7.96

19,977

229.9

08

1040 °C × 60 min, 15.0 MPa

9.98 × 8.02

19,296

222.5

09

1040 °C × 45 min, 15.0 MPa

9.98 × 8.00

14,115

176.8

10

1040 °C × 45 min, 15.0 MPa

10.00 × 7.98

14,380

180.2

11

1040 °C × 30 min, 15.0 MPa

9.98 × 7.96

13,636

170.8

12

1040 °C × 30 min, 15.0 MPa

10.08 × 7.98

11,407

143.0

13

1040 °C × 15 min, 15.0 MPa

9.99 × 7.99

4853

60.8

14

1040 °C × 15 min, 15.0 MPa

10.00 × 8.08

4298

53.8

198.6

211.5

226.2

178.5

156.9

57.0

(continued)

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5.3 Vacuum Diffusion Welding of Fe3 Al to Steel (Q235, 18-8 Steel)

267

Table 5.13 (continued) No.

Parameters (T × t, P)

Shear area/mm2

Max load F m /N

Shear strength σ τ /MPa

Mean shear strength σ τ /MPa

15

1040 °C × 80 min, 15.0 MPa

9.99 × 7.99

13,426

168.2

172.4

16

1040 °C × 80 min, 15.0 MPa

10.00 × 8.08

14,269

176.6

17

1060 °C × 60 min, 12.0 MPa

9.98 × 7.99

15,804

198.2

18

1060 °C × 60 min, 12.0 MPa

9.90 × 7.96

14,420

182.9

190.6

Fig. 5.30 Effect of heating temperature and holding time on the shear strength of the Fe3 Al/18-8 Steel diffusion weld interface

interface shear strength increased relatively slowly with increasing heating temperature was in the range of 1000–1040 °C. The shear strength of the Fe3 Al/18–8 steel diffusion welding interface gradually decreases as the heating temperature continued to increase above 1040 °C. This is due to the coarsening of the microstructure in the transition zone of the diffusion welded interface when the heating temperature is too high, which results in a decrease in the shear strength. At heating temperature of 1040 °C and welding pressure of 15 MPa, the atoms near the diffusion welding interface of Fe3 Al/18-8 steel were sufficiently diffused with each other and diffusion reactions occurred as the holding time increased from 15 to 60 min, the dense organization in the interface transition zone made a significant increase of its shear strength from 57 to 226 MPa, see Fig. 5.25b. The shear strength of the Fe3 Al/18-8 steel diffusion welded interface gradually decreased when the

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5 Joining of Iron-Aluminium Intermetallic Compounds

holding time exceeded 60 min and continued to increase. This is due to the increase in the width of the transition zone formed near the interface and the coarsening of the structure after the excessive holding time. Test results show that during Fe3 Al and 18-8 steel diffusion welding, in order to obtain a weld with good interface bond and high shear strength, the appropriate process parameters are: heating temperature is controlled at about 1040 °C, holding time is about 45–60 min, welding pressure is about 12–15 MPa.

5.3.3 Microstructural Characteristics of the Fe3 Al/Steel Diffusion Weld Interface (1) Compartmentalization of the transition zone When Fe3 Al is diffusion welded with steel, under the combined effect of process parameters (T, t, p) and concentration gradient, the elements in the base material continuously diffuse to the contact interface, when a certain concentration is reached, the diffusion reaction occurs, and a diffusion reaction layer with a different organizational structure from that of the base material is formed near the contact interface between the base material. These areas with different organizational structures from the two base materials are called the diffusion welding interface transition zone, and there is an obvious diffusion transition zone between Fe3 Al and Q235 interfaces. The Fe3 Al/steel diffusion weld interface transition zone consists of a mixed transition zone and a transition zone close to the sides of the welded material. The microscopic area formed near the original contact interface after diffusion welding of Fe3 Al to steel is the mixed transition zone; the characteristic area between the mixed transition zone and Fe3 Al and Q235 steel (or 18-8 steel) is the transition zone close to the sides of the base material. The compartmentalization of the Fe3 Al/steel diffusion weld interface transition zone is shown in Fig. 5.31a. Q235

Fe3Al

Fe3Al Transition zone near Fe3Al

Original contact surface

Interfacial transition zone Mixed transition zone Q235(or 18-8)

Transition zone near Fe3Al Transition zone near Q235(or 18-8)

Transition zone near Q235

Interfacial transition zone

Mixed transition zone (a) Division of interfacial transition zone

(b) Microstructure

Fig. 5.31 Illustration of the transition zone delineation at the Fe3 Al/steel diffusion welding interface

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5.3 Vacuum Diffusion Welding of Fe3 Al to Steel (Q235, 18-8 Steel)

269

The width of the diffusion welding interface transition zone depends on the state of the original contact interface between Fe3 Al and steel and the diffusion welding process parameters, with the smaller of the distance between the original contact interface, the higher of the heating temperature, the longer of the holding time and the higher of the pressure during diffusion welding, the wider of the diffusion welding interface transition zone will be. The tissue characteristics of the Fe3 Al/Q235 diffusion welded interface transition zone observed by SEM are shown in Fig. 5.31b. After diffusion welding of Fe3 Al/Q235 steel, the original contact interface disappeared due to the mutual diffusion of elements, a region with more enriched white particles was formed between Fe3 Al and Q235 steel, which is the diffusion welded interface transition zone. The area near the sides of Fe3 Al and Q235 steel, due to the diffusion of elements, the micro structure and structural morphology have a certain extent change, two transition zones near the base material are formed. (2) Microstructure of the transition zone at the diffusion weld interface of Fe3 Al/Q235 steel The Fe3 Al/Q235 steel diffusion welded joints were prepared as a series of metallographic specimens. As the corrosion resistance of Fe3 Al intermetallic compound and Q235 steel is different greatly, the Fe3 Al intermetallic compound side was corroded with aqua regia solution (HNO3 : HCl = 1:3) and the Q235 steel side was corroded with 3% nitric acid alcohol solution. The microstructure of the Fe3 Al/Q235 steel diffusion welded joint specimens were observed by using a metallographic microscope and JXA-840 scanning electron microscope (SEM).The tissue characteristics of the Fe3 Al/Q235 steel diffusion welded head region are shown in Fig. 5.32. As shown in the figure, the Fe3 Al/Q235 steel diffusion welding interface has obvious diffusion characteristics, the diffusion welding interface and the transition zone near the base material on both sides are interlaced. Interface transition zone near the Fe3 Al side of the microstructure across the diffusion welding interface to the (b)

(a) Q235

Q235

Fe3Al Fe3Al

25µm OM, 100

SEM, 1000

Fig. 5.32 Organizational characteristics of the diffusion welded head zone of Fe3 Al/Q235 steel

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5 Joining of Iron-Aluminium Intermetallic Compounds

Q235 steel side and continuous extension, the interface is inlaid occlusal type. In Fe3 Al/Q235 steel diffusion welding interface transition zone, the microstructure near the Fe3 Al side is coarse, discontinuous distribution of precipitation phase existed at the grain boundary due to the diffusion of Al, Cr and other elements, the micro structure was fine, and mostly appeared like equiaxed crystal distribution. During diffusion welding, a coarse organization was in the Fe3 Al side due to then grain growth, see Fig. 5.32a. The mutual diffusion of elements between Fe3 Al and Q235 steel results in a change of the organization near the diffusion welding interface and a change in the direction of grain growth. Due to the narrow transition zone at the diffusion welding interface and slow cooling rate, the grain is most suitable as a ready surface for crystallization at the interface connecting the base material, which is most favorable for crystallization. The diffusion welding interface organization is easily formed on the basis of the base material and grows along the heat conduction direction in a meritocratic epitaxial manner. To investigate the variation of microstructure and properties near the diffusion weld interface of Fe3 Al/Q235 steel, the grain size of a series of Fe3 Al/Q235 steel diffusion welded microstructure near the interface was rated using a XQF-2000 +3 microimage analyzer. According to the equation D2 = 1/2 N (D is the grain diameter and N is the grain size rating) for the calculation of grain diameter, the relative content of grain size and precipitated phases in the transition zone of the Fe3 Al/Q235 steel diffusion weld interface were tested and calculated as shown in Table 5.14. As shown in Table 5.14, the microstructure gradually refines as it transitions from the Fe3 Al matrix across the interface to the Q235 steel side, and the grain diameter is reduced from 250 to 112 μm. Therefore, the organization in the transition zone of the Fe3 Al/Q235 steel diffusion welding interface is finer than that of the Fe3 Al matrix, and the diffusion of elements near the interface is more uniform, which is conducive to improving the Fe3 Al/Q235 steel diffusion welded joint strength properties. The degree of microstructure coarseness in different parts of the transition zone of the diffusion welded interface of Fe3 Al/Q235 steel is related to the diffusion welding heating temperature and holding time. The microstructure characteristics of the Fe3 Al/Q235 steel diffusion welding interface transition zone at different heating temperatures and holding times are shown in Fig. 5.33. Table 5.14 Grain size and relative content of precipitated phases in the transition zone of the diffusion welded interface of Fe3 Al/Q235 steel Location

Fe3 Al

Fe3 Al transition zone

Mixd transition zone

Q235 transition zone

Grain degree

1.02

1.50

2.05

3.30

Grain diameter/μm

250

210

173

112

Precipitation/%

14.6, 12.8, 13.0 (13.5)

11.8, 13.0, 12.1 (12.3)

21.3, 26.5, 27.8 (25.3)

5.4, 5.6, 7.0 (6.0)

Note Data in parentheses are test averages

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271

(b)

Fe3Al

Q235

Transition zone 10µm Characteristics of transition zone

Precipitations

Fig. 5.33 Characteristics of the transition zone and precipitation phase at the diffusion weld interface of Fe3 Al/Q235 steel

With the increase of the heating temperature and the extension of the holding time, the width of the interfacial transition zone of Fe3 Al/Q235 steel diffusion welding gradually increased and the organization gradually coarsened due to the sufficient diffusion of the elements near the interface. When the temperature was increased to 1060 °C and the holding time was 60 min, the width of the interfacial transition zone of Fe3 Al/Q235 diffusion welding increased to 38 μm and the grain diameter of the microstructure reached 180 μm. Observed under scanning electron microscopy (SEM), there are some white precipitation phases in the transition zone of the Fe3 Al/Q235 steel diffusion welding interface, these precipitation phases are gathered at the junction of the mixed transition zone and the diffusion transition zone near the Fe3 Al side, mostly discontinuous distribution along the grain boundary, the microscopic morphological characteristics of the precipitation phases are shown in Fig. 5.33b. The composition analysis of some of the precipitated phases by electron probe (EPMA) showed that the C and Cr content in the particles of the precipitated phases in the transition zone of the diffusion welded interface of Fe3 Al/Q235 steel was higher and the Fe and Al content was lower than that of the matrix (Table 5.15). This is due to the fact that the C and Cr elements in the transition zone microstructure of Fe3 Al/Q235 steel do not have time to diffuse sufficiently during the diffusion welding process, which is the result of bias aggregation inside the crystal. The precipitated phase particles clustered at the junction between the diffusion welding interface and the diffusion transition zone near the Fe3 Al side are caused by the difference in fractional distortion energy. During the diffusion welding, the bigger difference in fractional distortion energy is bigger as the difference of radius of elements atoms is bigger. The radius of Cr atoms (RCr = 0.185 nm) is larger than the radius of Fe atoms (RFe = 0.125 nm) while the radius of C atoms (RC = 0.077 nm) is much smaller than the radius of Fe atoms during the diffusion welding process. The difference between the dot distortion energy caused at the junction of the diffusion welding interface and the diffusion transition zone near the Fe3 Al side

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Table 5.15 Electronic Probe (EPMA) analysis of the transition zone at the diffusion weld interface of Fe3 Al/Q235 steel Location Fe3 Al

Precipitation phase

Points

Composition/% Fe

Al

C

Cr

Nb

Zr

1

82.60

16.62

0.14

0.52

0.05

0.07

2

82.80

16.40

0.13

0.49

0.03

0.05

3

81.91

17.10

0.13

0.51

0.03

0.02

4

82.21

16.90

0.13

0.54

0.01

0.01

5

81.69

16.35

0.55

1.18

0.01

0.02

6

81.94

15.90

0.51

1.28

0.03

0.04

7

82.54

15.47

0.40

1.32

0.04

0.03

8

81.89

16.23

0.22

1.26

0.04

0.06

is large, resulting in the C and Cr atoms being biased in the region near the grain boundary and its interface. Also, the smaller the solid solution degree of the solute atoms (C, Cr), the greater the tendency to produce grain boundary adsorption in the Fe3 Al matrix. Therefore, C and Cr with very small solid solution in Fe elements will be biased to be concentrated near grain boundaries or interfaces and exist as precipitated phases in the transition zone of the Fe3 Al/Q235 steel diffusion welding interface. (3) Microstructure of the Fe3 Al/18-8 diffusion welding transition zone The 18-8 steel has more Cr and Ni elements, a microstructure with multiple morphological structures near the Fe3 Al/18-8 steel diffusion welding interface is formed due to the complex diffusion path of the elements during diffusion welding is more. Figure 5.34 illustrates the microstructure characteristics of the transition zone of the diffusion welded interface of Fe3 Al/18-8 steel. It can be seen that there are obvious diffusion features at the contact interface between Fe3 Al intermetallic compound and 18-8 steel, and three diffusion reaction layers A, B and C are formed in the Fe3 Al/18-8 steel diffusion welding interface transition zone, and each reaction layer is interleaved with each other. Due to the diffusion of Al, Fe, Cr, and Ni elements in the transition zone at the Fe3 Al/18-8 interface, the organization within the diffusion reaction layer is more complex. The organization within the diffusion reaction layer A near the Fe3 Al side is characterized by some white dots distributed on the matrix, while the diffusion reaction layer C near the 18-8 steel side is distributed with irregularly shaped precipitation phases, and the organization of the intermediate diffusion reaction layer B is finer, with both white dots and irregular precipitation phases present. In order to investigate the influence of the organizational and structural characteristics of the reaction layers A, B and C in the transition zone of the diffusion weld interface on the performance of the Fe3 Al/18-8 steel diffusion welded joints, the composition of each diffusion reaction layer A, B and C was analyzed using an electronic probe (EPMA) and the actual results were determined in Table 5.16.

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5.3 Vacuum Diffusion Welding of Fe3 Al to Steel (Q235, 18-8 Steel)

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Fig. 5.34 Microstructural characteristics of the transition zone at the diffusion weld interface of Fe3 Al/18-8 steel

Fe3Al

18-8 A B C 25µm

Table 5.16 Electron probe analysis of the reaction layer in the transition zone at the diffusion weld interface of Fe3 Al/18-8 steel Diffusion layer

Al

Cr

Ni

Fe

Ti

A

12.5

8.9

4.4

73.9

0.3

B

9.5

13.2

4.5

72.3

0.5

C

2.5

17.2

8.0

71.8

0.5

As seen in Table 5.16, the elemental contents of Fe and Ti do not vary much within the three diffusion reaction layers A, B and C, while the elemental contents of Al, Cr and Ni vary greatly. In the transition from diffusion reaction layer A to diffusion reaction layer C, the elemental content of Al decreases, while the elemental contents of Cr and Ni keep increasing. It can be determined that the matrix organization structure of diffusion reaction layer A is still Fe3 Al phase, and the organization structure of diffusion reaction layer C is α-Fe(Al) solid solution. The precipitates on the matrix are Cr and Ni-rich precipitation or reaction phases, which are due to the concentration gradient of elements between Fe3 Al and 18-8 steel, prompting the mutual diffusion and diffusion reaction of Cr and Ni elements.

5.3.4 Microhardness of Fe3 Al/Steel Diffusion Welded Joints In order to determine the changes of the organization properties of Fe3 Al/Q235 steel and Fe3 Al/18-8 steel diffusion welded joints, the microhardness of different areas near the diffusion weld interface of Fe3 Al/Q235 steel and Fe3 Al/18-8 steel was tested using a SHIMADZU microhardness tester with a loading load of 25 g and a loading time of 10 s.

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5 Joining of Iron-Aluminium Intermetallic Compounds 600

Transition zone

(a) Microhardness /HM

550

Fe3Al

Q235

Fe3Al

Q235

500 450 400 350

1060¡æ¡Á60min 1060¡æ¡Á30min

300

1020¡æ¡Á60min

25µm 250

0

10

20

30

40

Distance /μm Location of measurement

Microhardness distribution

Fig. 5.35 Microhardness distribution in the diffusion welded head area of Fe3 Al/Q235 steel

(1) Microhardness of diffusion welded joints of Fe3 Al/Q235 steel The microhardness measurements near the diffusion interface of Fe3 Al/Q235 steel are shown in Fig. 5.35 and it can be seen that the microhardness of Fe3 Al base material after diffusion welding is about 490 MH and the microhardness of Q235 steel is 340 MH and the microhardness of Fe3 Al/Q235 steel diffusion welding interface transition zone varies with the process parameters. When the heating temperature is 1020 °C and the holding time is 60 min, the microhardness of the diffusion welding interface transition zone from the Fe3 Al side across the interface to the Q235 steel decreases firstly and then increases, and the peak microhardness (550 HM) appears in the mixed transition zone. This is due to the diffusion reaction of Al elements near the Fe3 Al side which induced the disordered transformation of Fe3 Al phase structure, so that the microhardness of the transition zone of the interface near the Fe3 Al side is reduced; near the Fe3 Al/Q235 steel diffusion welding interface, the elements is not fully diffused but aggregated at a lower heating temperature (T = 1020 °C), the formed phase structure has a higher microhardness, a peak microhardness appears in the mixed transition zone. At a heating temperature of 1060 °C and a short holding time (t = 30 min), a decrease of microhardness was observed in the transition zone on both sides of the Fe3 Al/Q235 steel diffusion weld interface. This is due to the shorter holding time which induced that the elements do not diffuse sufficiently and microscopic voids formed by the Kirkendall (Kirkendall) effect do not disappear completely. The peak microhardness of the diffusion welded interface at higher heating temperature (T = 1060 °C) is 520 HM, which is slightly lower than the peak microhardness at the heating temperature of 1020 °C (550 HM). This is due to the full diffusion of elements at higher temperatures and different phase structures formed through the diffusion reaction. The Fe-Al alloy state diagram is divided into two phase zones, Fe-rich zone and Al-rich zone by composition, and each phase zone is roughly divided into two

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Table 5.17 Microhardness of Fe-Al alloys for possible formation of intermetallic compounds Compounds

Al/% Phase diagram

Microhardness/HM Chemical analysis

Fe3 Al

13.87

14.04

350

FeAl

32.57

33.64

640

FeAl2

49.13

49.32

1030

Fe2 Al5

54.71

54.92

820

FeAl3

59.18

59.40

990

Fe2 Al7

62.93

63.32

1080

regions, the high temperature zone and the low temperature zone. In The Fe-rich, γ, α, and B2 (i.e., Fe-Al ordered phase) phases appears in Fe-Al alloy when solidifying at high temperatures; when cooling to low temperature, the three main phases are B2, α, and D03 (i.e., Fe3 Al ordered phase). Solid solutions, intermetallic compounds and co-crystals can be formed between Fe and Al elements. The solubility of Fe in solid aluminum is very small, at 225– 600 °C, the solubility of Fe in Al is 0.01–0.022%; at the eutectic temperature of 655 °C, the solubility of Fe in Al is 0.53%. At room temperature Fe is almost completely insoluble in Al, so aluminum alloys containing traces of Fe will develop the intermetallic compound FeAl3 during cooling. The super dotted structure of Fe3 Al is formed at room temperature when Al content is 13.9–20%, FeAl is formed when Al content is 20–36%. As the Al content increases, brittle phases such as FeAl2 , Fe2 Al5 , and FeAl3 appear successively. The microhardness of the possible intermetallic compounds formed by Fe-Al alloys is shown in Table 5.17. As microhardness measurement results of the Fe3 Al/Q235 steel diffusion welding joints, when heating temperature was controlled at about 1060 °C, holding time was controlled in 45–60 min, high hardness brittle phase (such as FeAl2 , Fe2 Al5 , FeAl3 , Fe2 Al7 , etc.) does not appear in Fe3 Al/Q235 steel diffusion welding joint. This microhardness property determines that the Fe3 Al/Q235 steel diffusion welded joint has better tissue properties, which can improve the toughness of the diffusion welded interface region, prevent microcracking, and help to improve the macromechanical properties of the transition zone of the Fe3 Al/Q235 steel diffusion welded interface. In summary, Fe3 Al/Q235 steel diffusion welded joints are mainly composed of Fe3 Al phase and α-Fe(Al) solid solution, with the presence of a small amount of FeAl phase, but do not have Fe-Al brittle phase with higher aluminum content, which is conducive to improving the toughness and crack resistance of the joints and ensuring the quality of the welded joints. (2) Microhardness of diffusion welded joints of Fe3 Al/18-8 steel In order to determine the performance of the Fe3 Al/18-8 steel diffusion welded joints, the microhardness of the Fe3 Al/18-8 steel diffusion welded head area was determined using a microhardness tester. with the holding time of 60 min, heating

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5 Joining of Iron-Aluminium Intermetallic Compounds

temperature of 1000, 1040 and 1060 °C, respectively, Fe3 Al/18-8 steel diffusion welding joints of the microhardness measured point location and microhardness distribution is shown in Figs. 5.36 and 5.37. At a heating temperature of 1000 °C, from the Fe3 Al side through the diffusion weld interfacial transition zone to the 18-8 steel, the microhardness values in the interfacial transition zone near the Fe3 Al side decrease due to the presence of a large number of microscopic voids and the disordered transformation of Fe3 Al. In a narrow region of the diffusion welded interfacial transition zone, the microhardness increases abruptly to a peak of 720 HM. at a heating temperature of 1040 °C, the microhardness transitions from the Fe3 Al side to the 18-8 steel with an almost continuous change in microhardness, with microhardness values of 500 HM in the interfacial transition zone on the Fe3 Al side, increasing slightly to 520 HM in the 1000

(a)

900

Microhardness /HM

18-8 Fe3Al

50µm

800

18-8

700 600 500 400 300 200 100 0

-20

Location of measurement

Transition zone Fe3Al

0

20

40

60

80

100

120

Distance /μm Microhardness distribution

Fig. 5.36 Microhardness distribution in the diffusion welded head zone of Fe3 Al/18-8 steel

Fig. 5.37 Microhardness distribution in the diffusion welded head area of Fe3 Al/18-8 steel

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5.3 Vacuum Diffusion Welding of Fe3 Al to Steel (Q235, 18-8 Steel)

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diffusion welded interfacial transition zone, and then transitioning to the 18-8 steel side with microhardness decreases all the way down to 300 HM. At a heating temperature of 1060 °C, the peak microhardness of the transition zone of the Fe3 Al/18-8 steel diffusion welding interface was 700 HM, while the microhardness value on the 18-8 steel side was reduced as the lower heating temperature (T = 1000 °C), a microhardness was only about 280 HM, which was due to the gradual growth and coarsening of the austenite tissue in the 18-8 steel at higher temperatures. In conclusion, the lower of the heating temperature, the less adequate diffusion of the elements, which causes the elements within the intermediate diffusion reaction layer to aggregate and increase in concentration, resulting in the formation of a phase structure with a higher microhardness than the Fe3 Al matrix, peak points of higher microhardness appear in the transition zone of the Fe3 Al/18-8 steel diffusion weld joints.

5.3.5 Element Diffusion Near the Interface and Transition Zone Width (1) Diffusion of elements near the Fe3 Al/18-8 interface The Fe3 Al intermetallic compound has better oxidation and corrosion resistance than 18-8 steel and is cheaper, so diffusion welding of Fe3 Al to 18-8 steel has a promising application in production. The elemental distribution of Fe3 Al/Q235 steel diffusion welded joints is shown in Fig. 5.38. The actual electron probe measurements of elements near the interface of Fe3 Al/18-8 steel diffusion welding are shown in Fig. 5.39. The concentration of Cr Fig. 5.38 Elemental distribution of diffusion welded joints of Fe3 Al/Q235 steel

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Fig. 5.39 Elemental distribution of diffusion welded heads of Fe3 Al/18-8 steel

elements fluctuates at 10–25 μm from the interface on the 18-8 steel side, which is due to the influence of Al and Ni elements during the diffusion process, resulting in the deviation of Cr elements near the interface. In −20 to −5 μm interval range of the diffusion transition zone from the interface near the Fe3 Al side, the slope of the Al, Ni elements distribution curve is small, the concentration gradient changes slowly. The measured values of Al, Ni element concentration fluctuates in the interface near the 18-8 steel side from the interface 5–25 μm interval; Al element concentration gradually decreased to 0, Ni element distribution gradually increased to 9% which is the stable Ni concentration value of 18-8 steel. On the Fe3 Al/18-8 steel diffusion reaction layer near Fe3 Al side, the Al content is high, different types of Fe-Al intermetallic compounds are formed due to and the reaction between Fe and Al diffused mainly from Fe3 Al. x-ray diffraction (XRD) analysis shows that as the heating temperature increases from 1020 to 1060 °C, the Fe3 Al/18-8 steel diffusion compound formed on the near Fe3 Al side of the reaction layer gradually changes from (FeAl2 + Fe2 Al5 ) → (Fe3 Al + FeAl + Fe2 Al5 ) to (Fe3 Al + FeAl). At lower heating temperatures, the Al elements acquire low energy and poor diffusion activity, and only gather in the edge zone near the Fe3 Al interface before they have time to diffuse into the 18-8 steel. Therefore, the Al element concentration is high on the Fe3 Al side, and it combines with the Fe element in the Fe3 Al matrix to form FeAl2 and Fe2 Al5 new phases. FeAl2 and Fe2 Al5 have high brittleness and microhardness values up to 1000 HM due to the high Al content, and these two new phases tend to cause thermal vacancies during heating, leading to point defects, have lower plasticity and toughness at room temperature, and are prone to destructive fracture. Increasing the diffusion welding temperature can induce the diffusion of Al atoms in FeAl2 and Fe2 Al5 to form Fe3 Al + FeAl mixed phases.

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18-8 steel contains alloying elements such as Ni, Cr and Ti, which gain some energy during the diffusion welding process and diffuse into the Fe3 Al/18-8 steel contact interface, form various compounds with the Fe and Al elements in Fe3 Al. When the heating temperature is 1020 °C, the compounds formed in the Fe3 Al/18-8 steel diffusion welded joint are mainly α-Fe(Al) solid solution; while when the temperature is increased to 1040 °C, it includes not only α-Fe(Al) solid solution, but also Ni3 Al intermetallic compounds; when the temperature is up to 1060 °C, a small amount of Cr2 Al phase appears in the diffusion layer, which affects the toughness of the Fe3 Al/18-8 steel diffusion welded joints. (2) Width of diffusion welding interface transition zone When Fe3 Al is diffusion welded to steel, the elements diffuse from one side across the interface to the other side, obeying the one-dimensional diffusion law. The variation of the concentration of elements near the interface with distance and time obeys Fick’s second law one-dimensional infinite medium non-stationary conditions of the diffusion equation, diffusion welding interface transition zone width and holding time in accordance with the parabolic law. ) ( Q x2 = K p(t − t0 ), K p = K 0 exp − RT

(5.4)

where x—width of the interfacial transition zone, μm. K p —diffusion rate of the element, μm2 /s. t—holding times, s. t 0 —latency time, s. K 0 —temperature-dependent coefficient. Q—diffusion activation energy, J/mol. T—heating temperature, K. R—gas constant. The width of Fe3 Al and steel diffusion weld interface transition zone is related to the diffusion rate of the element in the transition zone. When calculating the diffusion rates of elements in the complex phase structure system in the transition zone of the Fe3 Al/Q235 steel and Fe3 Al/18-8 steel diffusion weld interface, the diffusion weld interface transition zone is considered as a superposition of reaction layers with high volume content of the phase structure, with little influence of other elements in the transition zone; and the diffusion reaction near the interface reaches a quasi-equilibrium state. The diffusion rates of elements at the Fe3 Al/steel diffusion welding interface at different heating temperatures are shown in Table 5.18. As the heating temperature of diffusion welding increases, the diffusion rate of elements in the Fe3 Al interface transition zone increases rapidly, the number of atoms undergoing diffusion migration increases due to the larger diffusion driving force obtained by the elements. The expression for the width of the Fe3 Al/Q235 steel diffusion welded interfacial transition zone is shown as below, which is calculated from the diffusion rates of the elements at different temperatures

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5 Joining of Iron-Aluminium Intermetallic Compounds

Table 5.18 Diffusion rates of elements at the Fe3 Al/steel diffusion welding interface at different heating temperatures Joints

Fe3 Al/Q235

Fe3 Al/18–8

Heating temperature/°C

1040

1060

1080

1000

1020

1040

1060

Diffusion ratio (K p ) /μm2 ·s−1

Al

1.2

7.7

17.1

0.98

1.0

3.9

9.1

Fe

1.9

4.9

14.5

0.08

0.44

2

2.4

Cr

–

–

–

0.34

0.85

0.98

2.5

Ni

–

–

–

0.78

1.0

1.6

2.1

x 2 = 4.8 × 104 exp(−

133020 )(t − t0) RT

(5.5)

The expression for the width of the transition zone at the diffusion weld interface of Fe3 Al/18-8 steel is x 2 = 7.5 × 102 exp(−

75200 )(t − t0) RT

(5.6)

The width of the interfacial transition zone of diffusion welding of Fe3 Al/Q235 steel and Fe3 Al/18-8 steel is mainly related to the heating temperature T and holding time t. As the heating temperature increases and the holding time increases, the width x of the interfacial transition zone gradually increases, which is conducive to promoting the bonding of the diffusion welded interface. The calculated and measured values of the width of the interfacial transition zone of Fe3 Al/18-8 steel are shown in Fig. 5.40. It can be seen that, under the given test conditions, a certain width of the interfacial transition zone of Fe3 Al/Q235 steel and Fe3 Al/18-8 steel diffusion welding can be obtained according to the determined heating temperature and holding time by using the relationship between the heating temperature and holding time, the bonding performance of the Fe3 Al/steel diffusion welding interface can be improved. The formation of the reaction layer in the interfacial transition zone of Fe3 Al/steel diffusion welding has a certain latent time t 0 . When the width of the interfacial transition zone is certain, the latent time t 0 shortens with the increase of the heating temperature T. Therefore, during determining the Fe3 Al/steel diffusion welding process parameters, in the condition of ensuring the appropriate width of the interface transition zone, increasing the heating temperature T at the same time can be appropriately shortened holding time t, that could improve welding efficiency.

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5.3 Vacuum Diffusion Welding of Fe3 Al to Steel (Q235, 18-8 Steel)

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Fig. 5.40 Comparison of calculated and measured values of transition zone width at the interface of Fe3 Al/18-8 steel

5.3.6 Effect of Process Parameters on the Interface Characteristics of Diffusion Welding (1) Three important parameters Diffusion welding parameters (heating temperature T, holding time t and pressure p) have an important influence on the bonding condition at the diffusion interface of Fe3 Al/Q235 steel. (1) Heating temperature The higher the heating temperature, the higher the energy gained by the atoms of the elements near the interface and the faster the diffusion rate. With the driving force of the concentration gradient, the elements in the base material will diffuse rapidly towards the interface. According to the analysis of the bonding characteristics of and degree of deformation Fe3 Al/Q235 steel and Fe3 Al/18-8 steel joint obtained under different process conditions, the diffusion interface bonding of Fe3 Al/Q235 steel and Fe3 Al/18-8 steel is gradually tightened with the increase of heating temperature. When heated to a certain temperature, a transition zone is formed near the Fe3 Al/Q235 steel and Fe3 Al/18-8 steel diffusion interface. (2) Holding time. The holding time determines the degree of homogenization of atomic diffusion near the interface of diffusion welding of Fe3 Al/Q235 steel and Fe3 Al/18-8 steel. As the holding time increases, the elements near the diffusion welding interface continue to diffuse to the interface, the distribution of elements becomes more and more uniform, and the width of the interface transition zone gradually increases and homogenizes.

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5 Joining of Iron-Aluminium Intermetallic Compounds

(3) Connection pressure. Connection pressure is the main factor to ensure the disappearance of interfacial microvoids and the degree of deformation in Fe3 Al/Q235 steel and Fe3 Al/18-8 steel diffusion welding joints. Under the condition of constant heating temperature and holding time, the higher the pressure, the larger the area of close contact at the diffusion interface, the interface microscopic voids are easy to disappear and gradually form a dense diffusion welding interface. When the pressure decreases, the contact area of the diffusion interface is smaller, and the interfacial microscopic cavities prevent the atoms on both sides from crossing the interface for diffusion migration, forming a non-dense diffusion welding interface or even inadequate interfacial bonding. However, when the pressure is too high, it will lead to significant plastic deformation of the diffusion welded joint. The pressure is generally determined according to the contact area of the weld to ensure that no macroscopic deformation occurs. The diffusion rate of different elements is different due to differences in the physical and chemical properties of the base material during Fe3 Al and Q235 steel (or 18-8 steel) diffusion welding, and the number of atoms is not equal through the interface to both sides of the base material, resulting in Kirkendall (Kirkendall) effect and the formation of microscopic cavities at the diffusion interface. In a certain heating temperature and holding time, these microscopic cavities gradually disappear, forming a dense diffusion welding interface, so the existence of microscopic cavities can be used as one of the important indicators to evaluate the bonding performance of the diffusion welding interface. Under pressure, diffusion welding joints will also produce certain macroscopic deformation due to the changes of high temperature performance, which affects the properties of the joint. (2) Effect of parameters on interfacial bonding and joint deformation (1) Fe3 Al/Q235 steel diffusion welding. The test shows that under the conditions of constant holding time and connection pressure (t = 60 min, p = 17.5 MPa), the Fe3 Al/Q235 steel interface does not form a sufficient diffusion bond at a heating temperature of 1000 °C, and a large number of microscopic voids can be observed under the microscope, see Fig. 5.41a; when the heating temperature is increased to 1020 °C, the Fe3 Al/Q235 steel interface was partially bonded, and microscopic cavities could still be observed under the microscope. When the heating temperature was 1040 °C, the microscopic cavities at the Fe3 Al/Q235 interface completely disappeared, the interface was well bonded, and a diffusion transition zone was formed near the Fe3 Al/Q235 steel interface, see Fig. 5.41b. When the heating temperature continued to increase to 1060 °C, no microscopic voids were observed at the Fe3 Al/Q235 steel interface, and the width of the interfacial transition zone increased, but slight plastic deformation occurred in the diffusion joint. The degree of macroscopic deformation of the diffusion joint gradually increased when the heating temperature was increased to 1080 °C.

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5.3 Vacuum Diffusion Welding of Fe3 Al to Steel (Q235, 18-8 Steel)

283

(b)

(a) Q235 Fe3Al

Q235 Micro-void 12.5µm T=1000℃

Fe3Al

20µm T=1040℃

Fig. 5.41 Bonding morphology of Fe3 Al/Q235 steel diffusion interface at different heating temperatures. (t = 60 min, p = 17.5 MPa)

No macroscopic deformation occurred in Fe3 Al/Q235 steel diffusion joints at heating temperature T = 1040 °C and pressure p = 17.5 MPa with holding time t = 15–60 min. And with the increase of holding time t, the interface bonding of Fe3 Al/Q235 was gradually tightened. Under the conditions of constant heating temperature and holding time (T = 1060 °C, t = 45 min) and welding pressure of 10 MPa, microscopic cavities could be seen in the Fe3 Al/Q235 interface. As the welding pressure increased from 12 to 17.5 MPa, the Fe3 Al/Q235 interface bonding gradually tightened. However, slight macroscopic deformation of the Fe3 Al/Q235 steel diffusion joint occurred at the pressure p = 17.5 MPa. Therefore, Fe3 Al/Q235 steel diffusion welding should be controlled at a certain heating temperature and holding time, and the welding pressure should not be too high to avoid macroscopic deformation of the diffusion joint. (2) Diffusion welding of Fe3 Al/18-8 steel. To obtain Fe3 Al/18-8 steel diffusion joints with good interfacial bonding, a series of tests were carried out using different heating temperatures T, holding times t and pressures p. The test results showed that under the conditions of constant holding time and pressure (t = 60 min, p = 17.5 MPa) and lower heating temperature (980 °C), there was a continuous distribution of microscopic voids at the Fe3 Al/18–8 interface and no good diffusion bond was formed, see Fig. 5.42a. When the heating temperature was increased to 1000 °C, the Fe3 Al/18-8 steel was partially bonded at the interface, and microscopic voids were still observed locally under the microscope. At a heating temperature of 1020 °C, the Fe3 Al/18-8 steel interface diffusion bonded well and a diffusion transition zone formed near the interface, see Fig. 5.42b. When he heating temperature continued to increase and the Fe3 Al/18-8 interface diffusion bonded more fully, but a slight plastic deformation occurred in the diffusion welding joint at a heating temperature of 1060 °C. More plastic deformation of the diffusion joint occurs at a heating temperature of 1080 °C.

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5 Joining of Iron-Aluminium Intermetallic Compounds

(a)

(b)

18-8

18-8

Fe3Al

Fe3Al 25µm

50µm T= 980℃

T= 1020℃

Fig. 5.42 Bonding morphology at the Fe3 Al/18–8 steel diffusion welding interface at different heating temperatures. (t = 60 min, p = 17.5 MPa)

Under the conditions of constant heating temperature and pressure (T = 1040 °C, p = 17.5 MPa), the Fe3 Al/18–8 diffusion bonding was gradually tightened with the increase of holding time. When the heating temperature and holding time are kept constant (T = 1060 °C, t = 45 min), the higher the pressure, the tighter the Fe3 Al/188 steel diffusion bonding is, and the microscopic voids in the transition zone are gradually reduced. However, at a pressure p greater than 17.5 MPa, Fe3 Al/18-8 steel diffusion welding joints underwent significant plastic deformation. The heating temperature, holding time and connection pressure determine the quality of the Fe3 Al/Q235 steel and Fe3 Al/18-8 steel diffusion welding joint. When the heating temperature increase, the holding time and welding pressure can be reduced accordingly; under the condition of ensuring that no macroscopic deformation occurs in the diffusion welding joint (the pressure is controlled within a certain range), the heating temperature can be reduced accordingly when extending the holding time. Therefore, the influence of process parameters on the microstructure properties of the interface of Fe3 Al/Q235 steel and Fe3 Al/18-8 steel diffusion welding should be taken into account. To obtain well-bonded Fe3 Al/steel diffusion welding joints, the optimum match of heating temperature, holding time and pressure should be determined by test. (3) Effect of parameters on the width of the transition zone at the Fe3 Al/steel interface (1) Effect of heating temperature. As the heating temperature increases, the elements diffuse fully, the width of the transition zone of Fe3 Al/Q235 steel and Fe3 Al/18-8 steel at the diffusion joint increases gradually, and the organization of the diffusion transition zone coarsens gradually. The measured width of the diffusion transition zone for Fe3 Al/Q235 and Fe3 Al/18-8 steels at the same holding time (t = 60 min) and different heating temperatures are listed in Table 5.19.

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Table 5.19 Width of transition zone at the interface of Fe3 Al/Q235 steel and Fe3 Al/18-8 steel at different heating temperatures (t = 60 min) Heating temperature/°C

1000

1020

1040

1060

1080

Width/μm

Fe3 Al/Q235

–

22.1

24.5

28.6

32.5

Fe3 Al/18–8

22.6

26.3

35.4

38.2

42.6

As seen in Table 5.19, the width of the interface transition zone of Fe3 Al/Q235 steel diffusion welding joint at a heating temperature of 1020 °C is 22.1 μm, and the width of the interface transition zone of Fe3 Al/18-8 steel is 26.3 μm. When the heating temperature is increased to 1080 °C, the width of the interface transition zone of Fe3 Al/Q235 joint increases to 32.5 μm, and the width of the interface transition zone of Fe3 Al/18-8 joint increases to 42.6 μm. The effect of heating temperature on the width of the diffusion interfacial transition zone obtained from the measured results is shown in Fig. 5.43. Based on the measured results, it is foreseen that the width of the transition zone at the diffusion interface of Fe3 Al/Q235 steel and Fe3 Al/18-8 steel will increase by continuing to increase the heating temperature. However, too high heating temperature will lead to a significant coarsening of the organization near the diffusion welding interface, which has a negative impact on the organization and mechanical properties of the diffusion joint. Therefore, the heating temperature should be limited. (2) Effect of holding time and pressure. Holding time and pressure are the main factors that determine the uniformity of element diffusion near the diffusion interface and the disappearance of microscopic voids. The width of the transition zone at the Fe3 Al/Q235 steel diffusion interface for a heating temperature of 1060 °C and different holding times and welding pressures is shown in Fig. 5.44.

Fig. 5.43 Effect of heating temperature on the width of the transition zone at the Fe3 Al/steel diffusion interface

286

5 Joining of Iron-Aluminium Intermetallic Compounds (b)

(a)

Q235

Fe3Al

Fe3Al

Q235

1060 ℃X30min, P=10MPa

1060℃X60min, P=12MPa

Fig. 5.44 Microstructure of the transition zone at the Fe3 Al/Q235 diffusion interface at different holding times

At a holding time of 30 min, even at a higher temperature (T = 1060 °C), unresolved microscopic voids were observed in the transition zone of the Fe3 Al/Q235 interface; at a heating temperature of 1040 °C and a shorter holding time (t = 30 min), there were obvious microscopic voids and insufficient diffusion of elements at the junction of the Fe3 Al/18-8 steel mixed transition zone and the interfacial transition zone near the 18-8 steel side. This is due to the short holding time and small pressure, the microscopic contact area of Fe3 Al/Q235 steel and Fe3 Al/18-8 steel diffusion welding interface is small, interfacial microscopic cavities hinder grain growth and diffusion migration of atoms across the interface, the atoms can not diffuse or diffusion is not sufficient, the formation of interface transition zone is narrow. The longer the holding time and the higher the pressure, the larger the close contact area of the diffusion welding interface, and the microscopic cavities near the interface will gradually disappear to form a dense diffusion welding interface transition zone. As the longer the holding time, the elements diffuse more fully, the more intense the mutual diffusion migration between atoms. The measured width of the interface transition zone for Fe3 Al/Q235 and Fe3 Al/18-8 diffusion joint at different holding times are shown in Table 5.20. The effect of holding time on the width of the interface transition zone for Fe3 Al/Q235 and Fe3 Al/18-8 steel diffusion welding based on the measured results is shown in Fig. 5.45. With the increase of holding time, the width of the interfacial transition zone of Fe3 Al/Q235 steel and Fe3 Al/18-8 steel diffusion welding joint gradually increased. Table 5.20 Width of transition zone at the interface of diffusion welding of Fe3 Al/Q235 steel and Fe3 Al/18-8 steel at different holding times Holding time/min Width/μm

15

30

45

60

80

Fe3 Al/Q235 (T = 1060 °C)

–

17.4

25.8

28.6

30.4

Fe3 Al/18-8 (T = 1040°C)

12.3

20.1

28.2

35.1

38.5

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5.3 Vacuum Diffusion Welding of Fe3 Al to Steel (Q235, 18-8 Steel)

287

Fig. 5.45 Effect of holding time on the width of the transition zone at the Fe3 Al/18-8 steel diffusion weld interface

The width of the interfacial transition zone increases faster when the holding time is less than 45 min; however, the width of the interfacial transition zone increases more slowly when it exceeds 45 min. At the initial stage of insulation, element diffusion is influenced obviously by the insulation time, the longer the insulation time, the more fully the diffusion of elements migrate. When a certain time is reached, the diffusion of elements by the influence of holding time decreases, and the diffusion migration of elements gradually reaches a quasi-equilibrium state, and a diffusion welding interface transition zone with a stable structure is formed near the interface of Fe3 Al/Q235 steel and Fe3 Al/18-8 steel joint. If the holding time is too long, the organization near the diffusion welding interface will grow, and the microstructure will be significantly coarsened and affect its mechanical properties. Therefore, holding time of Fe3 Al/Q235 steel and Fe3 Al/18-8 steel diffusion welding should be strictly controlled, not only to ensure that the diffusion welding interface transition zone has a certain width, but also make the organization not significantly coarsened. Heating temperature, holding time and pressure interact and affect together the organization properties of the interface transition zone of Fe3 Al/Q235 steel and Fe3 Al/18-8 steel diffusion joint. In order to obtain a well-bonded interface with adequate atomic diffusion and good properties of Fe3 Al/Q235 steel and Fe3 Al/188 steel diffusion joints, the heating temperature, holding time and joining pressure must be controlled in a coordinated manner.

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5 Joining of Iron-Aluminium Intermetallic Compounds

5.4 Other Welding Methods of Fe3 Al Intermetallic Compounds 5.4.1 Electron Beam Welding of Fe3 Al Intermetallic Compounds The poor fusion weldability of Fe3 Al intermetallic compound is evident in two ways. One is that Fe3 Al intermetallic compounds lead to high stress concentrations due to difficulties in cross-slip migration, resulting in high room temperature brittleness, low plasticity and susceptibility to cold cracking during fusion welding. Secondly, the low thermal conductivity of Fe3 Al leads to a large temperature gradient between the heat affected zone, the fusion zone and the weld, coupled with a large coefficient of linear expansion, which tends to produce large residual stresses during cooling, leading to thermal cracking. Electron beam welding uses an electron beam produced by an electron gun to focus on the workpiece, causing the weld metal to melt rapidly and then re-solidify and crystallize. The chemical composition and process parameters have a significant effect on the weldability of Fe3 Al. Vacuum electron beam welding (EBW) on Fe3 Al intermetallic compound sheets of 0.76 mm thickness showed that no delayed cracking occurred after welding as the welding process was carried out in vacuum, which suppressed the harmful effects of hydrogen. And the concentrated high energy input refines the organization of the weld fusion zone, the weld organization is columnar crystal with narrow width, growing along the heat conduction direction, and the heat affected zone is also very narrow, no cracks are generated at lower welding speed and the joint deformation is also small. Therefore, the Fe3 Al-based alloy rich in Cr, Nb and Mn does not show cracks after welding and the quality of the welding joints obtained is better. Good Fe3 Al welding joints can be obtained by controlling the welding speed below 20 mm/s when electron beam welding is used. The mechanical properties tests showed that the fracture occurred in the heat affected zone and the tensile fracture was a mixture of along-crystal and through-crystal destructive fracture, which was the same fracture mechanism as that of the base material before welding. It can be seen that although the heat input is concentrated in electron beam welding, the joint is still affected by the intrinsic brittleness of the Fe3 Al base material and shows brittle fracture characteristics. Vacuum electron beam welding of Fe3 Al-based alloys with a thickness of 1–2 mm, the process parameters used are: focusing current of 800–1200 mA, welding current of 20–30 mA, welding speed of 8.3–20 mm/s, vacuum of 1.33 × 10–2 Pa. The welding effect is better than tungsten arc welding because of the concentrated energy and the low concentration of H and O atoms in the vacuum atmosphere, which suppresses the action of hydrogen and makes it difficult for hydrogen delayed cracking to occur in the welded joint, and a crack and defect-free weld can be obtained with a narrow

5.4 Other Welding Methods of Fe3 Al Intermetallic Compounds

289

weld (about half that of tungsten arc welding) and a narrow heat-affected zone, with small post-weld deformation and less stress. Tensile and bending tests of electron beam welded Fe3 Al based alloy showed that fracture occurred in the heat affected zone of Fe3 Al base material at room temperature with a tensile strength of 289 MPa and the weld did not weaken the mechanical properties of the joint zone. Thus, using electron beam welding, the Fe3 Al-based alloy exhibited good weldability with an aesthetically pleasing weld profile and excellent performance. Vacuum electron beam welding of thin plate Fe3 Al-based alloy is fast and can be controlled in the range of 4.2–16.9 mm/s with high welding efficiency, which has good application prospects.

5.4.2 Electrode Arc Welding of Fe3 Al (1) Process parameters The butt weld tests of Fe3 Al/Fe3 Al, Fe3 Al/Q235 steel and Fe3 Al/18-8 steel were carried out using electrode arc welding (SMAW) without preheating and post-weld heat treatment. The welding (SMAW) was carried out using a Norwegian Master TIG MLS2500 welder and the available welding materials were four types of electrodes E308-16, E316-16, E309-16 and E310-16 with 2.5 and 3.2 mm diameters and the chemical composition and mechanical properties are shown in Table 5.21. The process parameters used for electrode arc welding (SMAW) are shown in Table 5.22. Table 5.21 Chemical composition and mechanical properties of weld materials Electrode Chemical composition of weld/wt.% type C Cr Ni Mn

Mechanical property Mo

0.5–2.5 ≤0.75

Si

Tensile Elongation strength δ 5 /% σb /MPa

E308-16

≤0.08

18.0–21.0 9.0–11.0

≤0.90 ≥550

≥35

E316-16

≤0.08

17.0–20.0 11.0–14.0 0.5–2.5 2.0–3.0 ≤0.90 ≥520

≥30

E309-16

≤0.15

22.0–25.0 12.0–14.0 0.5–2.5 ≤0.75

≤0.90 ≥550

≥25

E310-16

0.08–0.20 25.0–28.0 20.0–22.5 1.0–2.5 ≤0.75

≤0.75 ≥550

≥25

Table 5.22 Process parameters for electro-arc welding of Fe3 Al electrodes Diameter Φ/mm

Current I/A

Voltage U/V

Speed v/cm·s−1

Heat input E/kJ·cm−1 (η = 0.85)

2.5

100–120

24–26

0.20–0.30

8.8–13.3

3.2

125–140

24–27

0.25–0.35

9.2–12.9

290

5 Joining of Iron-Aluminium Intermetallic Compounds

(a)

(b) WM

PMZ

Fe3Al

25μm 250

500

Fig. 5.46 Microstructure of Fe3 Al/Q235 steel electrode arc weld

Too large or too small of welding heat input is easy to cause welding cracks. When the welding heat input is too small, the weld cools quickly and produces obvious surface cracks after welding. When the welding heat input is too large, overheating time of the fusion pool is longer, resulting in weld organization coarsening to induce cracks. In particular, the electrode arc welding slag adheres to the deposited metal resulting in slow heat dissipation and coarsening of the organization. The test results show that the electrode arc welding (SMAW) using E310-16 type welding rod as welding material and controlled welding heat input can obtain crack-free Fe3 Al joints. Microstructure of Fe3 Al/Q235 steel electrode arc weld is shown in Fig. 5.46. (2) Electrode arc overlay welding Electrode arc overlay welding can give the surface of the part wear resistance, corrosion resistance, heat resistance and other special properties. In petrochemical and thermal processing production, a large number of stainless steel overlay heat-resistant steel structure are used. If Fe-Al alloy can be used instead of stainless steel as the overlayer, such as Fe3 Al alloy was electrode arc welded (SMAW) on austenitic stainless steel, 2.25Cr-1Mo steel or other steel, its excellent performance can be achieved. The Fe3 Al alloy melted by the medium frequency induction furnace is cast into ingots, which are hot rolled and hot forged in multiple passes (temperature controlled above 900 °C) to make a 3.2 mm diameter bar, which is used as the core of the electrode. Low hydrogen potassium type flux is selected, and the composition of the core and the composition of the overlay metal is shown in Table 5.23.

Table 5.23 Fe3 Al weld core composition and % composition of overlay metal Materials

Al

Cr

Fe

Ni

Ti

Si

Fe3 Al core

16.00

5.10

78.70

–

–

0.20

Overlay

11.60

5.95

70.69

0.56

0.20

1.00

5.4 Other Welding Methods of Fe3 Al Intermetallic Compounds

291

To ensure stable composition, at least three layers should be overlaid. Using DC arc welding machine, the overlay voltage is about 25 V, the overlay current is taken as the lower limit (generally 90–110A), and the overlay electrode moves at a speed of about 12 cm/min. The spatter during overlay welding is small, but the deslagging is poor, and the residue should be carefully removed when overlaying the next layer, so that a crack-free Fe3 Al overlayer can be obtained. The workpiece is preheated to 300–350 °C before overlay welding and held for 30 min, and the weld is annealed at 700 °C × 1 h after overlay welding. The metal of the overlayer is dominated by coarse columnar crystals, and the Al content of the overlayer is lost during welding, resulting in the organization of the overlayer being dominated by α-Fe(Al) solid solution, but it does not affect the oxidation resistance of the overlayer. After oxidation at 800 °C × 70 h in the air furnace, the stainless steel matrix oxidized severely, while the Fe3 Al overlayer oxidized slightly, indicating that its high-temperature oxidation resistance is better than that of 18-8 stainless steel.

5.4.3 Argon Arc Overlay Welding and Characteristics of Fe3 Al Most of this overlay welding process is done by wire-filled tungsten arc welding with alloy filler material, which can be manipulated automatically or done manually. Surface oil and rust of 2.25Cr-1Mo steel plate of 40 × 20 × 6 mm is removal, and Fe3 Al alloy (Fe 84%, Al 16%) was overlay welded on 2.25Cr-1Mo heat-resistant steel using filler tungsten arc welding (GTAW) method, with welding current of 75A. The heat-resistant steel workpiece is preheated by 300 °C, and a post heat treatment of 600 °C × 1 h is carried out. The key to the technology is to choose an alloy wire that can form a large amount of Fe3 Al intermetallic compound in the overlay layer. Under such conditions, the overlayer has high hardness and wear resistance, corrosion resistance, although there may be some micro-cracks in the overlayer. As observed by scanning electron microscopy (SEM), the interface between the Fe3 Al overlayer and the 2.25Cr-1Mo heat-resistant steel was well bonded, forming a fusion zone of the overlayer with a width of about 300 μm. The organization of the overlayer was coarse columnar crystal, with a large number of needles distributed within each columnar crystal. By electron probe analysis, these needles contain a large amount of Fe and Al, which constitute α-Fe(Al) solid solution. The fusion zone is the weakest zone of the Fe3 Al and 2.25Cr-1Mo overlay joint, and the energy spectrum analysis of the chemical composition of the fusion zone of the Fe3 Al and 2.25Cr-1Mo overlay joint is shown in Table 5.24. The concentration gradient of the alloying elements Cr, Mo, and Al near the fusion zone between the Fe3 Al overlayer and the 2.25Cr-1Mo matrix varies more significantly. The Al in the overlay metal is heavily diluted, leading to low Al content in the overlayer, and a single-phase α-Fe(Al) solid solution is formed.

292 Table 5.24 Energy spectrum analysis of the chemical composition of the fusion zone of Fe3 Al and 2.25Cr-1Mo overlay joints

5 Joining of Iron-Aluminium Intermetallic Compounds Location

Al

Cr

Mo

1

1.07

2.18

1.29

2

1.22

2.42

1.21

3

2.02

2.05

0.85

4

3.04

2.01

0.97

5

3.31

1.85

0.94

Overlay

8.15

1.08

0.43

Matrix

–

2.43

1.19

Note The first 5 positions in the table are the measurement points taken at 100 μm intervals from the fusion line, respectively

Bibliography 1. Li Y, Ma H, Wang J (2011) A study of crack and fracture on the welding joint of Fe3 Al and Cr18-Ni8 stainless steel. Mater Sci Eng A 528:4343–4347 2. Ma H, Li Y, Puchkov UA, et al (2008) Microstructural characterization of welded zone for Fe3 Al/Q235 fusion-bonded joint. Mater Chem Phys (112):810–815 3. Mota MA, Coelho AA, Bejarano JM et al (1999) Directional growth and characterization of Fe-Al-Nb eutectic alloys. J Cryst Growth 198–199(1):850–855 4. Huang YD, Yang WY, Sun ZQ (2001) Effect of the alloying element chromium on the room temperature ductility of Fe3 Al intermetallics. Intermetallics 9:119–124 5. Ma HJ, Li YJ, Girasimov KL, et al (2007) Study of B2-D03 ordered structure transition mode of Fe3 Al intermetallic compounds under welding conditions. Chin J Nonferrous Metals 17(S1):25–29 6. McKamey CG, Horton JA (1989) Effect of chromium on properties of Fe3 Al. J Mater Res 4(5):1156–1163 7. Ding C, Chen C, Guozhi LS, et al (2000) Study on the organization and properties of Fe-Al alloy TIG welded heads. J Appl Sci 18(1):368–370. 8. Ma H, Li Y, Li J, et al (2007) Division of character zones and elements distribution of Fe3 Al/Cr-Ni alloy fusion-bonded joint. Mater Sci Technol 23(7):799–812 9. Ma H, Li Y, Juan W, et al (2006) Effect of heat treatment on microstructure near diffusion bonding interface of Fe3 Al/18-8 stainless steel. Mater Sci Technol 22(12):1499–1502 10. Senying L, Leiro JA (1999) Cr impurity effect on antiphase boundary in FeAl alloy. J Appl Phys 38(5):2806–2811 11. Joslin DL, Easton DS, Liu CT et al (1995) Reaction synthesis of Fe-Al alloys. Mater Sci Eng 192A(2):544–548 12. Ma H, Li YJ, Wang J et al (2006) Effect of reheating on the organization of Fe3 Al/18-8 diffusion welding near the interface. J Weld 27(5):35–38 13. Wang J, Li Y, Ma H (2006) Diffusion bonding of Fe-28Al(Cr)aloy with low-carbon steel in vacuum. Vacuum 80(5):426–431 14. Li Y, Wang J, Puchkov UA, et al (2007) Effect of Cr and Ni elements on the organization of the Fe3 Al/steel diffusion welding interface. Mater Sci Technol 15(4):470–475 15. Li Y, Gerasimov SA, Puchkov UA et al (2007) Microstructure performance on TIG welding zone of Fe3 Al and 18–8 dissimilar materials. Mater Res Innov 11(3):45–47 16. Wang J, Li Y, Liu P (2004) Microstructure and performance in diffusion-welded joints of Fe3 Al/Q235 carbon steel. J Mater Process Technol 145(3):294–298 17. Wang J, Li Y, Liu P (2003) Shear strength and tissue properties of Fe3 Al/Q235 diffusion welded joints. J Weld 24(5):81–84

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18. Li Y, Wang J, Yin Y, et al (2002) Phase structure analysis of Fe3 Al/Q235 diffusion welding interface of dissimilar materials, J Weld 23(2):25–28 19. Min XG, Yu XQ, Sun YS et al (2001) Histomorphology and oxidation resistance properties of hand-arc overlay welding Fe3 Al overlay layer. J Weld 22(1):56–58 20. Ma H, Li Y, Gerasimov SA, et al (2007) Microstructure and phase constituents near the fusion zone of Fe3 Al/Cr-Ni alloys joints produced by MAW. Mater Chem Phys 103(1):195–199

Chapter 6

Welding of Laminated Materials

The laminated material is a new type of material developed in recent years, which has received attention from Europe, America, Russia and other countries for its unique high temperature resistance and corrosion resistance performance. The thinner SuperNi cladding on the surface of the NiCr base layer can inhibit the expansion of microcracks in the NiCr base layer, prevent instantaneous damage when microcracks and defects exist in the structural parts, and improve the overall load-bearing capacity of the laminated material. As the Super-Ni laminated material (Super-Ni/NiCr laminated material) has the advantages of low density, corrosion resistance and high temperature resistance, it has a broad application prospect in aerospace, energy power and other fields, and the welding problem of laminated material is also getting more and more attention.

6.1 Characteristics and Weldability of Laminated Materials 6.1.1 Characteristics of Laminated Materials Super-Ni/NiCr laminated composite material is a new structural material developed in recent years, consisting of Super-Ni laminated on both sides and NiCr (or intermetallic compound) laminated in the middle, which is similar to a “sandwich” structure. The so-called Super-Ni means that the purity of the compound layer exceeds the Ni content level specified in the national standard. Super-Ni has better oxidation resistance, corrosion resistance, plasticity and toughness, which can be used as corrosion resistant structural parts under high temperature.

© Chemical Industry Press 2023 Y. Li, Joining Technology and Application of Advanced Materials, Advanced and Intelligent Manufacturing in China, https://doi.org/10.1007/978-981-19-9689-4_6

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The NiCr base layer of the laminated material is a porous material made of sintered Ni80Cr20 powder with a porosity of about 30–35%, which can reduce the structural mass. Porous, semi-dense or fully dense materials can be prepared by powder metallurgy technology now. Porous materials have many excellent properties, such as light weight, high specific stiffness, high specific strength, shock resistance, sound insulation, thermal insulation and so on. However, the porous material is rarely used as structural component directly, due to the reason that the structure is easy to deform and there are defects in the hole and the surface. The porous material is usually used to forming a composite structure with solid materials in order to play its unique material properties. Porous metal materials can be used as the core in the rigid sandwich composite structures, which have received attentions in the fields of aviation, aerospace, missile, and aircraft design. Super-Ni laminated material is a new type of high temperature structural material, composed of Super-Ni cover layer and Ni80Cr20 powder alloy base layer by vacuum pressing. The thickness of the cover layer of this laminated material is only 0.2~0.3 mm, Ni>99.5%. And the base layer is NiCr alloy with a thickness of about 2.0~2.6 mm (Ni content is 80%, Cr content is 20%). The earlier high-temperature alloy is the pressed nickel-based high-temperature alloy Nimonic80A developed on the basis of 80% Ni-20% Cr alloy, by adding a small amount of Ti and Al elements to improve the creep fracture strength and hightemperature oxidation resistance. NiCr alloy is often used as a substrate for other high-temperature alloys. Nickel-based high-temperature alloys are widely used in the aerospace field, especially for heated end components in turbine engines, such as combustion chambers, turbine blades, etc. The Super-Ni cover layer has excellent corrosion resistance, oxidation resistance and toughness, while the Ni80Cr20 powder is vacuum sintered to form a porous material with lower density. That can reduce the structural mass and improve the overall performance of the parts. Super-Ni/NiCr laminated composite can take full use of performance advantages of Super-Ni and NiCr base layers, which is better than single materials for certain applications. (1) Chemical composition and porosity of NiCr base layer There is a big difference of physical property parameters between Super-Ni/NiCr laminated composites and the conventional materials. The elemental content in the Ni80Cr20 base layer was determined using plasma emission spectroscopy, and the results are shown in Table 6.1. Table 6.1 Chemical composition of the laminated composite base layer Element

Ni

Cr

Fe

Mo

Al

Co

Wavelength/nm

231.6

284.3

259.9

204.5

167.0

231.1

Average content/%

64.86

17.44

0.4343

0.0467

0.0387

0.0133

6.1 Characteristics and Weldability of Laminated Materials

297

From Table 6.1, it can be seen that Ni and Cr are the main elements of the base layer and Fe, Co, Mo and Al are the trace elements. The microstructure characteristic of the laminated composites is shown in Fig. 6.1. Microstructure of the base layer is homogeneous and interfaces between the skeletal structures in white and the porosities in black are clear. Ni80Cr20 base layer is powder sintered alloy. The nominal porosity and density are important parameters reflecting the material properties. The nominal porosity and nominal density of the Ni80Cr20 base layer were measured using the area method according to the principle of bulk optics. Measuring the cross-sectional area of the pores AP , and the total area of the observation scope A. The area percentage of the pores is calculated according to Eq. (6.1), and the nominal porosity ε of the porous material can be calculated according to the theory of bulk optics (whose volume percentage is equal to the cross-sectional area percentage). ε=

Ap × 100% A

(6.1)

The nominal porosity of Ni80Cr20 alloy base was calculated to be 35.41% and the nominal density was 6.72 g/cm3 . (2) Structural characteristics of laminated material The laminated material is made of two materials with different properties by vacuum pressing or other special preparation methods. The laminated material has the advantages of each component obtaining the unique physical and chemical properties. At present, the United States, Russia, the United Kingdom, Germany and other developed countries have achieved remarkable achievements in the research and application of laminated materials. Our country started the relevant research in the 1960s. Important progress in the research and production has been made in recent years.

(a)

(b)

Ni cover layer

Ni80Cr20

400μm

Super-Ni cover layer/NiCr base layer Fig. 6.1 Microstructure of Super-Ni laminated composites

NiCr base layer

298

6 Welding of Laminated Materials

Figure 6.2 shows the service temperature of alloys and the percentage of service temperatures against the melting points. Materials for advanced aero-engines often serve at temperatures above 85% of their melting point and under high load conditions, put forward higher requirement on the high temperature properties of the materials. As can be seen from Fig. 6.2, the combination of two materials with different high temperature resistance and mechanical properties can give full play to the advantages of good high temperature resistance and mechanical properties of both materials to better meet the needs of special service environments. Composite materials can be classified as laminated composites, particle reinforced composites and fiber reinforced composites. Lamellar composites are composed with two or more materials with different properties by special processing methods. Therefore, the composites obtain the advantages of different components. The laminated composites can be divided into two types: laminated composites and micro-laminated composites, as shown in Table 6.2. Laminated composites are in the “sandwich” structure of cover layer + base layer + cover layer, and the cover layer is thin, generally less than 0.4 mm. The base layer mainly meets the requirements of structural strength and stiffness, and the cover layer meets the requirements of corrosion resistance, wear resistance and other special properties. The micro-laminated composite material is composed of two or three materials alternately, which is related to the preparation process of micro-laminated material, and the layer thickness of micro-laminated material is 100~300 μm.

Fig. 6.2 Service temperatures of alloys and the percentage of service temperatures against the melting points

6.1 Characteristics and Weldability of Laminated Materials Table 6.2 Classification of laminated composites

299

Lamellar composites

Structure

Layer thickness

Laminated composites

“Sandwich”

Cover layer 1 mm). The main problems during fusion welding of laminated materials include microcracks in the weld and fusion zone, burnout of Ni cover layer, fusion defects in NiCr base layer (e.g., incomplete fusion, microscopic holes and cracks) and so on. (1) Microcracks in the weld area The most prominent problem in fusion welding of laminated materials is cracking. The main problem in welds is the generation of thermal cracks, and cracking due to stress concentration during welding. The thermal expansion and contraction caused by the welding thermal cycle, is easy to induce micro cracks and their propagation at the poorly bonded large grain boundaries in the weld fusion zone. The cracks terminate at the sintering holes in NiCr base layer, which can play a role in stopping crack propagation. During fusion welding of Super-Ni laminated material, stress state of the joint and the low melting point inclusions aggregation leading by the physical metallurgical reactions may induce crack generation. Usually when the weld solidifies, the S elements and other elements are easy to form metal sulfides (FeS, NiS, etc.) low melting point eutectic with Fe and Ni elements, which are easy to gather at large grain boundaries and become a source of cracking. To further analyze the possibility of forming low melting point sulfides, the content of C and S elements in the weld and base material was tested using a carbon and sulfur analyzer, as shown in Fig. 6.4. The content of C and S elements in the weld were lower than the statutory content for steel welding, possibility of low melting point sulfides that could lead to cracking was small. Super–Ni laminated composite was welded to with austenitic steel (1Cr18Ni9Ti) by wire-filled tungsten arc welding (GTAW). Crack morphology in the weld joint is shown in Fig. 6.5. Microstructure of the weld seam was columnar and its growth direction was perpendicular to the fusion zone. Alloying element and some low melting point impurity phases accumulated in the remaining liquid phase at the end of the columnar crystal. This region was a weak area in the weld. If there is restraint stress in the welding process, it is very easy to induce solidification cracks during the welding process.

302

6 Welding of Laminated Materials

0.4 Ni-Cr base metal Weld Preferred in Weld

0.342

Element content /%

0.3 0.25 0.2

0.1 0.061

0.035 0.0047 0.0043

0.0

Fig. 6.4 Histogram of C and S content in the weld and Ni–Cr base alloy

(a)

(b)

100μm

crack initiation

100μm

crack propagation

Fig. 6.5 Micro cracks in the weld seam

Figure 6.5a shows the cracks induced by weld stresses. The cracks started from the weld surface and extended to the interior of the weld. These cracks usually formed during the welding cooling process. The presence of such micro cracks indicated that there was great residual stress in Super-Ni/NiCr laminated composite weld joint. As a result, the crack initiated at the weld center and distributed along grain boundaries. Crack. Figure 6.5b shows the cracks distribution at the end of columnar crystal in the weld. The cracks were obvious with existence of low melting point inclusions. During the filler wire tungsten arc welding (GTAW) between Super-Ni laminated material and austenitic steel (18–8 steel), due to the cover layer structure of the laminated material and the different thermal-physical parameters of NiCr base layer,

6.1 Characteristics and Weldability of Laminated Materials

303

Ni cover layer, austenitic steel and 0Cr25-Ni13 filler wire, a complex stress state formed in the joint after welding. As the weld joint cooled to room temperature, the plasticity of the weld metal decreased, forming a tensile stress effect. Thus the weak area in the welded joint was prone to crack. Cracks mainly formed from the root or surface of the weld and extended to the center of the weld further. At the same time welding thermal cycle and uneven weld microstructure further aggravated the generation of residual stresses. When welded without cover layer (only the NiCr base layer was welded), thermal cracking induced by stress was significantly lower. So it is necessary to reduce weld stress in the welding of laminated materials. When Super-Ni laminated material was welded to austenitic steel GTAW, the sintered filler separated from the Ni80Cr20 base layer may enter the molten pool under the stirring action of the arc force and gather at the end of the columnar crystal during the cooling process, which may become cause of crack formation. The results of the energy spectrometer tests indicated that the crack-initiating inclusions contained mainly B, C, O, Cr, Fe, and Ni (Table 6.3), which may form carbides of Cr and metal oxides. With the growth of columnar austenite grains during the crystallization of the weld, impurity elements gathered at the end of the austenite columnar grains, forming a weak region of the weld metal. Therefore, welding process parameters should be controlled to reduce the effect of arc blowing force and control the fusion ratio of the base layer. (2) Burnout of Super-Ni cover layer Super-Ni cover layer was easy to be burnt in the welding process. Due to the thickness was only 0.2 to 0.3 mm, the influence of arc heat was great. The thinner cover layer was heated preferentially in the welding process, and its thermal conductivity (67.4 J/cm s °C) was much bigger than Ni80Cr20 base layer. Thus the cover layer melted rapidly and the surface of welding seam was widened. If the welding arc act for a longer time, even excessive burnout can occur (when the welding current was higher), resulting in poorly formed welded joints. Table 6.3 % Elemental content of each measurement point Position

B

C

O

S

Cl

Cr

Fe

Ni

Total

1

13.16

22.66

5.74

–

–

17.78

18.75

21.90

100

2

–

52.32

28.47

–

0.48

14.80

0.84

3.09

100

3

–

36.18

18.99

–

0.40

15.73

1.78

26.93

100

4

–

52.12

17.57

–

0.78

18.81

3.54

7.18

100

5

–

41.40

28.55

0.28

0.98

15.09

2.81

10.89

100

Max

13.16

52.32

28.55

0.28

0.98

18.81

18.75

26.93

–

Min

13.16

22.66

5.74

0.28

0.40

14.80

0.84

3.09

–

Note All the results are shown in weight percentage

304

6 Welding of Laminated Materials

It is difficult to avoid the burnout of cover layer in the fusion welding of super-Ni laminated composite due to the different thermo-physical properties of the super nickel cover layer and Ni80Cr20 base layer. Therefore, it is necessary to control the welding heat input (process parameters) strictly. The larger welding arc power and unstable welding process parameters are prone to cause super nickel cover layer burnout. (3) Fusion defects in base layer The welding behavior and fusion state of Ni80Cr20 base layer have an important influence on the microstructure and properties of the joint, which are important factors in the weldability analysis of Super-Ni laminated materials. The analysis revealed that partial fusion, fusion zone holes, and fusion zone microcracks were the main fusion defects during the welding of NiCr base layers. Analysis of the fusion zone of the laminated material and 18–8 steel wire-filled GTAW joints indicated that there was partial fusion in NiCr base layer as shown in Fig. 6.6a. Discontinuous fusion zone was easily formed under the improperly controlled welding parameters. Microstructure of fusion zone of the NiCr base layer consisted mainly of austenite, with ferrite precipitated at the grain boundaries. Sintering additives in NiCr base layer has significant effect on the welding seam formation, which is different from the conventional cast or rolled alloys. There was a small amount of sintering additives transitioning from NiCr base layer to the fusion zone, resulting in the discontinuous fusion zone morphology. A series of holes would form between the weld filler metal and base layer in the GTAW of Super-Ni laminated composite. Similar phenomenon occurred in the welding of iron-based powder alloys. Sintering additives in base layer hinder the fusion of base material. It is the main reason for the formation of such large scale (~400 μm in length) holes defects. Microcracks in the fusion zone of NiCr base layer had significant effect on the performance of welded joints, as shown in Fig. 6.6b. Microcracks initiated from the

(a)

(b) Weld metal

Ni80Cr20 100μm

partial fusion

Fig. 6.6 Partial fusion and microcracks in Ni80Cr20 base layer

microcracks

6.1 Characteristics and Weldability of Laminated Materials

305

poor-bonded large grain boundaries, propagated along the large grain boundaries and finally terminated at the sintering holes in NiCr base layer. The sintering holes inhibited further expansion of the microcrack, which was beneficial for the welded joint to maintain its serviceability. (4) Stress and liquation cracking Ni content of the weld seam near the laminated composite reached 40%, Creq /Nieq < 1.52. The welding solidification mode was AF mode with obvious directional columnar crystal growth and thermal crack sensitivity. The porosities in NiCr base layer has an important influence on weldability, leading to the welding of NiCr base layer differ from that of conventional rolled materials. Due to the presence of porosities in NiCr base layer, the bone-shaped structure in the heat affected zone (HAZ) of NiCr base layer partially melted under the action of the large welding arc heat input and large size holes may appear after re-solidification and shrinkage. As the porosities exist in NiCr base layer, there is great difference of thermal expansion coefficients between the laminated material and 18–8 steel, which affects the stress distribution in the weld joint and even triggers liquation cracks. The Ansys finite element analysis was applied to investigate the stress distribution of wire-filled GTAW joint between Super-Ni/NiCr laminated composite and 18–8 steel stress. It was found that the stress concentrated in the fusion zone of laminated composite, the stress in super-Ni cover layer was greater than that in NiCr base layer and the interface between Ni cover layer and NiCr base layer was the weak region.

6.1.3 Research Status of Laminated Materials Welding Advanced welding technologies have an important advantage in realizing new ideas in structural design, such as reducing structural mass, reducing manufacturing costs, improving structural performance, etc. Research of the welding of SuperNi/NiCr laminated composites will provide theoretical and experimental basis for their application. Due to the specificity of the chemical composition and structure of Super-Ni/NiCr laminated composites, its weldability study involves the welding of nickel-based high-temperature alloys, powder high-temperature alloys and laminated composites. The special “sandwich” compound structure of laminated material is one of the important factors affecting its weldability. As the laminated composite combines the excellent properties of two metals, it can meet the requirements of many special applications. So its welding behavior research and application attract more and more attentions. Opening welding groove, welding cover layer and base layer separately and inserting transition layer are usually used in the welding of medium and thick composite materials, such as the welding of composite steel. There are also researchers studying the single-pass welding of the composite plate. However, the

306

6 Welding of Laminated Materials

welding methods for medium and thick composite materials are not suitable for laminated composite due to the thickness of the cover layer is only 0.2~0.3 mm. So solve the problems of laminated composite welding is the key to promote its application. TiAl-based alloy plates with a thickness of 2.7 mm were prepared at 1050 °C by Huang Boyun et al. of Central South University using the envelope rolling technique. Metallographic analysis showed that the thin plates had uniform and fine isometric crystals with an average grain size of about 3 μm. The envelope rolling technique can reduce the rheological stress during deformation of TiAl-based alloys, retard the rheological softening tendency and reduce the local rheological coefficient, thus improving the plastic deformation ability of TiAl-based alloys. The special cover layer structure of the laminated material is the key factor in its weldability. As the laminated composite consists of two or more materials with great differences in chemical compositions and mechanical properties, both cover layer and base layer should be taken into account in the welding process. Shandong University achieved the connection of Super-Ni/NiCr laminated composite with 18–8 steel by wire-filled tungsten arc welding (GTAW) and diffusion brazing, and obtained a good combination in the fusion zone. Since the thickness of Super-Ni cover layer was only 0.3 mm, it was prone to be burnout in the welding process. It is necessary to open a groove on the laminate side and control the arc towards to the 18–8 steel side. The weldability of double-sided ultra-thin stainless steel composite (thickness of cover layer < 0.5 mm) was studied. The cover layer and base layer was 18–8 steel and Q235 steel, respectively. The stainless steel composite plate (0.25 mm + 3 mm + 0.25 mm) was welded by tungsten argon arc welding, fusion electrode argon arc welding and micro-beam plasma arc welding. Comprehensive analysis of the advantages and disadvantages of various welding processes and research on the electrochemical corrosion performance and mechanical properties of the joint would promote the application of ultra-thin stainless steel composite. Nd:YAG pulse laser was used to butt weld the double-sided ultra-thin stainless steel composite plate (0.1 mm stainless steel + 0.8 mm carbon steel + 0.1 mm stainless steel). In order to ensure the corrosion resistance of the weld seam was consistent with the stainless steel cover layer, Fe alloy powder with high Cr and Ni content was used as the filler metal. The weld gap was combined with stainless steel cover layer and carbon steel base layer well. Tensile strength of the joint reached 92% of base material, and the elongation was 25% of base material. Some researchers studied the weldability of double-sided thin composite material, the cover layer was 18–8 steel, the base layer was Q235A, the thickness size was 0.8 mm + 5 mm + 0.8 mm. Learning from the welding method of medium and thick composite plate, the base layer was welded by manual arc welding and the base layer was welded by gas tungsten arc welding (GTAW). The joint meeting the use requirements was obtained. However, because the cover layer was very thin, the requirements for groove processing and welding operations were high, and the welding efficiency was low.

6.1 Characteristics and Weldability of Laminated Materials

307

Other researchers have studied the resistance spot welding behavior of two metal laminated materials, which were rolled from three layers of 0.5 mm thick steel plates using pure Zn or 95% Pb-5% Sn as an interlayer. The study showed that the two laminated materials exhibited good weldability. The joint strength of the laminated material with Zn interlayer was higher than that with 95% Pb-5% Sn interlayer. Fe-Zn and Fe-Sn intermetallic compounds generated in the joints. NiAl/V and NiAl/Nb–15Al–40Ti micro-laminated composites were prepared by Ohio State University, USA. NiAl powder was alternately laminated with V foil or Nb-15Al-40Ti foil and put in a stainless steel sleeve. The sleeve was vacuumized and sealed by electron beam welding, and then hot isobaric pressed at 1100 °C × 270 MPa for 4 h. The fracture toughness of the micro-laminated composites was investigated by three-point bending test with pre-existing cracks as shown in Fig. 6.7. For NiAl/V micro-lamianted composites, the initial crack stopped when it extended to the ductile layer. As the load increased, the cracks further extended after forming slip bands along 45° on both sides of the ductile layer, as shown in Fig. 6.7a. The ductile layers and brittle layers separated. The fracture of NiAl bulk material was brittle intergranular fracture. The separation region displayed fracture features of the dimple. Cracks in NiAl/Nb–15Al–40Ti micro-laminated composite propagated along the grain boundaries. When the thickness of Nb-15Al-40Ti layers was 500 μm, crack bridge formed as shown in Fig. 6.7b; When the thickness was 1000 μm, no crack bridge formed and the fracture surface was mixed fracture morphology. The density of the NiCr base layer is about 80% of the density of the dense material. Super-Ni/NiCr laminated material was welded with 18–8 steel using electron beam welding, micro-beam plasma arc welding and laser welding. Super-Ni/NiCr laminated composite was penetrated intensively during electron beam welding and micro-beam plasma arc welding. The weld spatter was serious and it was difficult to control the fusion zone of the laminated composite to obtain good welding appearance.

Fig. 6.7 Crack extension paths in micro-laminated composites a NiAl/V micro-laminated composites b NiAl/Nb-15Al-40Ti micro-laminated composites

308

6 Welding of Laminated Materials

In the laser welding process, when the laser power was 500–600W, the cover layer was bonded well but penetration of Super-Ni/NiCr laminated composite was insufficient. When the laser power was 700–1000W, discontinuous micro-hole appeared in cover layer. When the laser power increased to 1500W, continuous micro-holes appeared with obvious weld spatter and the fusion became sharply worse. Due to the special multi-layer structure of Super-Ni/NiCr laminated composites and the NiCr base layer is a powder sintered alloy, the impact on the NiCr base layer is high when using high-energy beam welding (including electron beam welding, plasma arc welding and laser welding, etc.). It is difficult to obtain good weld formation. Tungsten argon arc welding has flexible process parameters adjustment, which is commonly used in the powder alloy welding. Super-Ni laminated material is a promising new high-temperature structural material that combines the excellent properties of powdered high-temperature alloys with those of traditional high-temperature alloys. Welding is an important forming method of manufacturing technology. Good welding of laminated composite will improve its utilization rate and optimize component performance. The welding of Super-Ni laminated composite is very different from traditional metal materials. Groove and multi-pass welding are usually used in the traditional composite steel welding. But that are not suitable for the cover layer with the thickness of only 0.3 mm. Welding characteristics of Super-Ni cover layer and NiCr base layer become the focus of research on the weldability of laminated composites. It is necessary to study the special weld formation and microstructure of laminated composites and establish the connection between microstructure and joint property, which has important significance for clarify the weldability of laminated composites and promote their industrial applications.

6.2 Wire-Filled GTAW of Laminated Materials In many components, only a part needs to withstand high temperatures, high stress or corrosive medium. Therefore, composite structure formed by welding laminated materials with other materials can not only take full advantages of different materials, but also save precious metal materials, which has important economic value.

6.2.1 Process Characteristics of Wire-Filled GTAW of Laminated Materials Super-Ni laminated composites were welded by wire-filled gas tungsten arc welding (GTAW) method. The welding process parameters were accurately controlled to form a soft arc, which can realize the weld formation in a single process. GTAW was completed by inverter argon arc welding machine (welding current adjustment

6.2 Wire-Filled GTAW of Laminated Materials

(1 Super-Ni cover layer;

309

2 Ni80Cr20 base layer; 3-18-8 steel)

Fig. 6.8 Schematic diagram of assembly

range: 15~150 A), pulse wire filling. Firstly, the welding process test was carried out. Before welding, the groove of laminated material was processed, as shown in Fig. 6.8, and the assembly gap was less than 0.5 mm. Before welding, the surface of the samples (Super-Ni/NiCr laminated material, 18–8 steel) was machined. The oil pollution, corrosion, oxide film and other pollutants on the surface of the base metal and filler material (0Cr25-Ni13 alloy wire) were removed by chemical method. The mechanical and chemical surface treatment steps of welding test plate were: sandpaper grinding → acetone cleaning → water washing → alcohol cleaning → drying. 0Cr25-Ni13 alloy wire was used as filler metal in the welding process. Wire-filled gas tungsten arc welding (GTAW) was used and the process parameters used in the experiment are shown in Table 6.4. Diameter of the wire was 2.5 mm, and diameter of the tungsten electrode was 2.0 mm. Since the thickness of the super nickel cover layer was only 0.3 mm, small welding heat input was required and the arc direction should be strictly controlled. The macroscopic weld morphology is shown in Fig. 6.9. It was found in the test that the tungsten arc should be slightly biased to the 18–8 steel side. When the tungsten arc directly pointed to the Super-Ni laminated composites, the Super-Ni cover layer melted too fast, which did not synchronized Table 6.4 Welding process parameters used in the test Welding current Welding voltage Welding speed Argon rate Heat input Remarks /V /cm·s−1 /L·min−1 /kJ·cm−1 /A 80

10~12

0.08

8

7.6~9.0

Arc bias to laminate

80

10~-12

0.12

8

5.0~6.0

Arc centered

80

11~12

0.20

8

3.3~3.6

Arc bias to 18–8 steel

Note The effective arc heating coefficient η was 0.75

310

6 Welding of Laminated Materials

Fig. 6.9 Macroscopic weld morphology

Welding direction 120 mm 1Cr18Ni9Ti 120 mm

Weld metal

Super-Ni/NiCr 2.6 mm

Cutting specimen

Fig. 6.10 Cutting schematic diagram of GTAW butt welding joint specimens

with the melting of NiCr base layer. It was difficult to ensure the welding formation stability of Super-Ni cover layer. Microstructure and properties of Super-Ni laminated material and 18–8 steel GTAW welded joints were tested and analyzed. Firstly, the samples was cut and prepared, and the surface treatment and microstructure etching of the welding zone were carried out. A series of specimens were cut from the joint of Super-Ni laminated material and 18–8 steel GTAW by electric spark machine. The sample cutting schematic diagram of GTAW butt welding joint is shown in Fig. 6.10.

6.2.2 Fusion State of Welding Zone of Laminated Materials (1) Welding metallurgy of laminated materials and division of joint zones 1) Welding metallurgy of laminated materials GTAW process of Super-Ni laminated material to 18–8 steel with filler wire mainly involves two aspects of welding metallurgical process: First is the interaction of Ni, Cr, Fe elements. Interaction characteristics of several major elements are shown in

6.2 Wire-Filled GTAW of Laminated Materials

311

Table 6.5 Interactions of Fe, Cr and Ni Alloy elements

Melting Crystal point/ °C transition temperature/ °C

Lattice type

Atomic radiu/nm

Infinite

Form solid solution Finite

Form compounds

Fe

1536

910

α-Fe bcc γ-Fe fcc

0.1241

α-Cr,γ-Ni

γ-Cr,α-Ni

Cr, Ni

Cr

1875

–

bcc

0.1249

α-Fe

γ-Fe,Ni

Fe, Ni

Ni

1453

–

fcc

0.1245

γ-Fe

Cr, α-Fe

Cr, Fe

Table 6.5. The weldability is good when the elements are prone to form infinite solid solution metal. Second is the special pressing structure of the laminated material, which makes its welding behavior quite different from that of the conventional metal. When the NiCr base alloy is welded, it is easy to form a jagged fusion zone. Microstructure of the fusion zone has an important influence on the microstructure and properties of the laminated material joints. When Super-Ni laminated material and 18–8 steel were welded by wire-filled GTAW (with 0Cr25-Ni13 filler wire), the content of Cr in Super-Ni laminated material, 18–8 steel and 0Cr25-Ni13 filler wire were similar, but the contents of Fe and Ni were very different. So the welding metallurgical characteristics of Super-Ni laminated material can be analyzed by 20% Cr-Fe–Ni phase diagram (see Fig. 6.11). According to the different transition degrees of elements, 20% Cr-Fe–Ni alloy has four solidification modes: Alloy ➀, complete the solidification process with δ phase, the solidification mode is F; Alloy ➁, with δ phase as primary phase, as exceeding the AC surface, peritectic and eutectic reaction L + δ → L + δ + γ → δ + γ occur successively, and the solidification mode is FA; Alloy ➂, primary phase is γ, then the following reaction occurs L + γ → L + δ + γ → δ + γ, the solidification mode is AF; Alloy ➃, complete the whole solidification process with γ phase, the solidification mode is A. In the welding process of austenitic steel, hot cracks are prone to form in the weld and near-seam area, and the most common is weld solidification cracks. The solidification mode is related to the ratio of ferritizing and austenitizing elements (Creq /Nieq ) in the weld. The content of ferrite element is converted into content of Cr according to its ferritize capacity. Creq represents the sum of converted Cr content. Similarly, Nieq represents the sum of converted Ni content.

312

6 Welding of Laminated Materials

A

Fig. 6.11 20% Cr-Fe–Ni phase diagram

F solidification model

AF F

1600 L

A C

B 1400

L+δ

L+γ

T / ºC

L+δ+γ δ 1200 γ δ+γ 1000 20%Cr 800

20

10 ω Ni /%

0

It is found that the Creq /Nieq value determining the solidification mode is the key factor affecting the hot crack. When Creq /Nieq > 1.52, the primary phase is primarily δ ferrite phase. During the solidification process, δ ferrite transition to γ austenite, and finally the weld microstructure consists of γ austenite + a small amount of δ ferrite, which is generally not easy to induce hot cracks. When Creq /Nieq < 1.52, the primary phase is γ-austenite phase, and a small amount of δ-ferrite precipitates during the cooling process. The toughness of weld decreases significantly, and the tendency of hot crack is obvious. When the 18–8 steel was welded with 0Cr25–Ni13 alloy wire, the Creq /Nieq of the weld was between 1.5 and 2.0, and it was easy to form an austenitic weld containing a small amount of δ ferrite. The weld had good comprehensive mechanical properties, and the hot crack tendency was small. When the laminated material was welded with 18–8 steel, the laminated material was mainly composed of Ni element, transiting into the weld. When Creq /Nieq < 1.52, the higher the Nieq is, the smaller the ratio is, and the hot cracking tendency is obvious. It can effectively reduce the hot crack sensitivity by controlling fusion ratio reasonably, especially the fusion ratio of Super-Ni laminated material, reducing the Ni content of weld. The austenitic steel filler material was used in the test. The filler alloy composition can be determined by means of Schaeffler weld microstructure diagram (see Fig. 6.12).

6.2 Wire-Filled GTAW of Laminated Materials

313

26 100 Super-Ni cover layer

a 80

c NiCr base layer

e

f δ

Nieq=Ni+30C+0.5Mn /%

24 60

0%

18 h

A

16

W

d Cr25Ni13

A+M

12 10

20%

b M

8

0

2

.5

5% eq

10 %

g

14

=1 C r eq

i 5N

A+F

1Cr18Ni9Ti 12

14

16

18

20

22

24

26

Creq=Cr+Mo+1.5Si+0.5Nb /% Fig. 6.12 Scheffler weld microstructure

The base metal of laminated material belongs to NiCr alloy with high Ni content. According to the prediction of weld microstructure of dissimilar metals, when SuperNi cover layer (point a in Fig. 6.12) is welded with 1Cr18Ni9Ti stainless steel (point b) using 0Cr25–Ni13 welding wire (point d), the weld microstructure locates on point g. When NiCr base layer (point c) is welded with 1Cr18Ni9Ti stainless steel (point b) using 0Cr25–Ni13 welding wire (point d), the weld microstructure locates on point h. In the ideal state, the composition of weld metal should be controlled in the area W as shown in Fig. 6.12 to ensure that the weld has good hot crack resistance. The mass fraction formula of an element in dissimilar metal joints is: ωW = (1 − θ )ωd + kθ ωb1 + (1 − k)θ ωb2 In the formula: ωW ωd ωb1 , ωb2 k θ

- Mass fraction of an element in weld metal; - Mass fraction of an element in deposited metal; - Mass fraction of an element in base metal 1, 2; - Relative fusion ratio of two base metals; - fusion ratio.

(6.2)

314

6 Welding of Laminated Materials

The Ni content in the weld is related to the fusion ratio and relative fusion ratio of the base metal, so it is necessary to strictly control the fusion ratio (γ) of the base metal to ensure that the weld metal microstructure falls in the zone W as shown in Fig. 6.12. The control of fusion ratio is related to the composition of base metal and welding process parameters (heat input). In order to ensure the Creq /Nieq > 1.52 in weld metal, the fusion ratio of laminated material and 18–8 steel should be less than 10%. Fusion ratio can be controlled by opening groove and reducing welding heat input. 2) Division of welding joint zone of laminated materials In order to facilitate the analysis of the microstructure characteristics of SuperNi/NiCr laminated materials and 18–8 steel GTAW joints in different regions, the GTAW welded joint of laminated composites can be divided into three characteristic regions, as shown in Fig. 6.13. (i)

The transition zone between Ni cover layer and weld, including fusion zone and heat affected zone on the side of Ni cover layer; (ii) The transition zone between Ni80Cr20base and weld, including fusion zone heat affected zone on the side of NiCr base layer; (iii) The weld center region, including columnar crystal region and equiaxed crystal region. The GTAW formation of Super-Ni laminated material and 18–8 steel is complex. The transition zones near the Ni cover layer and Ni80Cr20 base layer have the greatest influence on the microstructure and properties of laminated material, which is the focus of the weldability analysis of laminated material. Fusion zone HAZ of Ni cover layer

Weld

Ni cover layer

NiCr base layer

Columnar zone Fusion zone

HAZ of NiCr base layer Fig. 6.13 Characteristic zone division of Super-Ni laminated material joints

6.2 Wire-Filled GTAW of Laminated Materials

315

A weld with a certain penetration and uniform transition can be obtained by wire filler tungsten argon arc welding (GTAW) of Super-Ni/NiCr laminated composite and 18–8 steel. Complete welded joint area includes four typical areas: (i) (ii) (iii) (iv)

Transition zone between Ni cover layer and weld; Transition zone between Ni80Cr20 base layer and weld; Central weld zone; Transition zone near 18–8 steel.

The microstructure of Super-Ni laminated composite GTAW welded joint with 18– 8 steel is shown in Fig. 6.14a. Super-Ni cover layer and the weld seam combined well and the weld surface was smooth. A good transition zone was formed between Ni80Cr20 base layer and the weld seam. The fusion zone of Ni80Cr20 alloy base layer was different from that of traditional casting or rolling alloys. The serrated fusion zone was formed due to the existence of sintering and pressing pores, which was very similar to the welding of iron-based powder alloy. The transition zone at the laminated composite side is shown in Fig. 6.14. Super-Ni cover layer combined well with the weld metal. Good weld appearance of Super-Ni cover layer was conducive to maintaining the unique heat resistance and corrosion resistance of laminated materials. Due to the effect of welding arc temperature gradient, grain size of the weld metal near the transition zone was fine. The transition zone between Ni80Cr20 base layer and the weld is shown in Fig. 6.14b. The bond between weld and NiCr base layer was weak and the interface was sawtooth-like indicating partial fusion. Different from conventional nickel-based superalloys, the microstructure of fusion zone of NiCr base layer was also different due to its special skeletal structure. The sawtooth-like fusion zone had great influence on the strength, high temperature resistance and corrosion resistance of the weld. When the GTAW was executed in small welding heat input (such as small current soft arc) with suitable filler alloy wire, the burning loss of Super-Ni cover layer was greatly reduced. A better fusion of NiCr base layer was obtained under the action of soft arc blowing, which improved the overall performance of the weld joint.

(b)

(a)

Weld Ni cover layer Weld Ni cover layer

Ni80Cr20

100μm

Fusion state of Ni coating

Ni80Cr20

Fusion state of base metal

Fig. 6.14 Super-Ni laminated material fusion zone and weld microstructure

100μm

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6 Welding of Laminated Materials

(a)

(b)

100μ equiaxed zone

50μm weld central

Fig. 6.15 Microstructure of weld metal center

The growth of austenite columnar crystals and the existence of low melting point segregation impurities in the weld weakened microstructure between large grain boundaries, which increased the thermal crack sensitivity. The weld center consisted of equiaxed austenite grains with uniform size, as shown in Fig. 6.15. (2) Microstructure characteristics of the weld near laminated materials side Due to the special structure of Super-Ni laminated material, two typical transition zones were formed after filler wire GTAW welding: the transition zone between Ni cover layer and weld, and the transition zone between Ni80Cr20 base layer and weld. 1) Transition zone between Super-Ni cover layer and weld metal The thickness of Super-Ni cover layer was only 0.3 mm, and the microstructure of the weld metal is shown in Fig. 6.16. The bonding between Ni cover layer and weld seam was good, and the fusion transition zone was clear. Obvious fusion zone and heat affected zone (Fig. 6.16 a) formed in the transition zone between Ni cover layer and weld. The grains in the HAZ of Ni cover layer recrystallized due to the welding thermal cycle and evolved from the original rolling elongated structure to massive structure. Near the fusion zone, grains in heat affected zone tended to coarsen. The columnar crystal structure in the weld was perpendicular to the interface and the grain was fine. The microstructure sensitization zone was formed near the fusion zone, and the grain boundary morphology throughout the fusion zone was formed. It shows that the good metallurgical bonding between the base metal and the weld was beneficial to improve the bonding strength near the fusion zone, so as to ensure the strength of the whole weld. 2) Transition zone between Ni80Cr20 base layer and weld metal

6.2 Wire-Filled GTAW of Laminated Materials

(a)

(b)

Ni cover layer

Ni80Cr20

317

Weld

100μm

OM

SEM

Fig. 6.16 Microstructure of fusion zone of Super-Ni cover layer

Ni80Cr20 base layer is a powder sintered alloy, and its fusion zone is different from conventional metals. Microstructure and appearance of the fusion zone are important factors affecting the performance of Super-Ni laminated material weld joints. The fusion zone between the Ni80Cr20 base layer and weld formed well. Because of the existence of internal pores in the sintered powder alloy, the fusion zone microstructure was completely different from that of the traditional casting or rolling metal. The NiCr metal particles in the powder alloy base layer melted at high temperature and formed metallurgical bonding with the filler metal, and the fusion zone was sawtooth-like and discontinuous. The microstructure of the fusion zone was columnar and the grain size was smaller than that of the NiCr base layer. The grain growth direction was perpendicular to the boundary line between the fusion zone and the heat affected zone. Large grain boundaries were formed between different columnar crystal groups, and finally solidified during the cooling process. The transition zone of NiCr alloy base was more obvious than that of super-Ni cover layer. The fusion zone of super-Ni cover layer was narrow. The thermal conductivity of super-Ni cover layer was much higher than that of NiCr base layer. In the weld cooling process, the temperature gradient near the fusion zone of NiCr base layer was larger, with obvious columnar crystal growth morphology. In the fusion zone near the super-Ni cover layer, the temperature gradient of the weld metal during cooling was small. The grains were not obvious columnar crystal morphology but equiaxed crystal morphology. However, due to the convective cooling effect of air on the weld surface, the temperature gradient increased and the columnar crystal morphology was noticeable. There were large-scale holes in the fusion zone between NiCr base layer and weld metal. This was for the reason that the wettability of liquid filled metal on sintered powder alloy matrix was worse and the metallurgic combination was difficult. The existence of such holes had an adverse effect on the bonding strength of the fusion zone, which can be controlled by adjust process parameters (heat input).

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6 Welding of Laminated Materials

(b)

(a)

100μm

100μm

columnar crystal zone

equiaxed crystal zone

Fig. 6.17 Columnar and equiaxed crystal regions in the weld seam

During the welding process, the original bone-shaped structure in HAZ of NiCr base layer changed and large-size ‘holes’ appeared. The ‘holes’ were different from those in the fusion zone. Under the heating of welding arc, low temperature phase among the bone-shaped matrix in NiCr lase layer local melted and recrystallized, resulting in new connection and local shrink with holes. This phenomenon existing in the welding of sintered powder alloys, also appear in the welding of iron-based powder alloys. 3) Microstructure characteristics of weld center Microstructure of Super-Ni laminated material and 18–8 steel wire-filled GTAW weld is shown in Fig. 6.17. The weld microstructure near 18–8 steel side was directional austenitic columnar crystal, growing perpendicular to the fusion zone (see Fig. 6.17a). The columnar crystal morphology of 18–8 steel side fusion zone was not as flat as that of laminated composite material side, and the microstructure scale was smaller. Columnar crystals on both sides of the weld grew towards the weld center, and gradually transformed into equiaxed austenite at the weld center (see Fig. 6.17b). A small amount of δ ferrite distributed on the austenite matrix. When there is 4~8% δ ferrite in austenitic stainless steel weld, it is beneficial to ensure the toughness of weld metal and prevent thermal cracks. The cooling rate was slower at the weld center, the austenite growth mainly parallel to the weld seam. While austenite columnar crystals interlaced near the weld surface. Due to partial Ni cover layer or NiCr base layer melted into the weld pool and Ni is austenitizing element, austenite content in the weld metal increased and the δ ferrite content decreased. Due to the transition effect of Ni element in the base metal on the side near the Super-Ni laminated material, Creq /Nieq < 1.52 in the local area of the weld was formed. The transformation from austenite to ferrite occurred. However, the weld at the 18–8 steel side, Creq /Nieq > 1.52, ferrite structure was first formed during the weld cooling process, and the transformation from ferrite to austenite occurred (FA solidification mode).

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319

The obvious transition from columnar crystal to equiaxed crystal was formed in the weld. Due to the different welding thermal cycles in different parts of the weld, the weld microstructure near the base metal on both sides showed columnar crystal morphology. The center area of the weld was uniformly heated with the formation of equiaxed crystal morphology.

6.2.3 Microstructure and Properties of Joint Between Laminated Material and 18–8 Steel (1) Effect of heat input on microstructure of laminated material joint Welding heat input affects the microstructure and weld formation of SuperNi laminated material GTAW joints, and different welding heat input can be realized by changing the welding speed. The welding heat input determined in the test was 3.3~3.6 kJ/cm, 5.0~6.0 kJ/cm, and 7.6~9.0 kJ/cm, respectively. By comparing the microstructure characteristics at the same position of the GTAW joint under different welding heat inputs, the regularity between welding heat input and microstructure of the joint was revealed. 1) Weld microstructure under different welding heat inputs The weld microstructure characteristics of wire-filled GTAW under different welding heat inputs are shown in Fig. 6.18. Microstructure transition from columnar crystal to equiaxed crystal formed under different welding heat input conditions. With the increase of welding heat input (from 3.3~3.6 kJ/cm to 5.0~6.0 kJ/cm and 7.5~9.0 kJ/cm), the welding speed became slower, the equiaxed grains in the weld center were coarsened. The faster the welding speed was, the more uneven the weld microstructure was. It is also found in the test that the weld microstructure with large welding heat input had partial remelting characteristics.

(b)

(a)

200μm

E=3.3~3.6 kJ/cm

100μm

E=7.5~9.0 kJ/cm

Fig. 6.18 Weld microstructure under different welding heat inputs

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6 Welding of Laminated Materials

In summary, as the welding heat input increased, the equiaxed grains in the weld center changed from fine to coarse. The faster the welding speed was, the finer the weld microstructure was, but the more uneven it was. With the decrease of welding speed, the welding heat input increased, length size of columnar crystal near base metals decreased. The weld formation was poor when welding heat input was too large. The microstructure of welds with heat input of 3.3~3.6 kJ/cm and 7.5~9.0 kJ/cm is shown in Fig. 6.18(a, b). 2) Microstructure on the side of laminated material under different welding heat inputs The microstructure on the side of Super-Ni laminated material also changed under different welding heat inputs. When the heat input was 3.3~3.6 kJ/cm, the laminated material and the filler alloy wire combined well with obvious columnar crystal morphology in the weld metal. When the heat input was 5.0~6.0 kJ/cm, the laminated material and the filler alloy wire also combined well with obvious columnar crystal morphology in the weld metal. However, the growth of columnar crystals was hindered by other columnar crystals, the length of columnar grains were shorter than the length when the heat input was 3.3~3.6 kJ/cm. As the welding speed decreased, the cooling rate decreased. So the length and size of columnar grains become smaller. When the heat input was 7.5~9.0 kJ/cm, the weld area near the surface and Ni cover layer displayed obvious columnar crystal morphology. The bonding between laminated composite and filler wire was worse and the joint microstructure was coarse especially the HAZ of laminated composite. (2) Effect of heat input on microhardness of laminated material GTAW joint In order to determine the change of microstructure and properties of GTAW joint between Super-Ni laminated material and 18–8 steel. Microhardness near the fusion zone of Super-Ni laminated material under different welding heat inputs was measured by Shimadzu microhardness tester of Japan, with a load of 50 gf and a loading time of 10 s. When the welding heat input was 3.3~3.6 kJ/cm, the microhardness of SuperNi laminated material side fusion zone was shown in Fig. 6.19. The microhardness near the fusion zone was higher than that of Super-Ni cover layer and weld metal, with a peak 190 HM. In the welding process, the cooling rate of the fusion zone was fast, firstly solidification and crystallization with a hardening tendency. The chemical composition of the weld metal near the fusion zone was not uniform. With the growth of columnar crystal structure, the composition was gradually uniform. The microhardness showed little change (mean 165 HM), indicating that there was no obvious brittle and hard phase in the weld. There were a large number of Cr elements in the Ni80Cr20 base layer and weld metal. While the Ni cover layer side has a dilution effect on the original Cr element in the weld metal due to the melting of Ni element into the weld metal. The hardness of high Cr phase was higher than that of low Cr phase. The microhardness near the fusion zone of Ni80Cr20 base layer was higher than that near the fusion zone of Super-Ni cover layer. The microhardness of NiCr base layer HAZ (mean 135

6.2 Wire-Filled GTAW of Laminated Materials

321

Microhardness /HM

(a)

200 (b) Base material 180

Weld 2

160

1

140

1. Fusion zone of Ni-Cr base layer 2. Fusion zone of Ni cover layer

120 100 -0.2

Measured position

-0.1

0.0 0.1 0.2 Distance/mm

0.3

Microhardness distribution

Fig. 6.19 Microhardness near fusion zone of laminated materials side (E = 3.3~3.6 kJ/cm)

HM) was higher than that of Ni cover layer HAZ (mean 108 HM). The temperature gradient at different positions varied greatly during the cooling process of welded joints, which was also an important reason for the different microhardness of weld microstructure. When the welding heat input was 5.0~6.0 kJ/cm, the microhardness test results near the fusion zone of Super-Ni laminated material are shown in Fig. 6.20. The average microhardness of weld (199 HM) was obviously higher than that of SuperNi laminated composite base metal. The average microhardness of NiCr base HAZ was 135 HM, and the average microhardness of Super-Ni coating HAZ was163 HM. The NiCr base was a powder alloy matrix, and the elastic effect was obvious in the microhardness testing, which made the microhardness fluctuation range larger. When the welding heat input was 7.5~9.0 kJ/cm, the microhardness test results near the fusion zone of Super-Ni laminated material are shown in Fig. 6.21. The average microhardness of weld was 166HM, while that of NiCr base layer HAZ was 240

(a) Microhardness /HM

220

(b) Base material

200

Weld

180 1

160

1. Fusion zone of Ni-Cr base layer 2. Fusion zone of Ni cover layer

140 120 100

100μm

2

-0.2

-0.1

0.0

0.1

0.2

Distance /mm

Measured position

Microhardness distribution

Fig. 6.20 Microhardness near fusion zone of laminated materials side (E = 5.0~6.0 kJ/cm)

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6 Welding of Laminated Materials

(a) Microhardness /HM

180 (b)

Weld

160

1

2

140 120 100

Base material -0.2

Measured position

1. Fusion zone of Ni-Cr base layer 2. Fusion zone of Ni cover layer

-0.1 0.0 0.1 Distance /mm

0.2

Microhardness distribution

Fig. 6.21 Microhardness near fusion zone of laminated materials side (E = 7.5~9.0 kJ/cm)

134HM, and that of Super-Ni cover layer HAZ was 128HM. The microhardness variation tendency was almost consistent with that when the heat input was 3.3~3.6 kJ/cm and 5.0~6.0 kJ/cm. The weld formation under large heat input (7.5~9.0 kJ/cm) was poorer than that under the first two smaller heat input. The change of welding heat input directly affected the transition of alloy elements and the solidification and crystallization of weld. The microhardness analysis near the fusion zone of three different welding heat inputs (3.3~3.6 kJ/cm, 5.0~6.0 kJ/cm, 7.5~9.0 kJ/cm) (see Table 6.6) shows that with the increase of welding heat input, the microhardness near the fusion zone of SuperNi laminated material side increased first and then decreased. The microhardness of the weld also increased first and then decreased. The microhardness of the Super-Ni cover layer HAZ also increased first and then decreased. Since the super nickel cover layer was only 0.3 mm, it was greatly affected by the thermal effect of welding arc, and the microhardness also showed obvious changes. In contrast, the microhardness of NiCr base layer HAZ did not change obviously. The microhardness of weld at 18–8 steel side decreased gradually, and the microhardness of HAZ increased first and then decreased. In summary, from the view of microhardness testing results, microhardness near the fusion zone increased slightly while microhardness of the other area tended to be consistent, indicating good structure uniformity. The microhardness near the fusion zone was higher, while the microhardness of the weld and NiCr base layer was lower. Table 6.6 Relationship between microhardness (HM) and welding heat input (E) near fusion zone of laminated materials side

Heat input /kJ·cm−1

Mean microhardness /HM Ni80Cr20 base layer

Super-Ni cover layer

Weld metal

3.3~3.6

136

108

165

5.0~6.0

135

163

199

7.5~9.0

134

128

166

6.2 Wire-Filled GTAW of Laminated Materials

323

6.2.4 Process Characteristics of Diffusion Brazing of Laminated Materials Super-Ni/NiCr laminated composite was prepared by putting Ni80Cr20 powder between the pure-Ni (Ni > 99.5%) cover layers and then hot pressing in vacuum. Super-Ni/NiCr laminated composite was in the thickness of 2.6 mm, while the cover layers were 0.3 mm and the NiCr base layer was 2.0 mm. NiCr base layer consisted of bone-shaped γ-Ni(Cr) solid solution. Porosity fraction of NiCr base layer was 35.4% and the density was 6.72 g/cm3 . The density of pure Ni was 8.90 g/cm3 . By comparison, the structural weight can be reduced by 24.5%. (1) Diffusion brazing equipment Heating temperature of the diffusion brazing is low, which has little influence on the Super-Ni/NiCr laminated composite and 18–8 stainless steel. This method can avoid problems resulting from the melting of base metal in the fusion welding, such as burn through of cover layer and shrinkage cavity of base layer. The whole assembly is heated in the diffusion brazing process and generates small deformation, leading to lower thermal stress and guarantees the dimensional accuracy. No flux is used in the vacuum diffusion brazing and no pollution is brought in the brazed joint. The vacuum diffusion brazing process of Super-Ni/NiCr laminated composite and 18–8 steel was conducted in Workhorse-II vacuum diffusion bonding equipment produced by the America Centorr Vacuum Industries. The appearance and structure of the equipment is shown in Fig. 6.22. The equipment mainly consists of vacuum furnace, automatic vacuum pumping system, hydraulic system, heating system, water circulation system and control system. Vacuum environment promotes the decomposition of oxide film and the oxidation of the brazing filler metal, which can ensure the quality of diffusion brazed joint. (2) Brazing filler metal Super-Ni/NiCr laminated composite has been widely concerned in the areas of aerospace, missile, aircraft design due to its special high temperature performance and corrosion resistance. In order to take full use of the performance advantages of laminated composite, it is better to choose filler metals with high temperature oxidation resistance and corrosion resistance, such as nickel based filler metal and cobalt based filler metal. While the main component of the filler metal is the same as the base metal, and the wettability of the filler metal on the surface of the base metal is better. During the cooling process of the brazed region, the primary phase with the same composition of the base metal tend to grow attaching to the grains of base metal and form good combination with the base metal. That is beneficial to promote the bonding strength. The brazing joint types used in this study were butt joint and lap joint. Element Cr added in the Ni-based filler metal can improve the oxidation resistance and bonding strength of the joint. Elements Si, B and P are added in the filler metal, which can bring down the melting point and improve the fluidity and wettability.

324

6 Welding of Laminated Materials

Fig. 6.22 Workhorse-II vacuum diffusion bonding equipment

However, the melting-point depressant elements (B, Si) would form brittle phases such as boride and silicide in (or nearby) the brazed region during the diffusion brazing process. That plays a great influence on the brazing quality. So the Ni– Cr-P and Ni–Cr-Si-B Ni-based filler metals with different melting-point depressant elements were used. Chemical composition and melting point of Ni–Cr-P filler metal and Ni–Cr-Si-B filler metal are shown in Table 6.7. Chemical composition of Ni–Cr-P filler metal is eutectic, which has a lower melting point and good fluidity and wettability. There is no B in Ni–Cr-P filler metal. So it is suitable for the brazing of thin-walled components due to its weak corrosion of base metal. Ni–Cr-P filler metal will not absorb neutrons and can be applied in the nuclear area. Ni–Cr-Si-B filler metal exhibits good high temperature performance and excellent bonding strength. It is appropriate for the components subjected to high stress in high temperature, such as turbine blades and jet engine. Furthermore, amorphous filler metal was used in this study. The composition of amorphous filler metal is uniform and its microstructure and thickness can be controlled. The amorphous filler metal foil can be preformed into different shapes according to the workpiece structure. The brazing assembly process will be simplified Table 6.7 Chemical composition and melting point of filler metals Filler metal

Chemical composition /%

Ni–Cr–P

Bal 13.0~15.0 9.7~10.5

Ni

Cr

Ni–Cr–Si–B Bal 6.0~8.0

P ≤0.02

Si

B

Fe

C

≤0.1

≤0.02

≤0.2

≤0.06

≤0.05 890

≤0.06

≤0.05 970~1000

4.0~5.0 2.75~3.5 2.5~3.5

Ti

Melting point /°C

6.2 Wire-Filled GTAW of Laminated Materials

325

and the brazing joint has high structure accuracy. But it is difficult to guarantee the precise clearance for the difficult-to-machine materials and components with complicated brazing mating surface. For these situations, the conventional filler metals in the form of paste or powder exhibit good flexibility. While the brazing clearance is larger than 100 μm, it is easily to form brittle phases of intermetallic compounds. Thus the holding time and the brazing temperature should be optimized in order to prevent the formation of intermetallic compounds. Conventional filler metal and amorphous filler metal of Ni–Cr-Si-B were used to braze Super-Ni/NiCr laminated composite to 18–8 steel in vacuum. Brazed clearance for the conventional filler metal was 100~150 μm. The diffusion- solidification process of the joint was studied, which provided fundamental basis for the control of brittle phase formation. (3) Technological parameters Super-Ni/NiCr laminated composite and 18–8 steel were cut into the size of 30 mm × 10 mm × 2.6 mm via wire-electrode cutting machining and cleaned in acetone. The faying surfaces were prepared by grinding with metallographical sand paper, and cleaned in alcohol. The specimens were assembled on the Mo plate. The specimens and Mo plate were separated by a piece of graphite paper between them, avoiding their combination through the spreading filler metals. Schematic diagram of butt joint assembly between Super-Ni/NiCr laminated composite and 18–8 steel was shown in Fig. 6.23. Ni–Cr-P solder paste was spread on the surface of the joint gap. In order to control the size of the joint gap, Mo wire with the diameter of 150 μm was put between the brazing surfaces. The assembled specimens were fixed by stainless steel block and put into the vacuum chamber. Technological parameters of the diffusion brazing cycle are illustrated in Fig. 6.24. Vacuum degree was 1.33 × 10–4 ~1.33 × 10−5 Pa. Brazing temperature was 940~1060 °C and holding time was 15~25 min for Ni–Cr-P filler metal. Brazing temperature was 1040~1120 °C and holding time was 20~30 min for Ni–Cr-Si-B filler metal. As the space in vacuum chamber was larger, heating process was proceeded step by step with several insulation stages, in oder to equalize temperature of vacuum chamber and

1

Super-Ni cover layer

2 NiCr base layer

Fig. 6.23 Schematic diagram of butt joint assembly

3

filler metal

4

18-8steel

326

6 Welding of Laminated Materials

Fig. 6.24 Technological parameters of the diffusion brazing cycle

specimens. The cooling process was performed in the vacuum chamber by circulating water until the temperature reached 100 °C. The initial cooling rate was about 10 °C/min. Then the brazed joints were cooled in the chamber spontaneously. Effects of brazing parameters on spreading of filler metals and bonding of laminated composite/18–8 steel joints were shown in Table 6.8. When the joint was brazed with Ni–Cr-P filler metal at 940 °C, although the brazing temperature was 50 °C higher than the melting point, good combination was not established. Molten Ni–Cr-P filler metal was inclined to spread to the side of Super-Ni cover layer, indicating that Ni–Cr-P filler metal had good fluidity and wettability on Super-Ni cover layer. Ni–Cr-P filler metal exhibited better fluidity than Ni–Cr-Si-B filler metal at same brazing temperature for the reason taht the melting point of Ni–Cr-P filler metal was lower. (4) Preparation of brazed joint samples In order to analysis microstructure and performance of laminated composite/18–8 steel brazed joints, transverse section of the brazed joints were ground and polished. Laminated composite has lower hardness compared with 18–8 steel, so the force should be applied evenly to prevent partial grinding. Thickness of Super-Ni cover layer was only 0.3 mm, but that was the key area to be analyzed. So it should ensure that cover layer and base layer were on the same horizontal plane. Finally, a series of samples were etched by a solution of 80 ml HCl, 13 ml HF and 7 ml HNO3 . The etching time was 1~2 min for optical microscopy observation, while a longer time 2~3 min was needed for scanning electron microscope (SEM).

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Table 6.8 Effects of brazing parameters on spreading of filler metals and bonding of laminated composite/18–8 steel joints Filler metal types Ni–Cr–P

Conventional Ni–Cr–Si–B

Amorphous Ni–Cr–Si–B

Brazing temperature/°C

Holding time /min

Bonding and spreading appearance

940

20

No combination, filler metal aggregated as particles

980

20

Good combination, filler metal spread to the surface of laminated composite

1040

20

Good combination and flat brazed seam, filler metal spread to the surface of laminated composite

1060

20

Good combination and flat brazed seam, filler metal spread to the surface of laminated composite

1040

20

Ordinary combination, a certain thickness on the surface of brazing joint

1060

20

Good combination, a certain thickness on the surface of brazing joint

1080

20

Good combination, a certain thickness on the surface of brazing joint

1100

20

Good combination and flat brazed seam, a certain thickness on the surface of brazing joint

1120

20

Good combination, filler metal spread sufficiently

1060

20

Good combination and flat brazed seam, a certain thickness on the surface of brazing joint

1080

20

Good combination and flat brazed seam, a certain thickness on the surface of brazing joint

1100

20

Good combination and flat brazed seam, small thickness on the surface of brazing joint (continued)

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6 Welding of Laminated Materials

Table 6.8 (continued) Filler metal types

Brazing temperature/°C

Holding time /min

Bonding and spreading appearance

1120

20

Good combination, filler metal spread sufficiently

6.2.5 Bonding Interface of Laminated Composite/18–8 Steel Diffusion Brazed Joint (1) Division of characteristic zones Ni–Cr-P and Ni–Cr-Si-B filler metals were used to bond Super-Ni/NiCr laminated composite to Cr18–Ni8 steel by diffusion brazing. Interactions between filler metal and base metal mainly involve two aspects: (1) element diffusion from filler metal to base metal; (2) dissolution of base metal to molten filler metal. According to the diffusion-solidification characteristics of Super-Ni/NiCr laminated composite and 18–8 steel diffusion brazed joint, the joint were divided into five characteristic zones as shown in Fig. 6.25. ➀ ➁ ➂ ➃ ➄

Diffusion affected zone (DAZ) of Super-Ni cover layer; DAZ of NiCr base layer; Isothermal solidification zone (ISZ); Athermal solidification zone (ASZ); DAZ of 18–8 steel.

P, Si and B were added in Ni–Cr-P and Ni–Cr-Si-B nickel based filler metal as melting-point depressant elements in order to improve the fluidity and wettability. During the temperature holding stage, P, Si and B diffused from the liquid interlayer into Super-Ni/NiCr laminate composite and 18–8 steel. At the same time, some base metal dissolved into the molten filler metal. So the melting point of the liquid near the base metal increased. When the melting point reached the brazing temperature, isothermal solidification occurred with the formation of γ-Ni solid solution. As the isothermal solidification continued, the interface between the liquid and solid moved towards the center line of the brazed region. Excess solute elements were enriched in the residual liquid near the solid/liquid interface. During the cooling process, athermal solidification occurred in the residual liquid with the formation of brittle phases, such as phosphide, boride and silicide. Moreover, P, Si and B diffused into Super-Ni/NiCr laminate composite, generating new phases with Ni and Cr. That influenced microstructure and properties of the diffusion brazed joint. Porosities in NiCr base layer generate significant capillary force pulling molten filler metal into the porosities and there are three different cases to consider. When the capillary force is too great, it will pull excessive filler metal into the porosities leaving an insufficient amount of filler metal in the joint gap, which is unfavorable to create a strong bonding. Besides, filler meal in the porosities will react with

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329

(a)

DA ISZ

ASZ

ISZ DAZ of Super-Ni

(b) Super-Ni

NiCr base layer Cr18-Ni8

DAZ of NiCr base layer (a) Division of characteristic zones (b) microstructure of brazed joint Fig. 6.25 Division of characteristic zones for diffusion brazed joint of Super-Ni/NiCr laminated composite to 18–8 steel

base metal inducing internal stress and microcracks. Infiltration to short distance is advantageous in creating a strong bond by increasing the contact area. A sufficient amount of filler metal will be remained in the joint gap and a good bonding without voids and flaws will be obtained between Super-Ni/NiCr laminated composite and 18–8 steel. When there is no filler metal infiltration into porosities, the contact area is diminished which will be the weak area in the brazed joint. (2) Microstructure of brazed region Diffusion brazing of Super-Ni/NiCr laminated composite to 18–8 steel was performed with Ni–Cr-P filler metal at 1040 °C for 20 min. Microstructure of the

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6 Welding of Laminated Materials

Fig. 6.26 Microstructure of laminated composite/18–8 joint (Ni–Cr-P filler metal, T = 1040 °C)

Brazed region Super-Ni

Cr18-Ni8

NiCr base layer

joint is shown in Fig. 6.26. Ni–Cr-P filler metal exhibited good wettability to both the cover layer and base layer. There was no porosity, crack, non-fusion or other defects through the whole joint. There was a sufficient amount of filler metal remaining in the joint gap with minimal infiltration into the porosities of NiCr base layer, which was in favor of joint strength. It was for the reason that γ-Ni solid solution formed during the isothermal solidification, which prevented the filler metal infiltrating into NiCr base layer. This could avoid the formation of porosities and keep stability of porous structures. In order to analysis the microstructure characteristic in detail, the brazed region was amplified (Fig. 6.27) and chemical composition of the special phases was analyzed by EDS. Net-shaped eutectic formed in the center of the brazed region. The content of Ni and P for point 1 was 69.66% and 22.03%, respectively. The Ni/P ratio was nearly 3. Point 2 was rich in Ni, the content of P was about 8.42%. According to Ni–P phase diagram, when the content of P exceeds 0.32 at.%, Ni–P binary eutectic will separate out. Therefore, The phases (labeled by 1 and 2) was just the binary eutectic of Ni3 P and γ-Ni(P) solid solution. The phase in the brazed region near to base metals was γ-Ni(Cr) solid solution. Due to the Fe in 18–8 steel diffused into the brazed region, the Fe content inγ-Ni solid solution near 18–8 steel (5.58%, point 4) was higher than that near laminated composite (1.75%, point 3). (3) Effect of brazing temperature on microstructure of brazed region When the joint was brazed at 940 °C, Ni–Cr-P filler metal aggregated as particles at the sides of joint gap. Although the filler metal has molten, the fluidity was not good enough for the molten filler metal to spread sufficiently at the low brazing temperature. Therefore, brazed joint of Super-Ni/NiCr laminated composite to 18–8 steel cannot be established when brazed at 940 °C for 20 min.

6.2 Wire-Filled GTAW of Laminated Materials

DAZ

SSZ

ASZ

SSZ

331

DAZ

1+

3+

2+

4+

(a) brazed region; (b) eutectic zone; (c) interface between the brazed region and NiCr base layer; (d) interface between the brazed region and 18-8 steel Fig. 6.27 Microstructure of brazed region for laminated composite/18–8 joint (Ni–Cr-P filler metal, T = 1040 °C)

Microstructure of joint brazed at 980 °C for 20 min is shown in Fig. 6.28. The brazed region was mainly composed of γ-Ni solid solution and Ni–P eutectic. There were unmelted filler metal islands embedded in the brazed region. The composition of filler metal was not homogeneous and there were not only one phase in the filler metal. So the low melting point composition divorced with the high melting point composition in the heating process, which was the so-called composition segregation. When the temperature reached the liquid line, the low melting point phases melted and flowed firstly and the high melting point phases aggregated as islands due to its weak fluidity. Microstructure of joint brazed at 1060 °C for 20 min is shown in Fig. 6.29. The filler metal exhibited good wettability and fluidity on Super-Ni/NiCr laminated composite and 18–8 steel. There was no porosity, crack, non-fusion and other defect through the brazed region. As the brazing temperature increased, the Ni–P eutectic scope decreased and the solid solution scope near the base materials increased.

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6 Welding of Laminated Materials

(a)

(b)

Super-Ni

filler metal island

18-8

NiCr base layer

Brazed region

Filler metal island

Fig. 6.28 Microstructure of laminated composite/18–8 brazed joint (Ni–Cr-P filler metal, T = 980 °C) (Ni–Cr-P filler metal, T = 1060 °C)

Fig. 6.29 Microstructure of laminated composite/18–8 joint

(4) Diffusion-solidification process of the brazed region Because of there was concentration gradient between filler metal and base metal, element interdiffusion occurred. Elemental distribution through brazed region of Super-Ni cover layer and brazed region of NiCr base layer were shown in Fig. 6.30 and Fig. 6.31. 1) Ni–Cr-P filler metal Diffusion brazing process of Super-Ni/NiCr laminated composite and 18–8 steel was divided into four stages as shown in Fig. 6.32.

6.2 Wire-Filled GTAW of Laminated Materials

ISZ

ASZ

333

ISZ Super-Ni

Cr18-Ni8

(a) Measured position ISZ

ASZ

ISZ

(b) Elemental distribution Fig. 6.30 Elemental distribution through brazed region of Super-Ni cover layer (Ni–Cr-P filler metal, T = 1040 °C)

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6 Welding of Laminated Materials

Cr18-Ni8

NiCr

(a) Measured position ISZ

ASZ

ISZ

(b) Elemental distribution Fig. 6.31 Elemental distribution through brazed region of NiCr base layer (Ni–Cr-P filler metal, T = 1040 °C)

6.2 Wire-Filled GTAW of Laminated Materials

335

(a) heating stage (room temperature100 μm), let alone their high costs. The second is SiC fibers manufactured using organic polymers, the high temperature performance of fibers is affected by the existance of impurities such as oxygen and free carbon, resulting in a significant reduction in strength above 1000 °C. Silicon nitride fibers are actually multiphase fibers composed of Si, N, C, and O. These fibers are also manufactured using organic polymers and their properties are close to that of silicon carbide fibers. Besides, they have the same problems as silicon carbide fibers. Carbon fiber is the best developed fibers and have the optimal performance, they have been widely used as reinforcing fibers for composites. The high-temperature performance of carbon fibers is also very outstanding, they can maintain their strength without loss in an inert atmosphere at temperatures above 2000 °C, and surely they have the best high-temperature performance among the reinforcing fibers. However, the biggest problem carbon fiber faced with is the poor oxidation resistance at high temperature. weight loss in oxidation and strength reduction occur above 360 °C if the fibers are exposed to the air. The mentioned problem can be solved by the surface coating technology, so coated carbon fibers are the best candidates for reinforcement of continuous fiber reinforced ceramic matrix composites. Currently, the commonly used matrix materials for ceramic matrix composites are Si3 N4 , SiC, ZrO2 , Al2 O3 , etc. In addition, the newly developed high-performance nanocomposite ceramics are also a class of promising composites. (2) Joining characteristics of ceramic matrix composites The joining of ceramic matrix composites has the same difficulties with ceramics. For example, their high melting point and high temperature decomposition of some ceramics make it hard to fusion welding of them, the electrical insulation of ceramics makes it impossible to be joined by arc or resistance welding, the inherent brittleness of ceramics makes it impossible to withstand the thermal stress of welding, the poor plasticity and toughness of ceramics makes it unable to bear great pressure for solidphase bonding, the ceramics are difficult to be brazed because of the poor wettability due to their chemical inertness, etc. Joining of ceramic matrix composites should be done with the following in mind. 1) The selection of joining methods and consumable materials should give full consideration to their adaptability to the matrix materials and reinforcing materials. 2) Unfavorable interfacial reactions between the reinforcing phase and the matrix should be avoided, which may cause oxidation and properties degradation of the reinforcing phase (e.g., fibers) etc. Therefore, the joining temperature and the holding time should be limited to a certain level. For example, when joining

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385

SiCf /SiC composites using Si intermediate layer at 1425 °C, the SiC performance was severely reduced with the holding time of 45 min, while the performance of the matrix stayed unaffected with the holding time reduced to 1 min. 3) Due to the poor or restricted pressure resistance of the fiber reinforced ceramic matrix composite, high pressure cannot be applied during the joining process. The commonly used joining methods for ceramic matrix composites are brazing, pressureless solid-phase reaction bonding, transient liquid phase diffusion welding, microwave joining, etc. Brazing of ceramic matrix composites is basically the same as that of ceramics, which can be active brazed using fillers containing Ti, Zr and other elements. It can also be performed by metalizing the composite surface firstly and then to be brazed using common fillers. The pressureless solid-phase reaction bonding of ceramic composites was performed via formation of compounds generated by reactions between active elements with the ceramic matrix at high temperature, and high pressure cannot be applied during the process. Dense joints with high temperature resistance can be obtained using this joining method, but the joints have low mechanical properties and cannot be applied to loads.

7.3 Welding of Continuous Fiber Reinforced Metal Matrix Composites 7.3.1 Problems in Welding of Continuous Fiber Reinforced MMC Continuous fiber reinforced metal matrix composites (MMC) consist of a base metal and reinforcing fibers.The welding of such materials involves the welding between metal matrixs, the welding between metal and non-metal reinforcing phases, as well as the welding between reinforcing phases. The matrix is usually ductile metal with good plasticity and weldability. While the reinforcing phase is non-metal fiber with high strength, high modulus, high melting point, low density and low linear expansion coefficient, and its weldability is poor. Therefore, the weldability of fiber reinforced metal matrix composites is also poor, and the main problems encountered during welding of such materials are as follows. (1) Interfacial reactions The matrix and the reinforcing phase of metal matrix composites are usually in a unstable thermodynamic state and chemical reactions between the two are prone to happen at their contact interface at high temperatures. The generated brittle phases are usually harmful to the material properties. Preventing or suppressing of interfacial reactions and the generation of brittle phases is one of the key points to ensure the welding quality, the problem can be solved by both metallurgical and technological measures.

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7 Welding of Advanced Composites

1) Metallurgical measures. The interfacial reaction can be prevented by adding some elements that are more active than the base metal or elements that can prevent the interfacial reaction. For example, the Al was replaced by the added Ti to react with SiC when welding of SiCp /Al composites, the foramtion of harmful compounds Al4 C3 is avoided, and the generated TiC particles can act as reinforcing phases. The interfacial reaction between Al matrix and SiC can be suppressed by increasing the Si content in the Al matrix or using welding wire with high Si content. 2) Technological measures. To suppress the interfacial reaction by controlling the heating temperature and welding time. For example, a solid phase welding process or a fusion welding process with low heat input can be used to suppress the interfacial reaction of SiCf /Al composites. (2) Poor fluidity of the molten pool and poor wettability of the base metal to the fibers The melting point difference between the metal matrix and the reinforcing fibers is large. When using the fusion welding, the presence of a large number of solid fibers in the molten pool impedes the flow of liquid metal, it can easily lead to welding defects such as porosity, incomplete penetration and lack of fusion. (3) High residual stress in the joint During the heating and cooling processes of welding, large internal stress forms easily near the interface due to the large difference of linear expansion coefficient between the reinforcing fibers and the matrix. The residual stress is liable to cause interfacial peeling of the bonding interface. Therefore, these materials have high hot cracking susceptibility. (4) Unfavorable changes of the fibers distribution state During diffusion welding or pressure welding, the reinforcing fibers will break if the pressure is too high. The fibers at the interface between weldments are almost impossible to butt joining, and the reinforcing fibers are discontinuous at the bonding area. The factors mentiond above result in a much lower strength and stiffness of the joint than the composite.

7.3.2 Joint Form Design for Continuous Fiber Reinforced MMC The discontinuity of fibers in the joints affects the strength and stiffness of fiber reinforced metal matrix composites. Therefore, the joint form must be designed properly to improve the performance of the joint. When a lap joint is used, the strength of the joint can be improved by adjusting the lap area. The joint strength increases with the increasing of lap area. When the lap area

7.3 Welding of Continuous Fiber Reinforced Metal Matrix Composites

387

Fig. 7.7 Proper joint forms for continuous fiber reinforced metal matrix composites

is increased to a certain value, the joint strength can reach that of the base material. However, the lap joint increases the mass of the welded structure and the form of the joint is discontinuous, so the applications of this joint form are limited. The ideal joint form is the butt joint with stepped or bevel grooves, which is characterized by the dispersion of discontinuous fibers into different cross sections. The number of steps and the bevel angle can be designed according to the force states on the workpiece. To ensure the continuity of the reinforcing fibers, Examples for proper design of the joint form are shown in Fig. 7.7 (d) and (e).

7.3.3 Characteristics of Welding Process for Fiber Reinforced MMC The main welding methods applicable to fiber reinforced metal matrix composites (MMC) are arc welding, laser beam welding, diffusion welding, brazing, etc. The friction welding is not suitable for welding of fiber reinforced metal matrix composites, because the process causes large plastic deformation of the compsites at the bonding interface. Table 7.12 gives examples of commonly used welding methods for fiber reinforced MMC and the strength of the joints. (1) Arc welding Arc welding has also gained considerations in the welding of metal matrix composites. Only butt joints and lap joints can be used in arc welding. The main problems of using arc welding methods including the bring about interfacial reactions, causing breakage of fibers, etc. In order to prevent interfacial reactions, the welding process is usually performed by pulsed tungsten arc welding (P-GTAW). The interfacial reactions can be suppressed by strictly controlling of the welding heat input and shortening the existence time of molten pool. By adding appropriate filler wire, the direct action of the arc on the fibers can be reduced and the damage to the fibers can be reduced.

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Table 7.12 Examples of common welding methods and joint strengths for composite materials The joint

Welding method

Joint form

Joint strength/MPa

Notes

Bf /Al joint

Brazing

Lap joint, Double strapped butt joint, Bevelled joint, Double pronged strapped butt joint

590 820 640 320

—

Dissimilar joint of Bf /Al with Ti-6Al-6 V-2Sn

Brazing

Double lapped joint 496

—

Dissimilar joint of SiCf /Al with Al

Diffusion welding

Butt joint

60

—

SiCf /Al joint

Diffusion welding

Butt joint

60

—

CO2 laser beam welding

Overlaying

—

—

Nicalon SiCf /Al joint

Diffusion welding

Lap joint

96

Shear strength

Cf /Al joint

CO2 laser beam welding

Overlaying

—

—

GTAW

Butt joint

—

—

Brazing

Lap joint

—

—

Resistance spot welding

Lap joint

—

—

Nb-Ti/Cu joint

Diffusion welding

bevelled joint

300

—

SiCf /Ti joint

Laser beam welding

Butt joint

550

—

Diffusion welding

Butt joint

850

—

12° bevelled joint

1380

—

Double strapped butt joint

1300

—

—

850–991

—

Dissimilar joint of SiCf /Ti with Ti-6Al-4 V

Laser beam welding

(2) Laser beam welding As a high power density welding method, laser beam welding has both advantages and disadvantages when welding fiber reinforced composites. The advantages of the laser beam welding of fiber reinforced composites are in the following. 1) The heat affected zone can be controlled in a small scale and the existence time of molten pool can be greately shortened.

7.3 Welding of Continuous Fiber Reinforced Metal Matrix Composites

389

2) The mechanical impact the fibers suffered is low without the direct irradiation from laser beam. So breakage and displacement of fibers can be prevented by controlling the position of the laser beam. The disadvantage in laser beam welding of composites is that the reinforcing phases with high resistivity is heated preferentially thanks to the ultrahigh temperature of the molten pool, it easily leads to melting, dissolution and sublimation of the reinforcing phases, as well as interfacial reactions between the matrix and the reinforcing phase. So the laser beam welding is not suitable for composites prone to interfacial reaction, such as Cf /Al and SiCf /Al. The method can only weld some composites who have good chemical compatibility between reinforcing phase and matrix, such as SiCf /Ti, etc. The key to laser beam welding of fiber reinforced metal matrix composites is to strictly control the laser beam position so that the fibers are outside the laser beam irradiation range, i.e., outside the “keyhole” of the molten pool. For example, when laser beam welding of SiCf /Ti-6Al-4V composites to Ti-6Al-4V titanium alloy, the laser beam should be properly offset towards the titanium side, as shown in Fig. 7.8a, so that the SiC fibers are outside the keyhole of the molten pool. During laser beam welding of SiCf /Ti-6Al-4V composites, a Ti-6Al-4V intermediate layer with a thickness of about twice of the keyhole diameter (approximately 300 μm) should be sandwiched between the composites to be joined. So that the fibers of composites at both sides are outside the keyhole (Fig. 7.8b). The composites at both sides are melted by heat conduction and then the molten materials are mixed together to form the joint. It has been reported that even such measures being taken, the SiC fibers in the molten pool can still react with the liquid titanium. However, the reactions can be limited to a low degree due to the short residence time of the molten pool.

Fig. 7.8 Schematic diagram of the laser beam position

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7 Welding of Advanced Composites

The strength of the laser welded dissimilar joint of SiCf /Ti-6Al-4V composite with Ti-6Al-4V titanium alloy mainly depends on the welding parameters and the distance (X) between the laser beam center and the edge of the composite. With fixed welding parameters, there is an optimum distance X* at which the joint tensile strength reaches its maximum value, as shown in Fig. 7.9. When X < X*, the joint strength decreases with the decreasing of the distance X due to the escalating damage of SiC fibers and the segregation of C and Si near the fibers. When X > X*, the joint strength decreases with the increasing of the distance X. It is thought that the decrease of strength values has relation with the massive formation of grain boundaries at the bonding interface and the lack of fusion defect. As shown in Fig. 7.9, during CO2 laser beam welding with the laser power of 1.5 kW and the welding speed of 50 mm/s, the joint tensile strength reaches a maximum of 991 MPa when the X = 250 μm. The joint tensile strength is higher than 850 MPa when the X is in the range of 225 μm to 280 μm. For joints with X outside this range, the joint tensile strength can be improved by post-weld heat treatment (holding at 900 °C for 60 min), the X range at which the joint tensile strength exceeds 850 MPa has been extended to 190–310 μm. The main reasons for the strength improvement are different at different cases. For joints with smaller X, the segregation of C and Si near the damaged fibers disappeared due to the heat treatment. For joints with larger X, the heat treatment causes the migration of grain boundaries at the bonding interface. With fixed thickness of the intermediate layer, the joint strength of SiCf /Ti-6Al4V composite depends mainly on the laser power, and there is an optimum laser power at which the joint has the maximum strength value. With insufficient laser power, the intermediate layer at the bottom of the weld can not be completely melted

Fig. 7.9 Effect of laser beam position on the tensile strength of Ti-6Al-4V and SiCf /Ti-6Al-4V dissimilar joints

7.3 Welding of Continuous Fiber Reinforced Metal Matrix Composites

391

and therefore the joint strength is influenced. If the laser power is too high, a large amount of brittle phases will form due to the dramatically increased interfacial reactions between the fibers and the matrix, which would result in reduction of the joint strength. (3) Diffusion welding The weldments stay in solid state during diffusion welding, the erosive effect of the molten metal on the reinforcing fibers is avoided, so the diffusion welding is considered as one of the best joining methods for fiber reinforced metal matrix composites. However, there are still some problems exist in the diffusion welding of fibre-reinforced metal matrix composites, the main problems are as follows. ➀ Interactions between the fibres and the matrix may occur due to the long holding time of diffusion welding. ➁ The reinforcing fibers have high strength and stiffness. Deformation and close contact of the weld surfaces may be impeded by the direct contact of fibres in the two weld surfaces, making diffusion welding difficult to achieve. ➂ During diffusion welding of the composite to its corresponding base metal, excessive deformation will happen at the base metal side. ➃ For diffusion welding of fibre-reinforced metal matrix composites, The joint strength mainly depends on the bonding strength between the metal matrix at the bonding interface, so the joint strength is directly proportional to the area percentage of matrix at the bonding interfaces. Conversely, the larger the percentage of fibres, the lower the strength of the joint. In other words, the larger the volume fraction of fibers in the composite, the worse the composite weldability. 1) Selection of welding temperature and holding time The chosen diffusion welding temperature and time should ensure that no significant interfacial reactions occur. The principles for the selection of welding parameters are discussed below using the example of diffusion welding of SiC(SCS-6)f /Ti-6Al4V composites. SCS-6 is a SiC fiber surface coated with a C-rich layer of 3 μm in thickness and its diameter is approximately 140 μm. The fiber is specifically designed for reinforcing titanium matrix composites. The relationship between the thickness of the interfacial reaction layer and the holding time at different welding temperatures in duffusion wedling of SiC (SCS6)f /Ti-6Al-4V composites is shown in Fig. 7.10. It can be seen that the higher the welding temperature, the higher the increasing rate of the reaction layer thickness. However, after holding for a certain time, the increasing rate of the reaction layer thickness becomes slower. So the reactions between SiC (SCS-6) fibers and titanium matrix can be divided into two stages. According to thermodynamic analysis, the reactions between SiC (SCS-6) fibers and titanium matrix that occur at high temperatures are as follows. Ti + C = TiC This is the reaction that occurs in the first stage of, which relies on the diffusion of Ti or C. Since C diffuses much more rapidly in TiC than in Ti, C continuously

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7 Welding of Advanced Composites

Fig. 7.10 SiC (SCS-6)f /Ti-6Al-4 V interfacial reaction layer thickness versus holding time t

diffuses outward through the resulting TiC layer and reacts further with the titanium matrix until the C-rich layer on the fiber surface is completely depleted. And then two reactions with smaller gibbs free energy changes then proceed. 9Ti + 4SiC = 4TiC + Ti5 Si4 8Ti + 3SiC = 3TiC + Ti5 Si3 This is the reactions happen in the second stage, the reactants are two silicides and TiC. For these two reactions, Ti must firstly passes through the previously formed reaction layer before it reacts with SiC. Therefore, the rates of these two reactions are at a low degree due to the previously formed reaction layer and the low diffusion rate of Ti. The tensile strength of the SiC/Ti-6Al-4V composite joint decreased dramatically when the thickness of the reaction layer exceeded 1.0 μm. The required times for the reaction layer thickness to reach 1.0 μm at different temperatures are given in Fig. 7.11. For diffusion welding of SiC/Ti-6Al-4V composites, the parameter point formed by the welding temperature and holding time should lie below the curve shown in Fig. 7.11. 2) Intermediate layer and welding pressure During welding of SiC/Ti-6Al-4V composite to Ti-6Al-4V titanium, deformation and plastic flow are easy for that there is no direct contact of fibers at the bonding interface. So the bonding can be obtained more easily using both direct diffusion welding and transient liquid phase diffusion welding. However, when using direct diffusion welding, the unilateral deformation of the Ti alloy is too large thanks to the high required pressure. While with transient liquid phase diffusion welding, the

7.3 Welding of Continuous Fiber Reinforced Metal Matrix Composites

393

Fig. 7.11 Required times for the reaction layer thickness to reach 1.0 μm at different temperatures

welding pressure is much lower and the unilateral deformation of Ti alloy is smaller. For example, in order to achieve a joint strength of 850 MPa, the required welding pressure in direct diffusion welding is 7 MPa and the welding time is 180 min. While in transient liquid phase diffusion welding using Ti-Cu-Zr alloy as the intermediate layer, the required welding pressure is only 1 MPa, the welding time is 30 min. Besides, the unilateral deformation of titanium alloy reduced from 5% of the direct diffusion welding to 2% of the transient liquid phase diffusion welding. Direct diffusion welding of fiber reinforced metal matrix composites is very difficult, because the plastic deformation and close contact of the weld surfaces may be impeded by the direct contact of high strength and stiffness fibres in the two weld surfaces. To overcome the problem, an intermediate layer should be inserted into the interface between the composites being welded, so that direct fiber-to-fiber contact is avoided. The weldability for transient liquid phase diffusion welding of fiber reinforced metal matrix composites is also poor. The overall strength of joint muanufactured by transient liquid phase diffusion welding is still low due to the poor bonding between the fibers and the matrix, despite the good bonding between the matrix metals. Generally, in addition to the selected transient liquid phase layer, matrix metal foil with appropriate thickness should be also added into the interface as an additional intermediate layer. Figure 7.12 shows the schematic diagram of transient liquid phase diffusion welding of SiCf /Ti-6Al-4V composites using Ti-6Al-4 V foil as the intermediate layer and using Ti-Cu-Zr foils as the transient liquid phase layers. During transient liquid phase diffusion welding of SiCf -30%/Ti-6Al-4V composites, the effect of intermediate layer (titanium foil) thickness on the joint strength is shown in Fig. 7.13. It can be seen that the joint tensile strength reaches 850 MPa when the thickness of intermediate layer exceeds 80 μm, which is equal to the joint strength between SiCf -30%/Ti-6Al-4V composite and Ti-6Al-4V titanium alloy. In fact, once the thickness of Ti-6Al-4V intermediate layer reaches to a certain degree, the welding

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of the composites using welding method mentioned above changes into the welding between the SiCf /Ti-6Al-4V composite and the Ti-6Al-4V titanium alloy. The feature differ from the dissimilar welding is that two dissimilar joints are to be weldeded simultaneously. 3) Optimal design of joint forms The joint form has an important effect on the joint strength. To improve the joint strength of fiber reinforced metal matrix composites, the joint form can be designed as a miter joint. Figure 7.14a shows a schematic diagram of a diffusion-welded miter joint with an intermediate layer. At a joint strength factor of about 80%, the fracture starts at the area with discontinuous SiC fibers at the joint interface (point A in Fig. 7.14b). After initiation, the crack propagates perpendicular to the tensile direction through the entire composite section. The joint strength can not reach the strength of composite due to the discontinuity of the fibers at the bonding interface and the reinforcing effect of the fibers at the interface is greatly reduced. So craks initiate easily under the action of internal stress. Fig. 7.12 Schematic diagram of transient liquid phase diffusion welding using both intermediate layer and transient liquid phase layer

Fig. 7.13 Effect of intermediate layer thickness on the joint strength of SiCf -30%/Ti-6Al-4V composites

7.3 Welding of Continuous Fiber Reinforced Metal Matrix Composites

395

Fig. 7.14 SiC with interlayerf -30%/Ti-6Al-4V diffusion welded miter joint and fracture process

(4) Brazing For metal matrix composites, the brazing temperature is relatively lower, the base metal stays unmelted in the brazing process and it is less likely to cause interfacial reactions. By selecting suitable filler materials, the brazing temperature can be reduced below the temperature at which the fiber properties begin to deteriorate. Lap joint is the dominant joint form used in composite brazing, which simplifies the joining of composites into the joining of the metal matrix itself. So brazing is more suitable for welding of composites and it has become one of the main methods for welding metal matrix composites. 1) Brazing of fiber reinforced aluminum matrix composites ➀ Brazing. In the 1970s, Bf /Al composites were successfully joined using brazing techniques, the joined parts were used to manufacture reinforcing ribs on aircraft. The brazing temperature is in the range of 577°C–616 °C when using Al-Si, Al– Si–Mg and other brazing fillers. Thus Bf /Al composite is not suitable for brazing because the interfacial reaction between B and Al may become violent at 550 °C. the joint strength would greatly reduced by the formation of brittle phase AlB2 . But using the same process to braze B fiber reinforced Al matrix composite (Borsic/Al), in which the B fiber surface was coated with a SiC layer of 0.01 mm in thickness, the interfacial reactions can be completely avoided. It is because the B fiber is protected by the SiC coating due to the high reaction temperature between SiC and Al (593°C–608 °C). Brazing can be done by both vacuum brazing and dip brazing processes. The joint strength of dip brazing is relatively higher (fracture strength for Tjoint reaches 310–450 MPa), but the corrosion resistance of the joint is poor. The joint strength of vacuum brazing is relatively lower (fracture strength for T-joint is 235–280 MPa), while the corrosion resistance of the joint is much better. Single-layer Borsic/Al composite foils can be manufactured into multilayered plates or multilayered sectional materials using vacuum brazing. For example, a AlSi filler foil is sandwiched between Borsic/Al composite foils, then the structure is vacuum brazed with brazing temperature of 577°C–616 °C and applied pressure of 1030–1380 Pa, a multilayered plate can be obtained after holding for a certain time. The tensile strength of Borsic-45% (45% fiber volume fraction)/Al multilayered

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7 Welding of Advanced Composites

composites made by this method is 978–1290 MPa. Structural parts with complex cross-sections are more suitable for brazing using hot isostatic apparatus. The pressure required for vacuum brazing is lower than that required for diffusion brazing. Compared to diffusion welding, the strength of Bf /Al composite brazed joints is reduced by about 20% to 30%, but the cost is also lower. When joining SiCf /Al composites using brazing, there exists an optimum brazing temperature at which the joint has the highest strength. When the brazing temperature is lower than the optimum temperature, the joint is fractured in the weld. While the joint is broken in the base materials when the brazing temperature is higher than the optimum temperature. It suggests that although no interfacial reaction between SiC and Al occurs during the brazing, thermal cycle of the brazing has an effect on the properties of the composite. ➁ Soldering. Soldering of Bf /Al or Borsic/Al composites can be performed with three kinds of fillers 95% Zn-5% Al, 95% Cd-5% Ag and 82.5% Cd-17.5% Zn, the melting temperatures of these fillers are 383 °C, 400 °C and 265 °C, respectively. For soldering of composites, the surface treatment of the composites has a significant effect on the joint strength. In soldering of Bf /Al composites, the wettability between the solid composite and liquid filler is significantly improved by plating a Ni layer with a thickness of 0.05 mm on the composite surface, the joint strength is also increased. Compared with using electroplating, the joint strength was increased by 10% to 30% when chemical plating was used. It is because the exposed B-fibers are nonconductive and Ni cannot be reliably plated onto the B-fibers using electroplating. So the wettability between the filler and the B-fibers is still poor. The problem will not arise when chemical plating is used. Table 7.13 gives the shear strength of the dissimilar joints between Bf /Al composites and 6061Al (T6) aluminium alloy soldered using these three fillers, the flux soldering process was carried out using oxy-acetylene flame. Table 7.13 Mechanical properties of Bf /Al soldering joints Filler compositions

Shear strength/MPa

Test temperature/°C

Failure mode

95%Cd-5%Ag

81 89 69 47 29 5.6

294 366 422 478 533 588

1 1 1 3 2 2

95%Zn-%Al

80 94 30

294 366 588

1 1 3

82.5%Cd-17.5%Zn

74 90 59

294 366 422

1 1 3

Note 1—composite interlaminar shear; 2—fracture from the brazed seam; 3—both 1 and 2 occur

7.3 Welding of Continuous Fiber Reinforced Metal Matrix Composites

397

The joint soldered using 95% Zn-5% Al filler has high high-temperature strength and is suitable for working at 316 °C, however, the soldering process is difficult to control. The joint soldered using 95% Cd-5% Ag filler has high low-temperature strength (below 93 °C), the appearance of weld is excellent and the soldering process is easy to control. The joint soldered using 82.5% Cd-17.5% Zn filler is very brittle and fracture may occur during the cooling process. ➂ Eutectic diffusion brazing. The process of eutectic diffusion brazing starts with coating of an intermediate layer on the weld surfaces or the addition of intermediate material foil into the bonding interface. The weld assembly is heated to a suitable temperature so that a liquid low-melting point eutectic layer could form due to the interdiffusion between base material and the intermediate layer. And eventually, a joint with homogenous composition is formed after isothermal solidification and diffusion homogenization and other processes. Suitable intermediate materials for eutectic diffusion brazing of Al matrix composites include Ag, Cu, Mg, Ge and Zn, etc. The thickness of the intermediate layer should be controlled at about 1 μm. Compared with eutectic diffusion brazing of a single metal material, when eutectic diffusion brazing of composites, the rate of diffusion homogenization is drastically reduced because the free diffusion of the intermediate layer elements into the metal matrix is hindered by the reinforcing fibers. And it is difficult to completely eliminate the brittle layers in the joint by diffusion. Therefore, it is very important to control the thickness of the intermediate layer, the time for diffusion homogenization should also be extended appropriately to avoid severe degrading of the joint properties. Bf -45%/1100Al composites were eutectic diffusion brazed using Cu foil of 1.0 μm thickness, brazing temperature of slightly above 548 °C, homogenized temperature of 504 °C and holding time of 2 h. The interdiffusion between Cu and Al gradually occurred during the heating process, eutectic liquid phase (Al-Cu 33.2%) was formed once the temperature exceeded 548 °C. Then the holding process was performed, Cu continuously diffused into the matrix Al with the holding process proceeded, and isothermally solidification happened when the concentration of Cu in the eutectic liquid phase dropped below 5.65%. The inhomogeneous distribution of Cu in the joint was then further reduced after a homogenization process of 504 °C × 2 h. The tensile strength of the joint obtained by this method was 1103 MPa, and the effective coefficient of joint strength reached 86%. The homogenization process is much easier using Ag interlayer than that using the Cu interlayer, and the joint performance was somewhat higher. 2) Brazing of fiber reinforced titanium matrix composites Brazing thermal cycle generally does no damage to the properties of titanium matrix composites. The commonly used fillers are Ti-Cu15-Ni15 and Ti-Cul5 amorphous materials, a composite filler that roll bonded using a 50% Cu-50% Ni alloy sandwiched by two pieces of dull titanium can also be utilized. The brazing temperature is higher and the holding time is longer when using composite fillers, so the thickness of the diffusion layer is a bit larger. The SCS-6/β21S composite was successfully brazed using Ti-Cu15-Ni15 filler and a composite filler that roll bonded using a 50% Cu-50% Ni alloy sandwiched

398

7 Welding of Advanced Composites

by two pieces of dull titanium. β21S is a titanium alloy with compositions of TiMo15-Nb2.7-Al3-Si0.25. The tensile test results at both room temperature and high temperatures (649 °C, 816 °C) showed that the tensile property of the SCS-6/b21S composite was not affected the brazing process. Eutectic diffusion brazing of CSC-6/b21S composites using 17 μm thick Ti-Cu15 filler was carried out by a rapid infrared brazing process. The process was performed in an infrared furnace that filled with Ar, the heating rate is 50 °C/s. The thickness of the reaction layer was 0.19 μm, 0.44 μm, and 0.62 μm when heated at 1100 °C for 30 s, 120 s, and 300 s, respectively. The isothermal solidification was unfinished when the heating time was 30 s, while the diffusion homogenized was completely finished after heating for 120 s. Therefore, the ideal brazing temperature and holding time are 1100 °C and 120 s, respectively. High temperature shear tests (at 650 °C and 815 °C) were performed on joints that brazed using the ideal parameters. The results showed that all the joints fractured at other locations instead of the bonding interface. (5) Resistance welding Joining of fiber reinforced metal matrix composites using resistance welding has advantages of short heating time, good controllability, avoiding of interfacial reactions and also preventing cracks and porosity by applying pressure. By using lap joints, the welding of fiber reinforced metal matrix composites can be changed into the welding of Al matrix, so this method is suitable for welding of fiber reinforced metal matrix composites. However, the selection of welding parameters and the control of welding quality are difficult due to the complexity of the current distribution and the electrode pressure distribution, which is caused by the reinforcing phases. The main problems exist in resistance welding of fiber reinforced metal matrix composites are the fracture of the fibers and the spatters formed by the molten matrix during welding. To prevent fracture of the fibers, the electrode pressure should be reduced as much as possible. But if the electrode pressure is too small, spatters will be produced from the molten matrix at the bonding interface. So strict control of the electrode pressure is required. The welding heat input should also be minimized during the welding. Because if the heat input is too large, the fibers will be damaged, the molten matrix at the bonding interface will also splashed out, which may causes exposing of the fibers and deterioration of the bonding strength. In addition, delamination defects in fiber reinforced metal matrix composites may also lead to spattering. So the ultrasonic inspection should be performed before the welding and the welding spot shoul be locates at a delamination-free area. During dissimilar welding of fiber reinforced metal matrix composites to its matrix metals, the weld nugget is easily offset towards the composite side due to the high resistivity and low linear expansion coefficient of the composite. Therefore, the electrodes should be matched properly to keep the weld nugget in the middle position. The electrode with a smaller contact area and higher resistivity should be used at the matrix metal side; the electrode with a larger contact area and lower resistivity should be used at the composite side.

7.4 Welding of Discontinuously Reinforced Metal Matrix Composites

399

The volume fraction of the reinforcing fibers has a great influence on the weldability during resistance welding. The fluidity of the molten pool becomes poor with the increasing of the fiber volume fraction, which may causes the reduction of the joint strength. For example, when the fiber volume fraction increases from 35% to 50%, the joint strength decreases by about 10%.

7.4 Welding of Discontinuously Reinforced Metal Matrix Composites The applications of continuous fiber reinforced metal matrix composites are limited to a few fields such as aerospace and military industries due to the complex manufacturing process and high cost. Discontinuously reinforced metal matrix composites maintain most of the excellent properties of the components. In recent years the development of discontinuously reinforced metal matrix composites is rapid due to their simple manufacturing processes, low cost of raw materials and the convenient secondary processing. The weldability of these composites, although better than that of continuous fiber reinforced metal matrix composites, is still very difficult compared to the welding of single component metals and alloys. The common discontinuously reinforced metal matrix composites mainly are SiCp /Al, SiCw /Al, Al2 O3p /Al, Al2 O3sf /Al and B4 Cp /Al.

7.4.1 Welding Problems of Discontinuously Reinforced MMC According to their performance characteristics, the following problems may exist in welding of discontinuously reinforced metal matrix composites. (1) Interfacial reactions Interfacial reactions between the matrix and the reinforcing phases exist in most of the metal matrix composites (MMC) at high temperatures. The formation of some brittle compounds at the interface may causes reduction the composite performance. Particles or short fibers made by Al2 O3 do not react with Al at any temperature, and therefore Al2 O3 reinforced aluminum matrix composites have good chemical compatibility. SiC does not react with solid Al but liquid Al, the reaction equation between SiC particles and liquid Al is described as follows. 4Al (liquid) + 3SiC (solid) → Al4 C3 (solid) + 3Si (solid) The free energy of the reaction is. ∆G = 11390 − 12.06T lnT + 8.92 × 10−3 T 2 + 7.53 × 10−4 T −1 + 2.15T + 3RT lnα[Si]

400

7 Welding of Advanced Composites

where α [Si] is the chemical activity of Si in liquid Al. The reaction above consumes the SiC reinforcing phase of the composite, also the brittlement of the joint is dramatically increased by the generated Al4 C3 . Therefore, prevention of interfacial reactions is the primary issue to be considered in welding of such composites. Methods to prevent or suppress interfacial reactions are listed below. 1) To increase the Si content in the molten pool by using Al alloy with high Si content as the matrix or wire with high Si content as the filler. According to the free energy equation of the reaction, the driving force of the reaction (−ΔG) decreases and the interfacial reaction is weakened or suppressed with the increasing of Si activity. 2) By using the welding methods with low heat input and controlling the welding heat input strictly to reduce the temperature of the molten pool and shorten the contact time between liquid Al and SiC. 3) Increase the bevel angle (or size) of the joint to reduce the melting of the composite, which equivalently, reduce the amount of SiC entering the molten pool. 4) Special fillers may also be used, which should contain active elements that have a stronger binding capacity with C than Al and do not produce harmful carbides, such as Ti. Ti will replace Al to reacts with SiC if Ti is added into the molten pool. The generated TiC tiny particles do no harm to the weld performance but act as reinforcing phases. The interfaces of composites such as Al2 O3 /Al, B4 C/Al are relatively stable and are generally less susceptible to interfacial reactions. (2) High viscosity and poor fluidity of the molten pool The unmelted reinforcing phases in the molten pool increase the viscosity of the molten pool while reduce the fluidity of the molten pool. Which, further more, increasing the susceptibility to defects such as porosity, cracking, and unfusion. The fluidity of the molten pool can be improved by using wire with high Si content or increasing the bevel size (reducing the amount of SiC or Al2 O3 reinforcing phase in the molten pool). The use of high Si wire improves the wettability of the molten metal to SiC particles. The use of high Mg wire improves the wettability of the molten metal to Al2 O3 and prevents agglomeration of particles. (3) High susceptibility to porosity and crystallization cracks The hydrogen content in metal matrix composites, especially those made by powder metallurgy, is high. During the welding, the gas is difficult to escape due to the high viscosity of the molten pool and therefore the porosity sensitivity is high. Generally, the composites should be vacuum dehydrogenated before welding to avoid porosity. In addition, linear expansion coefficient of the weld is different to that of the composite, so high residual stress exists in the weld, which further aggravates the susceptibility to crystallization cracking.

7.4 Welding of Discontinuously Reinforced Metal Matrix Composites

401

(4) Segregation of the reinforcing phases, discontinuity of the joint Segregation of the reinforcing phases is easy to occur after remelting and it leads to uneven distribution of reinforcing phases in the weld, the reinforcing effect of reinforcing phases is reduced. Until recently, there are no special filler wires used for composite materials. The selection of filler wires in arc welding of composites is generally based on their matirx metals, it causes great reduction of reinforcing phase content in the weld and material discontinuity of the joint. Even if the problems mentioned above are avoided, it is still difficult to achieve high strength welding of composites that equal to base materials.

7.4.2 Welding Process Characteristics of Discontinuously Reinforced MMC Table 7.14 gives the advantages and disadvantages of the three kinds of welding methods (fusion welding, solid phase welding, and brazing) that can be used to weld discontinuous fiber reinforced metal matrix composites. (1) Arc welding Both Non consumable electrode arc welding and consumable electrode arc welding (GTAW, GMAW) can be used to welding of discontinuously reinforced metal matrix composites (MMC). Improper selection of heat input when welding SiCp /Al or SiCw /Al composites can cause severe interfacial reactions and the formation of needle-like Al4 C3 . Therefore, pulsed current argon arc welding (GTAW, GMAW) is preferred to reduce wleding heat input and attenuate or suppress interfacial reactions. The pulsed arc has stirring effect on the molten pool, which can partially improve the fluidity of the molten pool, the particle distribution in the weld and the crystallization conditions. The weldability of SiCp /Al or SiCw /Al composites differs significantly when the base metal is different. The higher the Si content of base metal, the weaker the interfacial reactions, the better the fluidity of the molten pool, and the lower the susceptibility to cracking and porosity. When the Si content of the base metal is low, a wire with higher Si content should be used for welding to suppress interfacial reactions and improve the joint strength. SiCp /Al or SiCw /Al has a high porosity sensitivity, and a large number of hydrogen pores are easily formed in the weld and the heat affected zone. More serious, stratiform distribution of pores may formed. So the material must be vacuum dehydrogenated before welding. The treatment process is carried out with a vacuum degree of 10–2 –10–4 Pa, heating temperature of 500 °C, holding time of 24–48 h. Unlike SiCp /Al composites, the interfacial reactions between the reinforcing phase and liquid Al does not happen when welding Al2 O3p /Al composites by arc welding. The main problems existing are the high viscosity but poor fluidity of the molten pool, as well as the poor wettability of molten metal to the Al2 O3 reinforcing phase.

402

7 Welding of Advanced Composites

Table 7.14 Advantages and disadvantages of various welding methods for composite welding Advantages

Disadvantages

TIG welding

(1) Can suppress interfacial reactions and improve wettability of molten metal to reinforcing phase by selecting proper filler wires (2) Low cost, easy to operate and wide application

1) High possibility of interfacial reactions 2) Low volume fraction of reinforcing phase in the weld and low joint strength when using single material filler wire 3) High susceptibility to porosity

MIG welding

Ditto

Welding methods Fusion welding

Solid phase welding

Ditto

Vacuum electron beam 1) Low susceptibility welding to porosity 2) Even distribution of reinforcing phases in the weld 3) High welding speed

1) Interfacial reactions occur when the welding parameters are improper 2) High cost

Laser beam welding

Low susceptibility to porosity, high welding speed

Difficult to avoid interfacial reactions

Resistance spot welding

Short heating time, small nugget volume and high welding speed

High susceptibility to segregation of reinforcing phases

Solid state diffusion welding

1) Can improve joint performance and prevent interfacial reactions using proper intermediate layer 2) Can used in dissimilar welding

Low productivity, high cost, difficult to select parameters

Transient liquid phase diffusion welding

Ditto

Ditto

Friction welding

1) Can fabricate Only can welding parts joints with equal with small size and strength to base simple shapes materials by postweld heat treatments 2) can joining dissimilar materials 3) No interfacial reactions (continued)

7.4 Welding of Discontinuously Reinforced Metal Matrix Composites

403

Table 7.14 (continued) Welding methods

Advantages

Disadvantages

Brazing & Soldering

1) Low heating temperature and low susceptibility to interfacial reactions 2) Can joining dissimilar materials or complex components

Conducted in inert atmosphere or vacuum, postweld heat treatments are needed

The use of filler materials with higher Mg content can increase the fluidity of molten pool and improve the wettability of the molten metal to the Al2 O3 reinforcing phase. Examples of welding parameters and joint properties for a few kinds of discontinuously reinforced metal matrix composites are listed in Table 7.15. (2) Brazing Not all fillers that are used to braze aluminum alloys can be used to braze aluminum matrix composites, because brazing of aluminum matrix composites requires good wettability of molten filler both to the base metal and to the reinforcing particles or whiskers. Moreover, the brazing temperature should be controlled as low as possible to avoid the adverse effect of brazing thermal cycle on the reinforcing particles or whiskers. Al-Si, Al-Ge and Zn-Al fillers, which are commonly used for brazing of aluminum alloys, have good wettability for SiCw /6061Al, SiCP /LD2 and other composites and can be used for brazing of aluminum matrix composites. When using Al-Si and Al-Ge fillers, the main problem is the microstructure changes of the matrix that caused by the diffusion of Si or Ge into the composite matrix. During the holding process of brazing, Si or Ge diffuses into the matrix of the composite, the liquidus temperature of the duffusion zone decreases with the increasing of Si or Ge content. Once the liquidus temperature decreases to the brazing temperature, the diffusion zone will be partially melted. During the subsequent cooling and solidification process, the SiC particles or whiskers are pushed to the weld zone that has not yet solidified, where a SiC-rich layer is formed. The original uniform microstructure transforms into a lamellar structure consisting of a SiC-richw zone and a SiC-poorw zone, so the original microstructure of the composite is destroyed. The SiC-poor zone contains high concentrations of Si and Ge from the eutectics, which degrade the joint properties. Comparatively speaking, the interaction between the Zn-Al eutectic and the composite is weak and the diffusion of Zn into the mtrix metal is much slower. The interactions between the fillers and the composites are related to the processing state of the composites. The diffusion of Si and Ge in the extruded and cross-rolled SiCw /6061Al composites is greate. But in the hot pressed blanks of the same composite without secondary processing, the diffusion of Si and Ge is much weaker and does not cause changes of the composite microstrcture. This may relates to the increased dislocation density in the matrix after extrusion and cross-rolling.

12 11.5 12 11.5 12.8

154

145

149

147

147

GTAW

SiCp -20%/2028Al

100–110

GMAW

19–20

12–14

145–160

GTAW

12–14

Ip = 150 Ib = 50

SiCw -18.4% /6061Al

Arc voltage/V

Welding current/A

Pulsed current GTAW

Welding method

The welding parameters

SiCp -10%/LD2 -Al

Joint

4047

5356

–

–

–

–

–

5.7–7.1

16.5–19

–

LF6 (Al–Mg) 4043

–

argon flow rate/L·min−1

311 (Al-Si)

Welding wire

–

–

–

–

–

Solution and aging

T6

Vacuum dehydrogenation

As welded As welded

Vacuum dehydrogenation

Untreated

As welded As welded

Vacuum dehydrogenation

Untreated

As welded As welded

Untreated Vacuum dehydrogenation

As welded

Heat treatment condition

Vacuum dehydrogenation

Pre-welding treatment

Table 7.15 Examples of welding parameters and joint properties of discontinuously reinforced metal matrix composites

125

175

153

196

218

257

245

105

181

122

165

131

210

Tensile strength/MPa

404 7 Welding of Advanced Composites

7.4 Welding of Discontinuously Reinforced Metal Matrix Composites

405

The dislocations and the grain boundaries provide high-diffusivity path for the Si and Ge atoms. The brazing process parameters must be optimized to correctly match the brazing temperature and holding time when brazing such composites. (3) Friction welding Friction welding achieves joining using the heat generated by friction and the plastic flow deformation caused by the upsetting force, no melting of the base materials occurs during the whole welding process. Therefore it is an ideal method for welding of particle reinforced composites such as SiCp /Al, Al2 O3p /Al. While welding of fiber reinforced composites by this method is not suitable because of the needing large amount of plastic deformation near the weld surface. For particle reinforced metal matrix composites, the small sized reinforcing particles can move with the plastic flow of the base metal during the friction welding process. So the distribution of the particles is hardly changed by the welding process. The particles distribute evenly in the weld, the volume fraction of particles in the weld is similar to that in the base material. Besides, refining of the reinforcing particles at the interface occurs due to the violent collisions of particals during the friction welding. Based on the two factors above, the reinforcing effect will be enhanced. The treatment condition of the base material and the post-weld heat treatment have a great influence on the joint strength. For the T6-treated SiCp /357Al, the strength and hardness of the weld decreased significantly due to the massive dissolution of β'' -Mg2 Si particles during the welding process, but after the post-weld T6 heat treatment, the strength and hardness of the weld returned to the original level. While for the T3 tempered SiCp /357Al composites, the strength and hardness of the weld were increased compared to the base material due to the grains refinement and the increase of dislocation density. The mechanical properties of the two aluminium matrix composites are given in Table 7.16. Table 7.16 Mechanical properties of the two aluminium matrix composites Materials

Treatment conditions

Yield strength σ0.2 /MPa

Tensile strength σb /MPa

Elongation δ/%

SiCp /2618Al (matrix)

Solution and aging

396

455

4.2

SiCp /2618Al (joint)

As welded

–

386

1.8

SiCp /2618Al (joint)

Solution and aging

–

432

1.0

SiCp /357Al (matrix)

Solution and aging

315

352

3.6

SiCp /357Al (joint)

As welded

207

268

3.0

SiCp /357Al (joint)

Solution and aging

313

348

3.1

406

7 Welding of Advanced Composites

Friction welding of Al2 O3p /6061Al to Al alloys such as 6061-T6, 5052-T4 and 2017-T4 has been studied. It was found that sufficient mixing between the composite and Al alloys occurred in the weld. In the weld, the reinforcing particle size and the grain size were reduced compared to the composite. The particles in the composite were pushed into the Al alloy during the welding, and the traveling distance increased in the order of 6061-T6, 5052-T4, and 2017-T4. With a low volume fraction of the reinforcing particles in Al2 O3p /6061Al composite, the hardness of the heat affected zone was significantly reduced compared to the base material, while the hardness of the heat affected zone dropped slightly when the particles volume fraction was higher. (4) Diffusion welding During diffusion welding of aluminum alloys, the interdiffusion of elements is hindered by the the stable and dense oxide film on the Al surfaces. So direct diffusion welding of Al-based composites is difficult and it requires high temperature, pressure and vacuum. Thus the intermediate layers are needed. The addition of intermediate layers enables diffusion welding at lower temperatures and pressures, and also changes the original reinforcing phase-reinforcing phase (P-P) contact into a reinforcing phase-matrix (P-M) contact at the bonding interface, thereby the joint strength is improved, as shown in Fig. 7.15. The reason of the joint strength improvement lies in the fact that P-P bonding is almost impossible, while good bonding of P-M can be formed. According to the chosen intermediate layers, there are two diffusion welding methods, the solid state diffusion welding and the transient liquid phase diffusion welding. 1) Solid state diffusion welding using intermediate layers The key point of this method is the selection of the intermediate layer. The selection principles are that the oxide films can be removed at small deformation, the plastic flow is easy to occur, and adverse interactions between the intermediate layer and the base metal or reinforcing phase can be avoided. Metals and alloys that can be used as intermediate layers include Al-Li alloys, Al-Cu alloys, Al–Mg alloys, Al-Cu-Mg alloys and pure Ag.

Fig. 7.15 Interfacial bonding before and after adding the middle layer

7.4 Welding of Discontinuously Reinforced Metal Matrix Composites

407

Li is highly reactive and can reacts with Al2 O3 to form oxides such as Li2 O, LiAlO2 , LiAl3 O5 , etc., these oxides are easier to break or dissolve than Al2 O3 . So the Al-Li alloy can break the oxide films by chemical mechanism. Therefore, when welding SiCw /2124Al composite using Li-containing intermediate layer, joints with high strength (70.7 MPa) can be obtained at a lower deformation (40%) conditions. It is because the destruction of the oxide films is entirely dependent on the action of plastic flow when using this material as the intermediate layer. Under moderate deformation (20%-30%) conditions, the oxide films are difficult to remove and the shear strength of the welded joint is low. When Ag is used as the intermediate layer, a stable intermetallic compound layer of δ phase will be formed at the interface between the weld and the base material. The formation of δ phase is conducive to breaking the oxide film and promoting the bonding of the interface. However, if the δ-phase content is too large, especially when a continuous δ layer is formed, the joint will be greatly embrittled and the joint strength will be decreased. When the intermediate layer is thin enough (2 to 3 μm), the formation of continuous δ phases in the weld can be avoided and the joint can maintain its high strength. For example, when diffusion welding (470 to 530 °C, 1.5 to 6 MPa, 60 min) was performed with a 3 μm Ag layer coated on the weld surfaces, the shear strength of the joint reached 30 MPa. There are two mechanisms for breaking the oxide films to achieve welding. One is mechanical mechanism and the other is chemical mechanism. If the mechanical mechanism applided alone, such as using superplastic Al-Cu alloys as intermediate layers, large deformations of the bonding interface will occur, making it difficult to use in practical welding. When the chemical mechanism is too strong, brittle phases that harmful to the joint properties may form, e.g., when Ag is used as the intermediate layer, continuous brittle intermetallic compounds will form if the Ag layer thickness exceeds 3 μm, which will reduce the joint strength. Therefore, the most desirable way for removing oxide films is to combine these two mechanisms. 2) Transient liquid phase diffusion welding Because of the large number of dislocations, subboundaries, grain boundaries and phase interfaces in particle reinforced metal matrix composites, the diffusion time can be greatly reduced when the elements of intermediate layer diffuse along these regions. So the transient liquid phase diffusion welding of these composites is easier than that of the base metal. For example, when welding SiCp /Al with Ga as the intermediate layer, the required welding time at 150 °C is less than the aging time. So the welding process can be performed simultaneously with the aging. ➀ Selection of intermediate layer. The principles for selecting intermediate layer materials in transient liquid phase diffusion welding are that the intermediate layer can reacts with the matrix metal to form low melting point eutectics or alloys with a lower melting point than the matrix metal, the diffusion of elements into the matrix metal and the homogenization should be easy and no products that are harmful to the joint performance can form.

408

7 Welding of Advanced Composites

The metals that can be used as intermediate layers in the transient liquid phase diffusion welding of Al-based composites are Ag, Cu, Mg, Ge, Zn and Ga, etc. The alloys that can be used as intermediate layers are Al-Si, Al-Cu, Al–Mg and Al-CuMg alloys. When using Ag, Cu and other metals as the intermediate layer, the matrix metal at the interface melts during the eutectic reaction, the reinforcing phases are pushed by the re-solidification interface and gather at the bonding interface, which would result in reduction of the joint strength. Therefore, the holding time and the thickness of the intermediate layer should be strictly controlled. When the alloy is used as the intermediate layer, the transient liquid phase can be formed as long as the alloys are melted, and it is not necessary to form the transient liquid phase by interdiffusion between the intermediate layer and the base material, the segregation of reinforcing particles can be avoided. When the intermediate layer is too thin, the oxide films on the bonding surfaces cannot be completely removed, the transient liquid phase cannot fully wet the base metal, the P-P contact also cannot be avoided at the bonding interface, so the joint strength will not be high. When the intermediate layer is too thick, the layer is difficult to be consumed completely during the welding process, which is also unfavorable to the joint strength. Sometimes, intermetallic compounds harmful to the joint performance are formed. The strength values and the welding parameters of (Al2 O3 )sf -5%/6063Al composite joints welded using different intermediate layers are shown in Table 7.17. Although high joint strength can be obtained without intermediate layer, the select range of process parameters is very narrow. In contrast, when Cu, 2027Al or Ag is used as the intermediate layer, joints with strength close to that of the base material can be obtained in a wide range of welding parameters. ➁ Welding temperature and holding time. The eutectic temperatures of Al with Ag, Cu, Mg, Ge, Zn and Ga are 566 °C, 547 °C, 438 °C, 424 °C, 382 °C and 147 °C, respectively. When using these metals as intermediate layers, the welding temperatures of transient liquid phase diffusion welding should exceed their eutectic temperatures, otherwise it is actually a solid state diffusion welding with intermediate layers. Similarly, when using Al-Si, Al-Cu, Al–Mg and Al-Cu-Mg alloys as intermediate layers, the welding temperature should exceed the melting point of these alloys. The welding temperature should not be too high, and the temperature should be controlled as low as possible under the condition of ensuring the appearance of the liquid phase required for welding to prevent the unfavourable effects of high temperature on the reinforcing phase. As can be seen from Table 7.17, under the same condition, the strength decreases when the temperature is too high. The holding time is an important parameter affecting the joint performance. If the time is too short, the intermediate layer has no sufficient time to diffuse and a thick residual intermediate layer would remain at the bonding interface, which will limit the increase of the joint tensile strength. The thickness of residual intermediate layer decreases with the increasing of holding time but the joint strength. When the holding time increases to a certain degree, the residual intermediate layer disappears and the joint strength reaches the maximum. If the holding time continues to increase, the joint strength decreases instead of increases, that is because the welding thermal

Bonding temperature/°C

—

16

5

75

30

Materials

None

Ag

Cu

Al-Cu-Mg(A2017)

Al-Cu-Mg(A2017) 2

600

30

30

30

1

1

610

30

1 2

610

600

30

30

30

1

1

550

30

1 2

30

30

–

Holding time/min

610

2

2

Load pressure/MPa

600

600

600

Welding parameters

Thickness/μm

Intermediate layers

Table 7.17 Strength of (Al2 O3 )sf -15%/6063Al composite joints welded using different intermediate layers

187

177

173

184 181

161

119

162

179 181

125

188 145

98 97

Tensile strength/MPa

Bonding interface

Bonding interface

Bonding interface

Base material

Bonding interface

Bonding interface

Bonding interface

Base material Bonding interface

Bonding interface

Bonding interface

–

Fracture position

7.4 Welding of Discontinuously Reinforced Metal Matrix Composites 409

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cycle has a detrimental effect on the composite performance when the holding time is too long. For example, using Ag intermediate layer with a thickness of 0.1 mm to weld of Al2 O3sf -30%/Al composites with a welding temperature of 580 °C and a pressure of 0.5 MPa. When the holding time was 20 min, a thick residual intermediate layer remained in the joint and the average tensile strength of the joint was 56 MPa. When the holding time was 100 min, the joint tensile strength reached the highest value of about 95 MPa. But the joint tensile strength decreased to about 72 MPa when the holding time was 240 min, ➂ Welding pressure. The welding pressure has a great influence on the joint performance during transient liquid phase diffusion welding. If the pressure is too small, the plastic deformation of the bonding interface is small. Close contact between the bonding interface and the intermediate layer cannot be achieved, lack of fusion will appear, which would reduce the joint strength. If the pressure is too high, the liquid metal can be extruded from the bonding interface, resulting in the segregation of the reinforcing phases. And the liquid phase cannot fully wet the reinforcing phases, which will also causes lack of fusion. For example, during the welding of Al2 O3sf -30%/Al at a welding temperature of 580 °C using a 0.1 mm thick Ag as the intermediate layer, obvious microscopic holes appeared at the bonding interface when the pressure is less than 0.5 MPa or greater than 1 MPa, all the joint strength values are low. The joint tensile strength is less than 60 MPa in a 1 MPa, 120 min welding condition, but the strength reached about 90 MPa in a 0.5 MPa, 120 min welding condition. ➃ Treatment of the surfaces to be welded. The treatment of the weld surfaces has a great influence on the joint performance. Contrast analysis of the effect of three kinds of treatments, such as electrolytic polishing, mechanical cutting and brushing with a wire brush, on the performance of AlO23sf /Al joints indicated that the highest joint strength was obtained when treated with electrolytic polishing, the lowest strength was obtained when treated with brushing. This is due to the fact that when treated with the latter method, a number of tiny Al2 O3 wear debris accumulates on the weld surfaces, which hinders the close contact of the weld surfaces and causes the reduction of the joint strength. During electrolytic polishing, no Al2 O3 debris is present on the surface to be weldeded, but the fibers will be exposed out of the matrix surface. The electrolytic polishing time has a significant effect on the joint strength. If the electrolytic polishing time is too long, the exposed fibers become longer and will be broken under the action of load pressure. The exposed fibers impedes the contact between matrix metals and causes degradation of the joint performance. (5) High energy beam welding High energy beam welding using electron beam or laser beam is characterized by high heating and cooling rates, small sized molten pool and short existence time of the molten pool. These characters are beneficial to welding of metal matrix composites, but it is difficult to avoid the reactions between SiC and Al matrix when welding of SiCp /Al or SiCw /Al composites due to the high temperature of the molten pool.

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In particular, when using laser beam welding, the laser beam preferentially heats the reinforcing phases with high resistivity and causes severe overheating of the reinforcing phases, thus, the reinforcing phases were rapidly dissolved and violently reacts with the matrix. So laser beam welding is not suitable for welding of SiC/Al composites. During the laser beam welding of Al2 O3 /Al composites, although there are no reactions between the reinforcing phases and the matrix, the welding process stability is still poor due to the formation of sticky slag by the melting of Al2 O3 . The heating mechanisms of electron beam welding and laser beam welding are different. The electron beam heats the matrix and the reinforcing phases uniformly, so the interfacial reactions can be controlled at a low degree by adjusting welding parameters. Due to the impact of the electron beam and the rapid cooling of the molten pool, the reinforcing particles distribute uniformly in the weld. The electron beam weldability of SiC particle-reinforced Al-Si matrix composites is good, and the interfacial reactions are suppressed easily due to the high Si content in the matrix. Better results can also be obtained in electron beam welding of Al2 O3 particlereinforced Al–Mg-based or Al–Mg-Si-based composites. (6) Other welding methods Capacitor discharge welding is suitable for welding of metal matrix composites. Although the bonding interfaces are also melted during the welding, the interfacial reactions can be completely avoided because of the short discharge time (0.4 s), the fast cooling rate of the nugget (106 °C/s), and the fact that all of the molten metal is extruded. Moreover, defects such as porosity, cracks and fibre breaks do not occur in the weld, so the joint strength is high. The disadvantages of this method are that the weld area is small and the applications of the welding method is limited. Resistance spot welding has short heating time, small weld nugget, and good controllability. The interfacial reactions can be avoided effectively during resistance spot welding of composites. By using lap joints, the welding between fiber reinforced metal matrix composites can be changed into welding between Al and Al, so this method is suitable for welding of composites. However, severe segregation of the reinforcing phases in the weld nugget is likely to happen during resistance spot welding of discontinuously reinforced metal matrix composites. The reinforcing phases segregation can be mitigated by reducing the weld nugget size.

References 1. Wei Y (1987) Composite materials. China Machine Press, Beijing 2. Kenneth . G. Kreider. Metal matrix composites. Wen Zhongyuan, et al. Beijing: National Defense Industry Press, 1982. 3. Ibrahim IA et al (1991) Particle reinforced metal matrix composite - A review. J Mater Sci 26:1137–1156 4. Hirose A, Fukumoto S, Kobayashi KF (1995) Joining process for structure application of continuous fibre reinforced MMC. Key Engineering Material 104–107:853–872 5. Hall IW et al (1992) Microstructure analysis of isothermally exposed Ti/SiC MMC. J Mater Sci 27:3835–3842

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6. Wo Dingzhu. Encyclopedia of composites. Beijing: Chemical Industry Press, 2000 7. E. K. Hoffman, et al. Effect of braze processing on SCS-6/β21S Ti matrix composite, Welding Journal, 1994 (73), 8: 185s-191s. 8. Blue CA et al (1996) Infrared transient-liquid-phase joining of SCS-6/β21S Ti matrix composite. Metallurgical and Material Transactions 27A:4011–4018 9. Chen Maoai, Wu Renjie, Lu Hao, et al. Study on the weldability of metal matrix composites. Materials Development and Applications. 1997, 12(3): 34-40 10. Chen Maoai, Lu Hao, et al. Arc welding of SiC particle reinforced LD2 Al alloy matrix composites. Acta Metallurgica Sinica, 2000, 36(7): 770-774 11. Ren Jialie., Wu Aiping. Joining of advanced materials, Beijing: China Machine Press, 2000. 12. O. T. Midling, et al. A process model for friction welding of Al-Mg-Si alloys and Al-SiC MMC-I. HAZ temperature and strain rate distribution. Acta Metallurgica et Materialia, 1994, (42)5: 1595–1609. 13. Suzumura Akio, et al. Diffusion brazing of Al2 O3sf /Al MMC, Material Transaction, JIM, 1976, (37)5: 1109–1115. 14. Yu Qizhan, Shi Chunyuan. Welding of composite materials, Beijing: China Machine Press, 2012. 15. Maoai C, Junhua C, Jinqiang G (2005) Welding of composite materials. Chemical Industry Press, Beijing

Chapter 8

Connection of Functional Materials

Functional materials are special materials with physical also the mechanical properties, such as superconducting materials and shape memory alloys are typical of functional materials. Functional materials play a pivotal role in the development of high technology, which have been highly valued by countries around the world. It is difficult to achieve the connection of the superconducting materials or the shape memory alloys using the conventional welding methods, because it is very difficult to obtain the superconducting properties or shape memory effects equivalent to those of the parent material in the welded joints. In this chapter, the superconducting materials and shape memory alloys are used to illustrate the welding of these two typical functional materials.

8.1 Connection of Superconducting Materials to Metals The discovery of the superconducting materials in the 1980s turned a new page during the development of superconductivity technology. Practical superconducting magnets made of the Nb–Ti superconducting materials have entered the stage of large-scale, which not only puts more stringent requirements on the performance of superconducting materials, but also requires the longer the length of the conductor. Perversely, some superconducting devices with the practical superconducting materials weighing tens of tons and conductors of at least several thousand meters in length. The manufacturing of such lengths of superconducting wire is limited by the processing equipment, so the welding of superconducting materials has been received lots of attentions.

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8.1.1 Performance Characteristics and Applications of Superconducting Materials (1) Basic features Superconducting materials are materials whose resistance suddenly drops to 0 at very low temperatures, which should be in a superconducting state. Metals in general still have resistance at very low temperatures, and only superconducting materials reach a certain critical temperature (T c ) where the resistance drops abruptly to 0 and are characterized by full electrical conductivity. A superconductor has a permissible current density, and its full conductivity is destroyed when the current density exceeds a certain critical current density (J c ). Superconducting materials are also fully antimagnetic in that an applied magnetic field can not enter the superconductor when the material is in the superconducting state. The normal state of the material, which was originally in the magnetic field, which would completely exclude the magnetic field that was inside the conductor when the temperature drops below the critical temperature (T c ), and transforms into the superconducting state (this complete antimagnetism is called the Meissner effect). When the external magnetic field reaches a certain critical magnetic field strength (H c ), the magnetic field immediately enters the superconductor, causing the material to be in the superconducting state to return to its normal state, also superconductivity should be destroyed. The transition of a superconducting material from the normal to the superconducting state is limited by the three conditions: the critical temperature (T c ), the critical current density (J c ), and the critical magnetic field strength (H c ). A superconducting material could exhibit its superconducting properties only when it is below each critical point. Different superconducting materials have different critical values T c , J c , H c . How to obtain superconducting materials with the high critical temperature, the high critical current density and the high critical magnetic field strength, so they can be used in industry, which has been a matter of concern and a goal to be pursued. (2) Types of superconducting materials The types of superconducting materials are the pure metals (e.g., mercury, lead, indium, tungsten, etc., at superconducting critical temperatures T c close to absolute zero), the alloys, the compounds, oxide ceramics, and a small number of organic superconducting materials. The following three main types of the superconducting materials has been researched: the alloy superconductors, the intermetallic compounds superconductors (e.g. Nb3 Sn) and the oxide ceramics superconductors (e.g. Y–Ba–Cu–O, Bi–Pb–Sr–Ca–Cu–O). ➀ Alloy superconductors (e.g. Nb–Ti, Nb–Ti–Ta, Nb–Zr, etc.) are the most widely used representative superconducting wires, such as the Nb–Ti superconducting wire, a superconducting material operating at liquid helium temperature (4.2 K).

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➁ The intermetallic compound superconducting materials have the higher critical magnetic field strength (H c ) and higher critical transition temperature (T c ) than the alloy superconducting materials, which can be used as the superconducting wires for generating the high magnetic fields. However, intermetallic compounds are more brittle and special measures, which need to be considered for their design and fabrication. An example is the high superconductivity wire Nb3 Sn, which operates at liquid helium temperatures. ➂ In terms of superconductivity, oxide ceramic superconductors are the best, but the outstanding problem that hinders the development of oxide ceramic superconducting materials is the brittleness and the resulting difficulty in forming and processing, including the difficulty in welding. The first two types of the superconducting materials are the main ones that are available on a practical and the industrial scale of manufacture. Among them, the alloy superconducting materials have the best mechanical properties and better processing properties, also high current densities, which can be obtained at lower magnetic induction strengths (below 10 T). (3) Application prospects Due to their unique, fully conductive properties, superconducting materials allow the current to increase to a critical current density (J c ) to pass through in an unimpeded state, i.e., to transmit current without energy loss, providing that the conditions of the critical magnetic field strength (H c ) and the critical temperature (T c ). In addition, superconducting materials can be used in the superconducting magnetic levitation systems due to their unique and complete antimagnetic properties. Lots of applications of the superconducting materials are related to the power saving also the energy conservation. For example, the superconducting materials can be applied to the alternating (direct) current power transmission, the large electromagnets, the superconducting accelerators, the electromagnetic propulsion, the magnetic levitation trains, etc. Superconducting materials can be also widely used in the instrumentation and apparatus, such as the nuclear magnetic resonance imaging devices for medical devices, the superconducting quantum interference devices for the electromagnetic measurements, such as the geophysical measurements and the biomagnetism, etc.

8.1.2 Connection Methods for Superconducting Materials There are many joining methods for superconducting materials, and those used for large superconducting magnets include the explosion welding, the diffusion welding, the brazing, the cold-press welding, and the microwave welding. Table 8.1 shows the order of magnitude range of low temperature resistivity of superconducting joints obtained by different joining methods.

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Table 8.1 Comparison of low temperature resistivity of superconducting joints Welding method Energy storage impact welding Low temperature joint resistivity/Ω cm

10–13

Explosive welding

Cold pressure welding

Microwave welding

Diffusion welding

Brazing

10–9 ~ 10–10

10–8

10–9

10–8 ~ 10–9

10–8 ~ 10–9

The choice of superconducting material welding (connection) method, in addition to meet the above requirements for the superconducting material joint resistivity, more importantly, must also consider the possibility of the engineering applications and reliability. Several superconducting materials welding methods, both domestic and foreign have conducted experimental research, some of the methods have been applied in engineering. However, for the specific superconducting materials, the possibility of engineering applications of different welding methods and the reliability of the need for in-depth and systematic research will be discussed. (1) Explosion welding Explosion welding is the use of the explosion of chemical explosives as an energy source instantaneously and sharply releasing amounts of the energy, so that the surface of the welded metal produces a metal jet and mutual contact between pure metals to achieve solid metal bond. A NbTi–Cu multicore superconducting composite joint with cross-sectional dimensions of 6 × 6 mm was fabricated by the explosion welding at Lawrence Livermore Laboratory, USA, for the fabrication of a pair of long superconductors for a pair of cathode-type large coils for a controlled fusion reactor. The sample of this joint had a critical current of 750 A and a low-temperature resistance of 3 × 10–11 Ω at a temperature of 4.2 K and a field strength of 6 T. Imperial Metal Industries of the United Kingdom wound magnet coils with NbTi–Cu superconductors containing blast-welded joints, which had a cross-sectional size of 10 × 1.8 mm and a critical current of 1500 A and a low-temperature resistance of 3 × 10–9 Ω at a field strength of 6 T. The Northwest Institute of Nonferrous Metals in China also successfully welded samples with the small sizes of the NbTi–Cu multicore superconductors with specimen sizes of 7 × 3.6 mm, 204 cores and 3.6 × 1.8 mm, 174 cores using the explosion welding technique. The average current value of the superconductor joints was 180 A at a temperature of 4.2 K and a field strength of 5 T. This value was 10% lower than the critical current of this conductor under the same conditions, and the tensile strength at the room temperature and liquid nitrogen conditions was almost equal to that of the parent material. From the above experimental results of the Northwest Institute of Nonferrous Metals and the Lawrence Livermore Laboratory in the United States, both have achieved the more satisfactory results, and some of the research results have been applied to the productions of the medium-length superconducting long strips.

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However, from the experimental methods and results of both, still several issues should be further studied. ➀ The formation of the blast welded joint sections and the problem of the formation mechanism in blast welding, for the oblique joints, the high speed oblique impact generates lots of pressure and the superconductor at that place is subjected to a large shear effect. The plastic shear work is converted into the heat, and only a small portion of the heat dissipated by heat conduction, most of which drives the temperature rise at the joint. The shear strength of the material decreases with temperature rise, thus producing melting in a very narrow region at an interface contact. The instantaneous melting also cooling of the joint interface during the blast welding process produces a very thin new interfacial layer microstructure completely different from that of the superconducting alloy. This new microstructure leads to be changed in the superconductivity and mechanical properties of the joint. The relationship between the microstructures of the interface layers, the formation mechanism and the explosion welding parameters, also the influence of the formation of the interface layer on the superconductivity are still to be investigated. ➁ The problem of applicability of the explosion welding technology. The explosion welding technology has shown superiority in larger cross-sectional area of superconducting strips. The results of the Lawrence Livermore Laboratory in the United States and the Northwest Institute of Nonferrous Metals in China are good evidence of this. Large cross-section superconductor blast welding head configuration is easy to observe alignment, and conductor damage after blast welding is minimal. However, for fine or very small cross-section superconductors, the use of the explosion welding is extremely difficult, especially for the multi-core superconductors, because the configuration of the joint when the core wire at both ends of the conductor is difficult to align, a slight deviation from the welding effect has a great impact, which is also the shortcomings of the explosion to be welded the connection superconducting thin wire. (2) Diffusion connection The diffusion welding is a metallurgical bond achieved by the diffusion of atoms into each other. A beveled lap joint is also used for diffusion welding of superconducting materials, where the ends of the superconductors to be welded are beveled and angled so that the bevels are clean and free of oxides. The two conductors are lined up in a line at the top of the bevel and placed in a special press mold. The superconductor is heated under pressure and held and cooled slowly to achieve the required temperature for the purpose of the joining. Tests of diffusion welding of the NbTi–Cu multicore superconductors at the University of California, USA, have shown that the technique can be applied to the fabrication of conductors for large coils for controlled nuclear fusion. Two superconductor sizes were used in the tests: one with a diameter of 5.4 mm and a core diameter of 600 μm, and the others with a diameter of 1.5 mm and a core diameter of 200 μm. The optimum welding conditions for the tests were: a conductor bevel

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angle of 15°, a welding temperature of 450 °C, a welding pressure of 600 MPa, and a holding time of 30 min. But the electrical properties parameters (critical current and low-temperature resistance) could not be given. Diffusion welding tests were conducted at Northern Jiaotong University using diffusion welding methods on flat wire (Length 30 mm) of 2.1 × 1.54 mm and short specimens of round wire with a diameter of 0.75 mm and a length of 10 mm, and the diffusion welding temperature was 360–380 °C. For the flat wire, the critical current dropout rate was greater than 20% and the tensile strength dropout rate of the joint at room temperature and liquid nitrogen temperature was greater than 15%; for the fine diameter round wire, the critical current dropout rate was less than 12% and the tensile strength dropout rate at room temperature and liquid nitrogen temperature was less than 10%. Theoretically, the applications of the diffusion welding method to the connection of the Nb–Ti superconductors is feasible, but still some issues should be further explored from an engineering practical points of the views as follows: ➀ Applicability of the diffusion joining From the results of available studies, the application of the diffusion welding process to the joining of superconductors is still limited to the specimens with small sizes or experimental stages. For the production of thousands or even tens of thousands of meters long strips, the diffusion welding method is constrained by the equipment and is very difficult to apply in the engineering practice. ➁ Selection of diffusion welding process parameters a. Diffusion welding temperature. According to the equation D = D0 exp(−t/RT ) (D is the diffusion rate of the element, μm2 /s; D0 is the temperature-dependent coefficient; t is the diffusion activation energy, J/mol; T is the heating temperature, K; R is the gas constant), the diffusion process accelerates with increasing temperature, and the diffusion temperature should conform to the relation T D = 0.7T M . However, the diffusion welding temperature of Nb–Ti superconductor is governed by its own heat treatment regime. Nb–46.5Ti (mass fraction) alloy aged at 420 °C, the critical current density (J c ) of the sample at 5 and 8 T reaches 3700 and 1560 A/mm2 respectively. Diffusion welding above its aging temperature will destroy the superconducting tissue, resulting in the loss of superconductivity. Too high temperatures tend to generate Cu–Ti compounds, increasing the resistivity of the joint interface, eventually resulting in the loss of superconductivity. b. Diffusion welding pressure. The diffusion welding process applies a large pressure to bring the bonding surfaces into close contact; such a large pressure generates a large strain in the superconducting joint, which may destroy the microscopic substructure formed by plastic deformation and extension along the tensile direction, leading to the destruction of the superconductivity. (3) Brazing connection Brazing is a commonly used connection method in the field of superconductivity. Many specimens in the electrical property measurements of superconducting specimens are connected by brazing. The brazing method was used at an Oak Ridge

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National Laboratory, USA, to braze several kilometers of large cross-sectional NbTi– Cu multicore superconducting composite for the plasma magnetic column device of a controlled fusion reactor, and the wettability and fluidity of the brazing material were researched extensively, also the more systematically and joint tensile performance tests had been done. The brazing materials consisted primarily of Pb–1.5Ag–1Sn and 95Sn–5Ag (wt%); the brazing flux consisted of ZnCl2 + NH2 Cl + HCl + H2 O. Resistance brazing studies on copper-based Nb–Ti composite superconducting wires with a cross-sectional size of 2.8 × 1.2 mm and 178 cores were also carried out at Northern Jiaotong University. High containing of the lead, tin and indium-containing brazing materials were used, and the brazing thickness was 0.2 mm. The joint performance test results: the critical current degradation rate was 12% at 3 T and 4% at 4 T; the degradation rate of the brazed joint strength at liquid nitrogen temperature was less than 10%. The resistance brazing method has been successfully applied to the construction of the first stable strong magnetic field device in China (Hefei 20 T hybrid magnet system), where the five conductors in the superconducting coil winding have to be connected one by one. The joint is in the form of two wedgeshaped surfaces welded together, with a foil strip of solder on the bonding surface, and is press-welded by a resistance welding machine. The superconductivity of the joint and the bending tension of the joint should meet the design requirements. The brazing temperature is governed by the heat treatment temperature of the superconducting material ( Af ). When the applied load exceeds the critical stress σ M while the martensitic phase transformation would occurre, also the martensites can be produced whil the deforming process. When the stress is removed, the strain disappears with the inverse transformation of martensite, and the original state of the parent phase is restored. The phenomenon of the superelastic shape recovery is essentially the same as the shape memory effect, which is caused by the inverse transformation of martensite, but this martensite is very unstable, also the corresponding deformation can be recovered once the load is removed. (3) Application of the shape memory alloys The TiNi shape memory alloy has excellent shape memory effect and super elasticity, high specific strength, corrosion resistance, wear resistance and good biocompatibility, etc. It has the broad application prospects in the fields of aerospace, marine development, instrumentation and medical devices.

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1. Applications in the aerospace Shape memory alloys have been used in aerospace and space devices. The United States National Aeronautics and Space Administration in the “Apollo” moon campaign with NiTi memory alloy manufacturing hemispherical unfolded antenna, its own size is quite large, in order to facilitate the rocket or space shuttle transport, scientists will first “compression” of this antenna, and then use sunlight to heat it back to its original shape. After delivery to the lunar surface, then use sunlight to heat it back to its original shape. For example, the hemispherical antenna is first made at normal temperature according to the predetermined requirements, then the temperature is lowered and it is compressed into a ball and put into the cryogenic container of the lunar module, which is taken out after delivery to the Moon, and when the temperature rises to about 40 °C under the sunlight, the antenna “remembers” its original shape and automatically expands into a hemisphere. Dutch scientists use the NiTi memory alloy to manufacture the artificial satellite antenna, also through the “compression” technology to put it rolled in the satellite body, when the satellite into orbit, and then use the sunlight heating, so that it restores the “memory” and in space automatically. The satellite is then heated by sunlight to restore its “memory” and automatically unrolls in space. In space, Russian-made shape-memory alloy devices have reached a practical level, such as the attachment and assembly of large antennas for the space programme and antenna masts for the MIR space station. In the United States, the space program’s application of shape-memory alloy drive pins to release post-launch payloads, which also has proven to be successful. The brittle pins are used in pre-pressurized cylinders that cause notched pins to fracture when the shape is restored, and it is much safer than the conventional explosive release devices. In addition, a shape memory release device with an openable container is used in satellites to protect the sensitive germanium detectors from contamination during assembly and launch. The shape memory alloys were used in the United States in 1970 to make cryogenic mating connectors on the F-14 fighter jet, and were subsequently used in millions of the connections. 2. Application in industrial automation The shape-memory alloys are used extensively in the automatic control technology. For example, the “thermostatic valve” used in residential heating systems, which can work with the help of the shape memory alloys. When the room temperature rises to a certain value, the memory alloy spring elongates and then closes the valve; when the temperature drops to a certain value, the memory alloy spring shortens and the valve is opened again, thus maintaining a constant room temperature. By adjusting the knob to change the pressure of the spring, the room temperature can increase or decrease. Actuators made with shape memory alloys can operate at low voltages and low currents, which is both safe and power-efficient, and some countries have used these small and delicate components in micro-robots. Because of their simple structure and flexible control, shape memory alloys have an unique technical advantage in

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light robots and miniaturized systems. A few examples of applications are shown as follows: ➀ One-way shape recovery used one-way shape memory effect, such as the pipe joints, antennas, collars, etc. For example, the Ti–Ni shape memory alloy pipe joints can be used for sealing connection of various types of liquid and gas highpressure or low-pressure fittings, and also for sealing connection and fastening of heterogeneous devices with stable performance. ➁ The exogenous two-way memory recovery, that is, the use of one-way shape memory effect and the use of external forces to do the repeated actions with the rise and fall of temperature, such as thermal components, robots, terminal blocks, etc. Memory alloy spring. For example, put this spring in hot water, its length elongated, and then put in cold water it will return to its original state. The shape memory alloy spring can be used to control the water temperature in the bathroom plumbing, and the “memory” function can be used to adjust or shut off the water supply line when the hot water temperature is too high. It can also be used as a fire alarm device and as a security device for electrical equipment. When a fire occurs, the spring made of memory alloy deforms and activates the fire alarm device to achieve the purpose of alarm. ➂ Endogenous two-way memory recovery, that is, the use of two-way memory effect with the temperature rise and fall to do repeated actions, such as thermal machines, thermal components, etc. However, such applications are not commonly used because of the fast memory decay also the poor reliability. ➃ Super-elastic applications such as springs, lug posts, eyeglass frames, etc. For example, eyeglass frames made of memory alloy can be restored to their original shape by heating them in hot water if they are touched and bent. The shape memory alloys can be used as the low temperature mating connectors in hydraulic systems in aircraft and in petrochemical and power systems. The availability of wide thermal hysteresis NiTiNb alloys has made shape memory alloy connectors and joining devices even more attractive. Another type of connection is in the form of a welded mesh wire, which can be used to make safety joints with braided layers of wire for conductors. Such connectors have been used in sealing devices, electrical connection devices, electrical engineering and mechanical devices and can work reliably at −65 to 300 °C. Sealing devices have been developed as the electrical component connections in harsh environments. The interconnecting cables of computer connection boards require a connector that closes when the contact resistance is minimized, preventing the damage to the electrical components. 3. Applications in medicine Memory alloys used in the medical field should meet the requirements for reliability in chemistry and biology (biocompatibility) in addition to having shape memory or superelastic properties. In practical terms, only alloys that form a highly stable passivation film upon contact with a living organism can be implanted in living

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organisms, of which only TiNi alloys meet the conditions for use and are currently the main memory alloys used in medicine. TiNi alloy is very biocompatible, in medical TiNi alloy is widely used in oral dental orthopedic wire, various orthopedic rods used in surgery, bone connectors, vascular clips, blood clotting filters, etc. ➀ Orthodontic wires Orthodontic wires made of superelastic TiNi alloy wires and stainless steel wires. Usually stainless steel wire and CoCr alloy wire are used for orthodontics, but these materials have the disadvantage of high modulus of elasticity and small elastic strain. With TiNi memory alloy orthodontic wires, no plastic deformation occurres even if the strain is as high as 10%, and the stress-induced martensitic phase transformation causes the elastic modulus to exhibit nonlinear characteristics, i.e., the orthodontic force fluctuates very little when the strain increases, which reduces the patient discomfort. ➁ Scoliosis orthosis Orthopedic rods made of shape memory alloy are placed and fixed only once. If there is a change in the corrective force of the rod, sufficient corrective force can be restored by heating the shape memory alloy outside the body to a temperature approximately 5 °C higher than the body temperature. 4. Application in flange sealing connections The failure of a flange seal connection, a common form of detachable connection in industrial installations such as pressure vessels, power machinery and connecting piping, has potentially catastrophic consequences. Bolts also exhibit creep relaxation under prolonged tension. To meet the sealing requirements of flange connections in specific operating conditions, such as nuclear power plants and aerospace facilities, it is necessary to ensure that the sufficient compression force is maintained on the sealing element during the long operating cycles. Pipe joints made of the shape memory alloys have been used in engineering, especially in the connection of hydraulic lines for aerospace applications. ➀ Performance of flange sealing alloy The creep relaxation of the flange sealing connection is the result of the interaction between the gasket, bolt and flange. When the bolted flange connection enters working condition, under the action of medium pressure, the bolt deformation elongates, the gasket deformation thins, and the compression force acting as a seal decreases. With the passage of time and temperature, the creep of each component gradually increases so that the pressure on the gasket is smaller and smaller, resulting in sealing failure. Especially in high temperature flange connection, the creep relaxation phenomenon is more obvious. Shape memory alloys have a shape memory effect, and alloys with thermoelastic martensitic phase transformation also exhibit superelasticity. The memory alloy is a stable parent phase above the end of the austenite transformation temperature (Af ) and becomes a martensitic phase below the end of martensite transformation temperature (M f ), with both phases coexisting between M f and Af . When a sample of a certain shape of the parent phase is cooled from above Af to below M f to form

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martensite, it will be deformed below Mf and heated to above Af with inverse phase transformation, and the material will automatically recover its shape in the parent phase. When memory alloys are constrained during the shape recovery process, they generate significant stresses to resist. For example, the TiNi alloys can generate a resistance of 700 MPa when the memory effect is prevented. This resisting stress can be applied directly or indirectly in bolted flange joint seals to compensate for the reduction in compression force caused by the creep relaxation. ➁ Shape memory alloy sealing element In bolted flange connections, progress has been made with bolts (or combination bolts) and gaskets made of shape memory alloy. When the bolted flange connection enters working condition after pre-tightening, with an enhancement of the working temperature and internal pressure, the bolts and gaskets show shape memory effect, producing inverse deformation, preventing the elongation of the bolts and maintaining the compression force loaded on the gaskets in a more constant range; the shape memory effect and super elastic properties in the subsequent long-term work, the creep relaxation, internal pressure and temperature field fluctuations caused by the reduction of compression force to play an active The shape memory effect and superelasticity actively compensate for decreases in compression force caused by the creep relaxation, internal pressure and temperature fluctuations during the subsequent long-term operation, resulting in the excellent sealing results. Figure 8.14 shows a new combination bolt made of two kinds of the materials, coaxially combined and made into one piece, which uses the memory effect to suppress the bolt stress relaxation behavior. Low-period repeated load tests on double-headed bolts made of shape memory alloy have shown that the shape memory bolts have a higher energy dissipation capacity relative to steel bolts at the same stress state (below yield strength for steel bolts and above martensitic transformation strength and below yield strength for memory alloy bolts). This has applications for improving the fatigue resistance and extending the seal life of bolted flange connectors in special applications. The memory alloy gasket shown in Fig. 8.15a consists of corrugated memory material and a protective film, and the gasket shown in Fig. 8.15b is made of a Vshaped band made of memory alloy wound spirally with the filling material, both of which use the shape memory effect to compensate for the drop in sealing compression force. The tests show that the sealing performance of the TiNi alloy flat gasket is better than that of the aluminum flat gasket, and the memory alloy gasket in the mother phase can still maintain the sealing effect through its superelasticity when there is a 20% fluctuation in the axial compression force. The advantages of shape-memory alloy shims are obvious, but they require a lowtemperature martensitics for preload and a high-temperature parent phase for the memory effect to produce the required resistance; the memory effect superelasticity also requires the shim to be operated at a temperature between the Af point and the M d (maximum stress-induced martensitic temperature) point, and suitable memory alloy shims should be developed for the different operating environments. ➂ Application of the several memory alloys in the sealing joints

8.2 Shape Memory Alloy to Metal Connection Fig. 8.14 Shape memory alloy combination bolt. 1–Shape memory alloy; 2–common alloy steel bolts; p–working internal pressure; T–working temperature

Fig. 8.15 Shape Memory Alloy Spacer. 1–Protective film; 2, 3–shape memory alloy; 4–filling material; p–working internal pressure; T–working temperature

441

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8 Connection of Functional Materials

Memory alloys can be classified as high temperature shape memory alloys (HTSMA) and low temperature shape memory alloys (LTSMA) according to the onset of martensitic inverse phase transformation temperature (As ). Alloys with subsequent thermal–mechanical treatment, under unconstrained stress conditions, As > 120 °C are classified as HTSMA and the vice versa as LTSMA. HTSMA and LTSMA are suitable for different applications in terms of extended seal life, reducing the maintenance and improved seal reliability. HTSMA is more valuable for the development of applications to solve the problems of nuclear reactors, turbine hot zones, geothermal Sealing of piping connections. The development of sealing elements suitable for bolted flange connection requirements requires consideration of the machinability, phase stability, mechanical stability and economics of the alloy material in addition to the applicable temperature of the material and seal design. a. Ti–Ni alloy. Ti–Ni alloy is one of the earliest memory alloys studied, and the addition of Series 3 elements to form Ti–Ni–X alloy can change the properties to meet the requirements of different applications. For the development of flange joint sealing components, Ti–Ni, Ti–Ni–Nb and Ti–Ni–Cu are of application value in LTSMA; Ti–Ni–Hf and Ti–Ni–Pd are of application value in HTSMA. b. Cu-base alloys. Cu base memory alloys have some characteristics inferior to the Ti–Ni alloys, but are favored in engineering applications due to easy processing and low cost. c. Cu-based memory alloys include two major alloy families, Cu–Zn and Cu–Al. The M s point of the Cu–Zn family of alloys is generally below 100 °C and has poor thermal stability, while the Cu–Al family of alloys is expected to be developed into HTSMA. c. Fe-base alloys. Fe-base alloys have the advantages of the high strength, good plasticity, easy forming and processing and low price, although the memory effect is not as good as Ti–Ni alloy, but there is the great potential for application. Among them, the Fe–Mn–Si alloy can be used to develop the bolts or bolt components, its alloy reverse phase transformation occurs at 100–200 °C, additions of Cr, Ni, Co can prevent rust, improving the corrosion resistance, the additions of the rare earth elements Re can improve the memory effect and the M s point. d. Ni–Al alloy. Ni–Al alloy martensite phase transition temperature (M s ) varies with the Ni content from −196 °C to about 950 °C, due to the alloy contains a large number of Al, presenting good high temperature oxidation resistance and thermal conductivity, suitable for the development of high temperature shape memory alloy, is currently considered as one of the highest development potential of high temperature shape memory alloy. In the bolted flange connection seal, the Ni–Al–Fe and Ni–Al–Mn alloys can be further developed.

8.2 Shape Memory Alloy to Metal Connection

443

8.2.2 Advances in the Welding of Shape Memory Alloys Current practical TiNi shape-memory alloys are mainly manufactured into the simple industrial fabrications (e.g. springs, wires and sheets), and welding TiNi shapememory alloys into more complex shapes is an important way to expand their applications. Research and development on the welding of TiNi shape memory alloys has focused on the welding methods also processes, and the effects on the tissue properties of the joints. Most of the current research is exploratory, but it is of the practical importance to advance its application. Numerous researchers have done lots of work in the development of the TiNi shape memory alloy connections, including the argon arc welding, the electron beam welding, the laser welding, the resistance welding, the friction welding, also brazing, etc. When the TiNi shape memory alloy are welded, a certain mechanical properties should be ensured, also the shape memory function should meet the required requirements. Therefore, it is more difficult to weld than general structural materials, and the welding process is subject to more restrictions, which poses great difficulties to its welding. (1) Fusion welding of TiNi shape memory alloy Because TiNi shape memory alloys’microstructures and mechanical properties are extremely sensitive to temperature changes, the affinity of Ti for N, O and H is particularly strong at high temperatures, and TiNi memory alloy can easily absorb these gases during the fusion welding process and form brittle compounds at the joint. During fusion welding, the joint forms a coarse as-cast structure and forms compounds such as Ti2 Ni and TiNi3 during solidification, which adversely affects the mechanical properties and shape memory effect of the joint. Therefore, it is necessary to prevent the intrusion of N, O, H, etc. and not to produce liquid phase as much as possible when joining these alloys. Due to the characteristics of the TiNi shape memory alloy, the solid phase joining methods, such as brazing, friction welding and resistance welding should be conducive to the joining of the TiNi shape memory alloys. Welding production in melt welding is the most widely used. The tungsten argon arc welding had been used to connect TiNi shape memory alloys in 1960s, but did not obtain the satisfactory results. The welding results of welding TiNi shape memory alloy using melt welding methods, such as the argon arc welding, the electron beam welding and the laser welding still should be improved. The main problems in the fusion welding of the shape memory alloys are shown as follows. (i) Brittle welded joints due to the dissolutions of N, O, H, etc. (ii) The cast crystalline microstructures produced in the weld zone, which should retard the martensitic phase transformation in a certain extent, also affect its shape memory effect. (iii) The shape memory effect of the weld heat affected zone is affected by the grains’growth that destroys the ordered matrix structure of the base materials.

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8 Connection of Functional Materials

➃ Easy formation of intermetallic compounds (e.g. Ti2 Ni, TiNi3 ), which adversely affects joint strength and shape memory effect. 1. Tungsten Argon Arc Welding By studying the effect law of N and O on the microstructure, shape memory effect and mechanical properties of the TiNi shape memory alloy tungsten arc welded joints, the results show that N and O have adverse effects on the microstructures and properties of TiNi shape memory alloy tungsten arc welded joints, with the increase of the N and O content in the joints, the second phase particles (such as TiN, Ti4 Ni2 O, etc.) would appear in the joint area, the phase transformation temperature decreases, the shape memory effect and reduced tensile strength of the joint. When He gas-shielded tungsten arc welding was used to join the TiNi memory alloys, the welds showed a fine dendritic microstructure, but the shape memory effect and mechanical properties of the joints were still poor. 2. Electron beam welding Mechanical properties tests with electron beam welding for welded joints of TiNi shape memory alloy sheets with a thickness of 1.16 mm showed that the fracture stress in the parent phase state was 740 MPa and elongation was 26% at the room temperature after the memory alloy was calendered and heat treated at 973 K × 60 min. The mechanical properties of the TiNi memory alloy electron beam welded joints are shown in Table 8.7. The fracture stress of the welded joints in the martensitic state is 410 MPa, and the fracture stress in the original parent phase state is 560 MPa; fracture occurs in the weld or in the semi-melted zone at the toe area, and small longitudinal and transverse cracks are present at the toe area. The cracks can be removed by grinding and the fracture stress rises to 710 MPa. the welded joint undergoes post-weld heat treatment (973 K × 120 min) with the grain refinement, elongation rises to 16% and the fracture stress is 660 MPa. but electron beam welding still has a negative effect on the shape memory effect. 3. Laser welding Laser welding can achieve the welding of shape memory alloy thin plate parts, and can obtain the shape memory effect and super elasticity similar to the base material, but the weld strength is low, and in the center of the weld is prone to cracking, which is mainly due to the melt zone of the joint produced a coarse cast organization and make the weld brittle. This conclusion was also confirmed by the Japanese scholars who used a 10 kW CO2 laser to weld thin plates of NiTi memory alloy with a thickness of 3 mm. For example, for the Ti–Ni50.7% memory alloy, the solid solution treatment condition for the base material is 973 K × 30 min and the aging condition is 673 K × 60 min in Ar gas. CO2 The process parameters for laser welding are: power of 6 kW, welding speed of 3.4 m/min and aging treatment of 673 K × 60 min in an Ar gas environment after the welding process. Table 8.8 lists the transformation temperatures of the Ti–Ni50.7% alloy, the base material after the aging treatment and

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Table 8.7 Mechanical properties of TiNi memory alloy electron beam welded heads Material

Status

Test conditions

Fracture stress/MPa

Elongation/%

Substrate

973 K × 60 min, water quenching

T < Mf

860

31

973 K × 60 min, water quenching

T > Af

740

26

No heat treatment, no grinding

T < Mf

410

9.8

No heat treatment, no grinding

T > Af

560

11

No heat treatment, grinding

T > Af

710

7.2

973 K × 120 min, heat treatment, grinding

T > Af

660

16

Welded joint

Table 8.8 Transition temperature of Ti–Ni50.7% alloy and laser weld weld metal Material

Transition temperature/K Af

As

Ms

Mf

Ti–Ni50.7% alloy

296

251

248

194

Substrate aging treatment (673 K × 60 min)

296

271

245

200

Laser welding metal

296

236

250

185

the laser welded weld metal, the data in the table shows that the phase transformation points of the base material and the weld metal are basically the same. The shape memory effect was evaluated at different test temperatures (from below the M s point to above the Af point), loaded at a strain rate of 1.6 × 10–4 /s, and after reaching a strain rate of 4% the load was removed and heated to the parent phase state to test its shape recovery and evaluate its shape memory effect. The test results showed that the laser welded head had the same shape memory effect as the parent material, but the tensile strength and fracture strain of the welded specimen were lower than that of the parent material (Table 8.9). Fracture occurred at the grain boundary of the columnar crystal in the center of the weld due to the fact that the grain boundary of the columnar crystal was perpendicular to the load and the presence of oxide inclusions on the grain boundary. Despite this, the welded specimens had a strain at break in excess of 6%, which is the maximum recoverable elongation in the polycrystalline TiNi metals. Therefore, the laser welding technique is very feasible for the Ti–Ni50.7% shape memory alloy. The functional properties of Ni–49.6% Ti shape memory alloy welded joints were investigated using a Nd:YAG laser welder. The results of tensile tests showed that the welded area of the specimen was less affected by the shape memory effect after annealing treatment at 900 °C × 1 h, while the superelastic properties of the specimen

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8 Connection of Functional Materials

Table 8.9 Mechanical properties of Ti–Ni50.7% alloy and its laser welded head Material

State

Test temperature/K

Tensile strength/MPa

Elongation/%

Ti–Ni50.7% metal

973 K × 30 min solution treatment

233

957

37

313

840

18

673 K × 60 min aging treatment

233

1224

15

313

1155

18

Welded

233

417

7.9

313

740

7.7

233

492

6.5

313

656

6.0

Laser welded joint

673 K × 60 min aging treatment

changed more than the unwelded specimen after the annealing treatment at 400 °C × 20 min. The shape memory effect and corrosion resistance of the joints were studied by welding Ti50 Ni50 and Ti49.5 Ni50.5 shape memory alloy sheets with a thickness of 2 mm using a CO2 laser. The results showed that the welded joints had a slight decrease in the martensitic phase change point and their shape memory effect was similar to that of the parent material. The Ti50 Ni50 alloy welds showed an increase in B2 phase, higher joint strength and lower elongation. The welded joints showed good corrosion resistance in H2 SO4 (1.5 mol/L) and HNO3 (1.5 mol/L) solutions. The results of superelasticity tests on Ti49.5 Ni50.5 alloy showed that the residual strain in the welded joints was high after the cyclic stress deformation due to the inhomogeneous weld microstructure. The 500 W pulse laser welder was used to laser spot weld Ti–50.6% Ni alloy wire with a diameter of 0.5 mm to research the microstructure and properties of the joint. The results show that: the melting zone of the laser spot welded joint consists of dendritic crystals, and the heat affected zone consisted primarily of the coarse equiaxed crystals near the weld part and fine equiaxed crystals near the base material part; laser welding causes the evaporation of Ni, reducing the Ni content in the joint and makes the joint. The evaporation of Ni caused by laser welding reduces the Ni content in the joint, and increases the phase transformation temperature of the joint; the tensile strength of the joint can reach 70% of the base material, and the recoverable strain reaches 92% of the base material. When TiNi alloy is used as a functional material with shape memory effect, the laser spot welding method is desirable. The above research results show that the melting welding method can be used to weld the TiNi shape memory alloys, the joint becomes brittle due to the dissolution of N, O, H and the generation of Ti2 Ni, TiNi3 brittle compounds; the heat affected zone metal is heated to make its microstructure coarse and the microstructure is varied, leading the shape memory effect and superelasticity of TiNi shape memory alloy to decrease. Therefore, in terms of ensuring the functional properties of the joint area,

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447

it is difficult to weld TiNi shape memory alloys by the conventional melt welding methods, except for the laser welding. (2) Solid-state welding of TiNi shape memory alloy Solid-state welding methods (such as resistance welding, friction welding and diffusion welding) have the advantages of small changes in the microstructure of the metal in the joint area, the ability to obtain joints at a lower temperature (compared to fusion welding) and the absence of molten metal, which are very advantageous for the welding of TiNi shape memory alloys and for ensuring the properties of the joint area. 1. Resistance welding ➀ Resistance spot welding For the spot welding test of cross-lap joints in the wire mesh structure of the Ti– 55.2% Ni shape memory alloy with a diameter of 0.5 mm, two process methods, the precision time-controlled AC spot welding and energy storage spot welding, were compared, and the effect of argon gas protection was investigated. The results show that the spot welding of the TiNi alloy tends to absorb N, O and H, which degrades the mechanical properties and shape memory effect of the joint. Therefore, using of argon gas protection during welding is essential. The shape memory recovery rate of the welded joints obtained by both process methods can reach more than 98%. The mechanical properties AC spot welding method is superior to the stored energy pulse spot welding, and the maximum shear strength of AC spot weld joints and stored energy pulse spot welded joints are 700 and 500 MPa, respectively; their maximum tensile strengths are 1200 and 1000 MPa, respectively. Analysis of the properties of the TiNi alloy base material, welded joints and post-weld heat treatment organization shows that the TiNi shape memory alloy is basically homogeneous in all regions of the welded joints and the base material compositions after the spot welding. The microstructure of the weld after welding without heat treatment is dominated by the high-temperature phase, and the weld joint after the same heat treatment as the base material is basically the same as that of the base material after the heat treatment, consisting of the high-temperature also the martensite phase. Through the variable temperature dynamic analysis of the welded joint, it is proved that the welded joint has the function of thermoelastic martensitic phase transformation and the property of shape memory effect. ➁ Resistance butt welding For resistance butt welding of the TiNi shape memory alloy wire with a diameter of 0.73 mm, the effect of weld top forging force and welding current on the mechanical properties and shape memory effect of the joint can be researched, and a map of the area suitable for the welding conditions and shape recovery rate should be given. The shape memory properties of the welded parts can be evaluated by using the bending test method.

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8 Connection of Functional Materials

For the precision pulse resistance butt welding of TiNi shape memory alloy, the effect of parameters, such as welding current, welding pressure, top forging pressure and shielding gas on the mechanical properties and shape memory effect of welded joints was analyzed. The parameters derived from the test to obtain the highest shape recovery rate welded joints are: welding heat of 75%, excitation current of 2 A, tuning extension length of 5.0 mm, welding retention of 2.5 mm, post heat treatment of 10% and post heat treatment time of 40–60 cycles. Resistance welding is a favorable method for joining TiNi alloys, but the flexibility of this method is limited. For example, by the complexity of the workpiece shape and joint, as well as the size of the dimensions. 2. Friction welding Using friction welding also the post-weld heat treatment, Ti50–Ni50 (molar fraction, %) metal rods with a diameter of 6 mm (length of 100 mm) could be successfully joined with the excellent results. The top forging pressure used in friction welding was 39.2–196.1 MPa, and the post-weld heat treatment conditions were 773 K × 30 min and the ice water quenching. The mechanical properties and shape memory effects of the welded joints after heat treatment were good, and the transition temperatures of the friction welded joints under different process conditions are shown in Table 8.10, which shows that the heat treated welded joints have almost the same transition temperature as the TiNi base material. Stress–strain measurements show that the shape memory effect of the friction welded joint after the heat treatment is better than that of the base material. Severe thermal extrusion deformation can be produced in the weld zone under the action of the friction welding, also the fine microstructure can be obtained, which is beneficial for the shape memory effect. However, the friction welding can not guarantee the geometric accuracy of the joint bond surface. Therefore, the geometric accuracy of the workpiece joint is an unavoidable problem in friction welding of TiNi shape memory alloys. Energy storage friction welding is capable of joining non-axisymmetric components, but it requires the applications of the rapid thermal cycling and high axial forces during a welding process to cause the thermally deformed plastic metal to extrude out of the bond surface to obtain a dense joint, but this can have a detrimental effect on the shape memory effect of the TiNi alloys. Table 8.10 Transformation temperature of friction weld under different process conditions Base metal and joint status

Transition temperature/K Ms

Mf

As

Af

Ti50–Ni50 memory alloy

309.0

277.5

314.2

331.0

Upsetting pressure 39.2 MPa, post weld heat treatment

309.5

279.0

316.3

332.0

Upsetting pressure 196.1 MPa, post weld heat treatment

309.2

276.3

316.3

334.5

Upsetting pressure 39.2 MPa, welded

245.0

216.4

287.4

310.0

Upsetting pressure 196.1 MPa, welded

267.6

216.7

286.9

309.8

8.2 Shape Memory Alloy to Metal Connection

449

3. Diffusion welding Diffusion welding achieves the joining of materials by applying a certain pressure at the high temperatures free of the significant macroscopic deformation of the joined workpiece. An intermediate alloy can be filled at the bond surface, which is a very promising method for joining the different shape-memory alloys. However, the temperature of diffusion welding is generally higher than the annealing temperature of the TiNi shape memory alloys, which is detrimental to the shape memory effect of the base material. The instantaneous liquid phase diffusion welding (TLP) study of NiTi alloy revealed that a layer of the Ni2 AlTi compound was formed at an joint interface. The diffusion of Ti in the NiTi alloy into the joint during the welding process leads to a decrease in the solid-phase line temperature of the NiTi alloy, resulting in its partial melting during a welding process. The diffusion of elemental Cr in the NiAl alloy into the joint and the NiTi matrix leads to the formation of α-Cr phase in the NiTi matrix, which can be eliminated by the post-weld heat treatment, reducing the effect on the memory effect of the NiTi alloy. It is shown that the use of instantaneous liquid phase diffusion welding method for joining TiNi shape memory alloys, where the chemical composition and microstructure of the welded joint can be brought close to the base material by prolonged diffusion or post-weld heat treatment, has great potential for joining the TiNi shape memory alloys, and its successful application depends on the optimization of parameters for a given alloy system. (3) Brazing of TiNi shape memory alloy 1. Brazing of homogeneous joints Japanese scholars have developed brazing materials and brazing fluxes capable of brazing Ti–55.75% Ni shape memory alloys in the atmosphere. The braze A-1 composition (mass fraction) based on BAg7 is Ag59%, Cu23%, Zn15%, Sn1%, Ni2%. The brazing agent composition (mass fraction) is 25% AgCl, 25% KF, 50% LiCl, which enables the Ag-based brazing material to wet well on the TiNi shape memory alloy. The brazing process is carried out in the two steps: the first step is pre-fused brazing, where the developed brazing flux is applied to the joint part of the specimen so that the brazing flux melts and is fused to the joint part of the specimen; the second step is joining, where the joint part of the specimen with the pre-fused brazing layer is coated with a universal brazing flux for silver brazing, and then the two specimens to be joined are assembled together, pressed with 100 g mass and brazed in a furnace. The test results showed that the A-1 braze with 2% Ni significantly improved the strength of the joint with a maximum shear strength of about 300 MPa compared to the conventional braze BAg7, with which the maximum shear strength of the BAg7 braze was compared to about 200 MPa. The Ti50 Ni50 shape memory alloy was brazed with pure Cu and Ti–15Cu–15Ni foils in an infrared heating furnace in an environment of the argon gas to research the microstructure of the braze seam and the shape memory properties of the joint.

450

8 Connection of Functional Materials

The results show that when pure Cu brazing material is used, the braze joint consists of the Cu-rich phases, the CuNiTi and Ti(Ni, Cu) phases, where the Cu-rich phase disappears rapidly within the first 10 s of brazing and the joint consists of CuNiTi and Ti(Ni, Cu) eutectics. The CuNiTi phase gradually decreased with the extension of brazing time; the shape recovery rate of the brazed joint reached 99.9% at 130 °C when the brazing temperature was 1150 °C and the brazing time was 300 s, which was comparable to that of the substrate, and the extension of the brazing time helped to improve the shape recovery rate of the joint. In contrast, when the Ti–15Cu– 15Ni brazing material was used, the Ti2 (Ni, Cu) brittle phase was formed in the joint, resulting in the unsuccessful bending test, and the brittle phases could not be eliminated by increasing the brazing temperature or extending the brazing time. 2. Brazing of heterogeneous joints Brazing the connections of TiNi shape memory alloys to 304 austenitic stainless steel were achieved using the Ag–Cu eutectic brazing material BAg28, BAg28 with 0.5% and 3% Ni additions (composition wt%: Ag 72.6, 71.5 and 77%, Cu 27.4, 28 and 20%, Ni 0, 0.5 and 3%, respectively). The brazed layer at the connection site was kept fixed and welded in an infrared heating furnace in an argon gas stream at a pressure of 0.5 MPa. The results are shown as follows: ➀ When brazing with BAg28 brazing material, a lower temperature or shorter holding time leads an uniform reaction layer to be formed on the joint surface with a joint strength of 200–250 MPa and a maximum strength of 270 MPa. The fracture of the welded part occurs near to the FeTi compound layer formed on the brazing material and the interface zone. ➁ When brazing with Ni-added brazing media, the dissolution of Fe and Ti is inhibited and no FeTi compound layer is formed, forming the Fe- and Ni-rich solid solution layer on the 304 stainless steel side and a Ni3 Ti layer on the shape memory alloy side. ➂ For brazed parts with Ni brazing material, breakage occurs in the Ni3 Ti layer and NiTi layer formed at the interface zone, as no FeTi compounds are formed, and the maximum fracture strength of the welded parts can increased to about 400 MPa. Brazed joints of TiNi shape memory alloy with the pure Ti can be achieved using the Ag–Cu (BAg-8) and Cu–Ti–Zr (MBF5004) brazing materials. The experimental results show that with BAg-8 brazing material, when the brazing temperature is lower than 1153 K, four compound layers are formed in the joint with a tensile strength of up to 330 MPa, and the fracture occurs in the Ti–Cu intermetallics’ layer between Ti and brazing material; when the brazing temperature is higher than 1193 K, the two layers are formed in the joint with a tensile strength of up to 350 MPa, and the fracture occurs in the brazed α-Ti also Ti2 (Ni, Cu) layers. The thickness of the diffusion layer at this point is three times higher than that of the brazing materials, indicating the partial melting of the TiNi memory alloy base material near to an interface zone during a brazing process.

8.2 Shape Memory Alloy to Metal Connection

451

MBF5004 brazing material can be used, the brazed joint microstructure and the fracture sites were similar to those using BAg-8 brazing material, but the joint strength was higher and close to the strength of pure Ti substrate. Microbeam plasma arc welding, energy storage welding and the laser brazing were used to compare the microstructure and properties of TiNi shape memory alloy and stainless steel joints, and the results showed that using microbeam plasma arc welding and energy storage welding, the weld joints were extreme brittle due to the melting of the stainless steel and the TiNi shape memory alloy, the formations of the cast tissue and brittle compounds at the joints, which varied the composition and the microstructure of the TiNi shape memory alloys. The tensile strength is low, and cannot withstand the bend loads, the hardness of the heat affected zone increases, and the joint is brittle fracture. Therefore, the TiNi shape memory alloy overheating welding should be avoided in order to improve the performance of the heterogeneous joints, and minimize the melting of the two base materials or welding the excess molten metal out of the weld zone. A new AgCuZnSn silver-based brazing material is suitable for brazing TiNi shape memory alloy with stainless steel can be used, which can be applied in the medical field. The silver-based brazing material composition (wt%): Ag51–53%, Cu21–23%, Zn17–19%, Sn7–9%. The solid-phase line temperature is 590 °C and the liquidphase line temperature is 635 °C. The brazing material mainly consists of α-Ag solid solution, α-(Cu, Zn) solid solution and the Ag–Cu eutectics. This brazing material can be used to braze the TiNi shape memory alloy with the stainless steel, and the metallurgical bond at the interface of the brazed joint was straight and dense. The joint strength can be up to 360 MPa by selecting the appropriate laser brazing process parameters, while the shape memory effect and superelastic properties of TiNi shape memory alloy are less lost. The composite orthodontic archwire made of the TiNi shape memory alloy orthodontic wire and stainless steel orthodontic wire, that were connected by the laser brazing, which was applied to the orthodontic clinic, resulting in the excellent orthodontics. For the TiNi shape memory alloy heterogeneous material connections, the brazing and instantaneous liquid phase diffusion welding can be used to obtain the weld joints with the better performance at the temperatures lower than the annealing temperature of TiNi shape memory alloy, and the shape memory effect and superelasticity of the base material are less affected and should be one of the concerns.

8.2.3 Resistance Brazing of TiNi Shape Memory Alloys For the TiNi shape memory alloy, Xue Songbai et al. conducted an experimental study by mean of using a resistance brazing method with the argon gas protection. Because of the poor thermal conductivity and high resistance of the TiNi alloy, the resistance brazing method has short time, the low welding heat, the concentrated heating, the low thermal influence, and good wetting of the brazing material to the

452

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Table 8.11 Chemical composition of TiNi shape memory alloy % Ti

Ni

Mn

Si

Fe

43.58

Bal

0.01

0.005

0.005

Table 8.12 Chemical composition of brazing materials (mass fraction) % Cu

Ni

Mn

Fe

Al

Si

55.65

42.20

1.47

0.50

0.10

0.084

base material, which are not only beneficial to the improvement of the brazing seam strength, but also can reduce the loss of shape memory effect of the joint. (1) Test material and welding method The size of the TiNi memory alloy used for the test was a 2.5 × 1.2 mm flat wire with the main chemical compositions, which was shown in Table 8.11. The size of the brazing material was a 1.0 × 0.24 mm thin strip of CuNi with the chemical composition shown in Table 8.12. The brazing agent used was a modified version of the existing brazing agent (mass fraction, %) AgCl25–KF25–LiCl50 with a certain amount of Ax By . The substrate and brazing material are cleaned and degreased with acetone before the welding process, then the base material is soaked in the hydrofluoric acid and nitric acid solution for 10–15 min (room temperature) to remove the oxide films on the surface, and finally the base material and brazing material should be cleaned with the alcohol and dried naturally. The welding equipment for the test is a self-developed DN25 CNC AC resistance welding machine with a rated primary current of 66.8 A, rated power of 25 kW and secondary no-load voltage of 4 V. It can realize the synchronous and precise controlling of the multiple parameters during a welding process, and can control the size of welding heat and the length of welding time, and the welding pressure is accurately controlled by a solenoid valve according to the welding requirements. It adopts internal water-cooled electrode and argon gas protection synchronized with welding. The welded joint is in the form of a lap joint, and the contact surface between the thin strip of brazing material and the base material is coated with a certain amount of the brazing flux (pre-stirred into a paste), and is protected with and without argon gas, respectively. After a pre-pressure stage, an energized welding stage and a maintenance pressure stage to complete a welding process. The selected welding process parameters are shown in Table 8.13. for 2.5 × 1.2 mm TiNi wire for the best process parameters: the welding heat adjustment for 1, welding pressure of 0.14 MPa, welding time for 5 cycles. After welding using Zwick I microcomputercontrolled electronic tensile testing machine to determine the shear strength of the joint.

8.2 Shape Memory Alloy to Metal Connection

453

Table 8.13 Welding process parameters Welding heat adjustment

0

1

2

3

4

Welding time t/cycles

–

–

5

–

–

Welding pressure p/MPa

–

–

0.14

–

–

(2) The effect of welding heat on the mechanical properties of the joint Since the weld heat regulation directly affects the heat input to the joint (when the weld heat regulation is 0, the conduction angle of the SCR is 30°; when the weld heat regulation is 10, the conduction angle of the SCR is 90°), it largely determines the quality of the joint. Figure 8.16 shows the relationship between weld heat regulation and the shear strength of the joint. It can be seen that with the welding heat regulation increases, the strength of the joint also increases, reaching a peak at a welding heat regulation of 1. The strength starts to decrease as the welding heat regulation increases further. The shear strength of resistance brazed joints of TiNi shape memory alloy can reach up to 577 MPa using thin strips of the CuNi brazing material with the flux. The analysis suggests that when the welding heat is regulated, the heat gained in the brazed joint area is small, the temperature in the joint area is lower, and the activity of the brazing flux is reduced. The thermoplastic deformation of the joint surface is small and insufficient to break the oxide film on the joint surface. This affects the removal of the oxide film and prevents the wetting of the brazing material to the substrate, and the joint brazing adhesion is low. When subjected to shear, the joint has a small force area at the joint surface and is easily broken at the lap part, so the strength is low. The macroscopic morphology of the joint section also proved this point. The original appearance of the base material in the joint area of the specimen was clearly visible, and no trace of the brazing material wetting was noted. When the welding heat regulation is large, although the joint action of the brazing flux and the thermoplastic deformation of the brazing joint surface enables the oxide Fig. 8.16 Relationship between weld heat regulation and shear strength

454

8 Connection of Functional Materials

film on the surface of the subsrate to be removed well, the high temperature leads to a violent reaction between the brazing material and the base material and the formation of a new compound phase. At the same time, the larger heat input has a greater thermal impact on the base material, the coarsening of the base material grain, the TiNi3 phase will grow sharply, the content of the brittle phase of the joint is extreme high, so the dislocation density increases, forming a potential source of the cracks in the joint and heat-affected zone, the plasticity of the joint is poor, and it is easy to fracture in the brazing seam and the heat-affected zone when the force is applied. When the welding heat is properly regulated, a braze joint with high strength and plasticity would be obtained. (3) The influence of welding parameters on the mechanical properties of the joint According to the formula for calculating the heat of resistance welding. Q = I 2 Rt

(8.1)

where I is the current intensity, A; R is the contact resistance, Ω; and t is the welding time, it can be seen that the heat Q is linearly related to the welding time t when the welding heat is adjusted to a certain level and the base material is treated with the same pre-weld treatment. The relationship between the welding time and the shear strength is shown in Fig. 8.17. It shows that when the welding time is short, the heat is small, the brazing flux is not active enough and the oxide film at the joint interface cannot be removed sufficiently. Meanwhile, the interaction time between the melted brazing material and the base material is short, the brazing material can not be completely melted, resulting in a low brazing rate, can not form a compact brazing seam. When the welding time is too long, the excessive heat makes the action of the brazing material and the base material violent, and the brittle phases can be formed at the brazing seam. At the same time, the base material grains are coarsened due to an excessive heat input, causing a reduction in joint strength. The strength of brazed joints protected with argon is about 25% higher than that without argon protection. Fig. 8.17 Relationship between welding time and shear strength

8.2 Shape Memory Alloy to Metal Connection

455

Without the argon protection, the degree of the interfacial bond plays a decisive role in the mechanical properties of the base material when the brazing time is short. With the extension of the brazing time, the heat gradually increases and the protective effect of argon is obvious. As the heat input is large, if there is no argon protection, the joint high temperature zone by the nitrogen, the hydrogen, the oxygen and the other intrusions and the generations of the various oxides or nitrides and the other inclusions affect the brazing filler, so that the joint is severely embrittlement, the joint strength is reduced. The welding time is not an independent parameter, it depends on the size of the welding heat regulation, which could determine the welding temperature. The relationship between welding pressure and mechanical properties of the joint is shown in Fig. 8.18. The strength of the joint increases with increasing pressure and reaches a maximum at 0.14 MPa. The effect of welding pressure on the strength of the joint is not as significant as the welding heat, and the range of variation in the strength of the joint is smaller as the welding pressure varies. The resistance brazing technology has a loose requirement for the welding pressure and a wide range of pressure adjustment. Increasing the pressure increases the thermal deformation of the joint metal, leading to grain refinement in the brazing bond area and improving the strength of the joint. However, excessive welding pressure will cause significant macroscopic deformation of the welded parts, which will reduce the strength of the joint. During the resistance brazing process, due to the different surface roughness of the base material, a certain amount of thermoplastic deformation is produced at the joint surface of the thin strip of brazing material and the base material under the action of moderate welding pressure, and a microscopic full contact is achieved between the interfaces, avoiding the spattering due to the steep increase in transient welding current caused by the microscopic local contact at the joint surface. The welding pressure is able to break the oxide film at the joint surface under the action of heat, and the brazing agent can penetrate between the base material and the oxide film through the oxide film rupture, and achieve the purpose of removing the film Fig. 8.18 Relationship between weld pressure and shear strength

Air cooling

456

8 Connection of Functional Materials

through the peeling and dissolving of the oxide film. However, the excessive pressure causes a large thermoplastic deformation of the joint surface, and the molten brazing material is extruded from the joint surface, resulting in the direct contact between the joint surface and the substrate, and it is difficult to achieve sufficient diffusions of the atoms due to the low temperature, which is not conducive to the crystallization process of the brazing seam, thus reducing the strength of the joint. (4) Analysis of base material and brazing organization The metallographic analysis of the base material and resistance brazed joints by optical microscopy showed that there was no obvious heat affected zone at the joints and the brazing seam was continuous and dense, the brazing process had little thermal effect on the base material and the molten brazing material was able to fill the seam adequately. The braze joint is mainly composed of β-phase and Ni3 Ti2 phase, which is due to the rapid cooling of the braze zone directly by the molten brazing material, which retains more high-temperature phases, which is beneficial to the TiNi shape memory effect and mechanical properties. The shape memory properties and superelasticity of the welded joint are influenced by the welding time and welding temperature, but also by the geometry of the base material itself, and also by the size of the heat capacity required for brazing. Comparing to the resistance spot welding, resistance brazing does not require the weld temperature to reach the melting point of the base material, only the low melting point of the brazing material, which reduces the thermal impact on the base material. The heat source of resistance brazing is generated inside the welded part and heats the weld zone more rapidly and intensively than the external heat source in fusion welding, where the internal heat source heats up the entire weld zone. In order to obtain a reasonable temperature distribution, water-cooled electrodes can be used to achieve heat dissipation by rapid cooling of the weld zone. At the same time, the relatively small amount of melted brazing material excludes the adverse effect of the brittle relative strength formed by the large amount of melted brazing material during the brazing process, and also reduces the thermal effect on the base material, thus minimizing the loss of shape memory effect of the joint. Compared with the conventional brazing methods, the resistance brazing compensates for the formation of the new phases, the grain coarsening of the substrate, and the low brazing seam strength due to the long heating and holding time. The results of the study showed that the shear strength of the joints brazed using resistance brazing increased by a factor of one over that of resistance spot welded joints and by about 70% over that of the joints brazed in the furnace. The shape memory effect of the joint is basically ensured due to the relatively small thermal influence on the base material. Therefore, the TiNi alloy wire has the good application prospects by using the resistance brazing technology.

8.2 Shape Memory Alloy to Metal Connection

457

8.2.4 Transition Liquid Phase Diffusion Welding of TiNi Alloy to Stainless Steel The TiNi shape memory alloy is expensive, in practical applications it is connected with excellent performance, the low price of the stainless steel is an important way to reduce the costs and expand its application. TiNi alloy and stainless steel physical and chemical properties (such as melting point, thermal conductivity, linear expansion coefficient, crystal structure, etc.) are very different, the use of fusion welding method when the joint is prone to the stress concentration and cracking, and the bonding interface is easy to form TiFe, TiFe2 , TiC and the other brittle phases, seriously affecting the performance of the joint. The using of the AgCu metal foil as an intermediate transition layer and the experimental study of transition liquid phase diffusion welding (TLP-DB) for TiNi/stainless steel can expand the application of the TiNi shape memory alloy. (1) Material and welding method The test materials were Ti50.2Ni49.8 (mass fraction, %) and 304 stainless steel (18-8 steel) with the physical properties shown in Table 8.14. AgCu28 metal foil with a thickness of 50 μm was used as an intermediate layer with a melting point of 779 °C and a tensile strength of 343 MPa at room temperature. a lap joint was used, with a specimen size of 30 × 10 × 2 mm and a lap length of the welded surface was first polished with the sandpaper, ultrasonically cleaned with acetone for 10 min and dried. The prepared materials were assembled in the order of TiNi/AgCu/304 stainless steel. The process parameters: the connection temperature T is 820–900 °C, the holding time t is 20–100 min, connection pressure p is 0 to 0.1 MPa, the vacuum is 1.0 × 10–2 ~ 1.0 × 10–3 Pa. Scanning electron microscopy and X-ray diffractometer were used to analyze the microstructure of the joint interface. Shear strength was used to evaluate the joint strength at each process parameter, and the average of the shear strength of at least three specimens was taken. The post-welded joints were subjected to the shear experiments on the MTS810 material testing machine with a load speed of 0.5 mm/min. (2) Effect of process parameters on joint strength ➀ Effect of heating temperature The variation of shear strength of the joint with the joint temperature for t = 60 min and p = 0.05 MPa is shown in Fig. 8.19. With the increase of the connection temperature, the shear strength of the joint first increases and then decreases. At a heating temperature of 820 °C, the AgCu interlayer only formed a diffusion layer with the parent TiNi with a width of about 2 μm, and the joint interface with the stainless steel did not react sufficiently and the demarcation line was obvious. The fracture morphology is a small number of tough nests distributed on a large lamellae, and brittle fracture is predominant. This indicates that only a small amount of reaction products were formed and grew at the interface, no continuous diffusion layer was

Density/g cm−3

6.4–6.5

7.9–8.0

Substrate

TiNi

304 stainless steel

1440

1310

Melting point/°C

17

10

Coefficient of linear expansion/× 10–6 °C−1

Table 8.14 Physical properties of the test base material

0.16

0.21

Thermal conductivity/J (cm s °C)

726

940

Tensile strength/MPa

379

444

Yield strength/MPa

59

9

Elongation of section/%

458 8 Connection of Functional Materials

8.2 Shape Memory Alloy to Metal Connection

459

Fig. 8.19 Effect of connection temperature on the shear strength of a joint

formed, and the interfacial metallurgical bonding rate was low. The maximum shear strength of the diffusion joint was 239.4 MPa when the temperature was increased to 860 °C. When the temperature increases to 900 °C, the interface zone between the middlelayer and the substrate on both sides disappears, and the weld is relatively narrow, but the thickness of the diffusion layer increases significantly. The fracture morphology shows the coarse grains, the low melting point eutectic melting of the intermediate layer fills the grain gap, and the grain morphology is not obvious. This indicates that the joint has shifted from ductile fracture to brittle fracture. Therefore, the shear strength of the TiNi/stainless steel transition liquid phase diffusion welded joints is related to the degree of interfacial diffusion reactions also the grain sizes. ➁ Effect of holding time The variation of shear strength of the joint with holding time at T = 860 °C and p = 0.05 MPa is shown in Fig. 8.20. At this temperature, the intermediate transition layer generates a low-melting eutectics, and the liquid metal is obtained at a lower temperature. With the welding time increases, the molten liquid intermediate layer gradually spreads to the surface of the base metal. At the same time, the TiNi and stainless steel interface at the partial dissolution, and diffusion into the liquid metal, so that the liquid metal continues to increase. With the diffusion reaction, the composition of the intermediate layer varies, the melting point of the liquid metal increases, and finally deposited on the surface of the substrate. With the extension of the holding time, the solid Ag enriched in the weld zone continuously diffuses into the substrate, so that the microstructure of the weld zone becomes gradually homogenized, a good performance weld joint is obtained. When the holding time is short, the intermediate layer elements do not have time to diffuse into the base material, and the interface has not formed a metallurgical bonding layer or the bonding rate is low, especially on the stainless steel side, there are also a large number of pores. Fracture morphology is laminar tearing, indicating that the interface metallurgical bonding is poor. After the holding time exceeded 60 min, the interfacial reaction layer thickened and increased the stress in the interfacial bonding area due to the mismatch of physical properties, resulting in cracks on the TiNi side of the joint, which led to a significant decrease in the joint strength. The

460

8 Connection of Functional Materials

Fig. 8.20 Effect of holding time on shear strength of joints

joints exhibited the mixed brittle and ductile fracture due to the precipitations of the second phase hard mass points at the interface zone during the reaction process, and grew into a lamella. The holding time determines the degree of diffusion of elements at the interface of the transition liquid phase diffusion welding and is an important parameter for the formation of a uniform reaction layer in the joint. ➂ Effect of pressure on joint strength As seen in Fig. 8.21, when the connection pressure is small (T = 860 °C, t = 60 min), only a small number of microscopic bumps on the surface of the material to be welded occur in physical contact, and the plastic deformation is small, providing little deformation energy, the weld fit rate is small, and the joint strength is not high. When the pressure increases to 0.05 MPa, the effective contact area and deformation energy increase, the gap between the intermediate layer and the base material decreases, the diffusion of interfacial elements accelerates, and the joint strength is higher. However, when the pressure is too high, the liquid intermediate layer may be extruded during the joining process, which reduces the reaction and diffusion of interfacial elements, and the joint strength decreases instead. The shear strength of the joint also tends to increase and then decrease with the change of the connection pressure, but the decrease is not large. The highest shear strength of the joint was found at the connection pressure of 0.05 MPa.

Fig. 8.21 Effect of connection pressure on the shear strength of a joint (T = 860 °C, t = 60 min)

References

461

Table 8.15 % energy spectrum analysis of the composition of the transition liquid phase diffusion weld head at each point Measuring point Ti

Ni

Ag

Cu

1

47.78 51.05 –

–

2

32.23 16.79 7.63

3.06

3

1.67

0.11 58.35 37.62

Fe

Cr

1.17 –

Si

Mn

Possible phase

–

–

–

33.07 7.22

–

–

TiFe

1.36 0.89

–

–

AgCu

2.04 –

TiFe2

4

23.95

7.87 14.72 3.15

40.39 7.88

5

–

7.73 –

70.29 18.88 1.63 1.48 –

–

(3) Joint organization and interfacial reaction layer The results of the energy spectrum analysis of the composition of TiNi/304 stainless steel transition liquid phase diffusion weld head at each point at T = 860 °C, t = 60 min, p = 0.05 MPa are shown in Table 8.15. According to the energy spectrum analysis, the Fe in stainless steel and Ti in TiNi cross the intermediate transition layer during the joining process and participate in the interfacial reaction occurring in the intermediate layer. The reaction layer on the TiNi side is mainly composed of Ti, Ni, Fe elements and a small amount of Ag, Cu, Cr elements, the interfacial reaction products are dominated by Ti(Ni, Fe). At 860 °C, the effective diffusion coefficient of Ag in Fe is larger than that of Cu in Fe, and the content of Ag in the interfacial bonding area on the stainless steel side is more than that of Cu. The fracture occurs at the reaction interface between TiNi and the interlayer. The X-ray diffraction test results show that in addition to the diffused α-Ag and brittle phases such as TiNi2 , TiFe at the interface, a Ti3 Ni4 compound phase with a co-grid relationship with the matrix TiNi is found. The Ti3 Ni4 phase is the main factor for the dual-range and full-range memory effect of the TiNi shape memory alloy, and the joint has some shape memory effect.

References 1. Ren JL, Wu AP (2000) Connection of advanced materials. Machinery Industry Press, Beijing 2. Wei W, Feng Y, Wu XZ et al (1991) A review of the research status of niobium–titanium low-temperature superconducting materials welding technology. Adv Titan Ind 1:12–16 3. Zou G, Hao AP, Ren JL, et al (2001) Status and prospects of research on the connection of high Tc oxide ceramic superconducting materials. Mater Guid 15(12):27–28 4. Wang F, Chen ZL (2002) Application of shape memory alloy materials. Mater Mech Eng 26(3):5– 8 5. Li MG, Sun DQ, Qiu SM et al (2006) Research progress on TiNi shape memory alloy joining technology. Mater Guid 20(2):121–125 6. Xue SB, Lv XC, Zhang HW (2004) Resistance brazing technique for TiNi shape memory alloy. J Weld 25(1):1–4 7. Yingling W, Hong L, Zhuoxin L et al (2008) Study on the instantaneous liquid phase diffusion welding process of TiNi shape memory alloy with stainless steel. Mater Eng 9:48–51 8. Zhu SC, Lu XF (2009) Application of shape memory alloys in flange sealing connections. Nucl Power Eng 30(3):136–140