High Interstitial Stainless Austenitic Steels [1 ed.] 978-3-642-33700-0, 978-3-642-33701-7

Citation preview

Engineering Materials

For further volumes: http://www.springer.com/series/4288

Hans Berns Valentin Gavriljuk Sascha Riedner •

High Interstitial Stainless Austenitic Steels

123

Sascha Riedner Deutsche Edelstahlwerke GmbH Kamen Germany

Hans Berns Ruhr-University Bochum Germany Valentin Gavriljuk G.V. Kurdyumov Institute for Metal Physics Kiev Ukraine

ISSN 1612-1317 ISBN 978-3-642-33700-0 DOI 10.1007/978-3-642-33701-7

ISSN 1868-1212 (electronic) ISBN 978-3-642-33701-7 (eBook)

Springer Heidelberg New York Dordrecht London Library of Congress Control Number: 2012949081 Ó Springer-Verlag Berlin Heidelberg 2013 This work is subject to copyright. All rights are reserved by the Publisher, whether the whole or part of the material is concerned, specifically the rights of translation, 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. Exempted from this legal reservation are brief excerpts in connection with reviews or scholarly analysis or material supplied specifically for the purpose of being entered and executed on a computer system, for exclusive use by the purchaser of the work. Duplication of this publication or parts thereof is permitted only under the provisions of the Copyright Law of the Publisher’s location, in its current version, and permission for use must always be obtained from Springer. Permissions for use may be obtained through RightsLink at the Copyright Clearance Center. Violations are liable to prosecution under the respective Copyright Law. 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. While the advice and information in this book are believed to be true and accurate at the date of publication, neither the authors nor the editors nor the publisher can accept any legal responsibility for any errors or omissions that may be made. The publisher makes no warranty, express or implied, with respect to the material contained herein. Printed on acid-free paper Springer is part of Springer Science+Business Media (www.springer.com)

Preface

A close cooperation between the Institute of Metal Physics in Kiev, Ukraine, and the Chair of Materials Technology at the Ruhr University Bochum, Germany, has been going on now for more than 20 years. During the first decade the joint effort centered on high nitrogen steels (HNS) and the partners published a book on this subject in 1999. Already then they had shown that combined alloying of martensitic stainless steels with carbon + nitrogen considerably improved structure and properties. These findings were transferred to the hardenable stainless bearing steel CRONIDURÒ used, e.g., in aviation and to SolNitÒ which allows case hardening of stainless steel with nitrogen instead of carbon. The second decade of partnership was dedicated to extending the beneficial C+N concept to new stainless austenitic grades called high interstitial steels (HIS). The results compiled in the present book have two major targets. On the scientific side the structure/ property relation starts at the electron structure and is carried on to the macroscale explaining the superior performance of HIS. The engineering aspects cover major steps of industrial manufacture and possible applications. Compared to similar HNS the new HIS do without costly pressure or powder metallurgy. Thus the contents are of interest to materials scientists working in R&D but also to engineers in design, manufacture, and materials selection. The authors thank Prof. Dr. Bela Shanina (Theoretic Physics), Dr. habil. Yuri Petrov (Electron Microscopy), Dr. Andrij Tyshchenko (Mössbauer Spectroscopy) in Kiev and Dr.-Ing. Fabian Schmalt, Dr.-Ing. Lais Mujica-Roncery, and Dipl. Ing. Nilofar Nabiran in Bochum for their most valuable contributions. Thanks are also extended to other researchers, students, and technical staff involved in the development of HIS and to Miriam Rockenbach, Agnes Krolik, and Dipl. Ing. Fabian Pöhl for preparing the final manuscript. The authors are grateful to Prof. Dr.-Ing. Werner Theisen, Head of Chair in Bochum, for his continuous support to the HIS project and to the German research foundation (DFG) for

v

vi

Preface

sponsoring part of the work. They feel also indebted to the National Academy of Science and the Science and Technology Center in Ukraine for financial support. Last but not least, we thank several companies for melting and processing new HIS (see Chap. 5). Summer 2012

Hans Berns Valentin Gavriljuk Sascha Riedner

Contents

1

Introduction . . . . . . . . . . 1.1 High Nitrogen Steels . 1.2 High Interstitial Steels 1.3 Aim . . . . . . . . . . . . . 1.4 Procedure . . . . . . . . . References . . . . . . . . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

1 1 2 3 4 5

2

Constitution . . . . . . . . . . . . . . . . . . . . 2.1 General Remarks . . . . . . . . . . . . . 2.2 Variation of Interstitial Content . . . 2.3 Effect of C/N Ratio . . . . . . . . . . . 2.4 Variation of Substitutional Content 2.4.1 Chromium and Manganese . 2.4.2 Molybdenum and Copper. . 2.4.3 Tramp Elements . . . . . . . . 2.5 Selection of Steels . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . .

. . . . . . . . . .

. . . . . . . . . .

. . . . . . . . . .

. . . . . . . . . .

. . . . . . . . . .

. . . . . . . . . .

. . . . . . . . . .

. . . . . . . . . .

. . . . . . . . . .

. . . . . . . . . .

. . . . . . . . . .

. . . . . . . . . .

. . . . . . . . . .

. . . . . . . . . .

. . . . . . . . . .

. . . . . . . . . .

. . . . . . . . . .

. . . . . . . . . .

. . . . . . . . . .

. . . . . . . . . .

. . . . . . . . . .

7 7 8 9 11 11 14 16 17 19

3

Structure . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.1 As-Quenched . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.1.1 Electron Structure: Calculated and Measured 3.1.2 Atomic Distribution. . . . . . . . . . . . . . . . . . 3.1.3 Chemical Nanoscale Homogeneity . . . . . . . 3.2 Structural Change by Loading. . . . . . . . . . . . . . . . 3.2.1 Tensile Straining. . . . . . . . . . . . . . . . . . . . 3.2.2 Effect of Subzero Temperature . . . . . . . . . . 3.2.3 Effect of Strain Rate . . . . . . . . . . . . . . . . . 3.2.4 Effect of Cyclic Loading . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

. . . . . . . . . . .

. . . . . . . . . . .

. . . . . . . . . . .

. . . . . . . . . . .

. . . . . . . . . . .

. . . . . . . . . . .

. . . . . . . . . . .

. . . . . . . . . . .

. . . . . . . . . . .

21 22 22 28 36 47 47 59 61 68 77

vii

viii

Contents

4

Properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.1 Mechanical Properties . . . . . . . . . . . . . . . . . . . . . 4.1.1 Tensile Properties at Room Temperature . . . 4.1.2 Tensile Properties at Subzero Temperatures . 4.1.3 Tensile Properties at Elevated Temperatures. 4.1.4 Creep Properties . . . . . . . . . . . . . . . . . . . . 4.1.5 Hardness . . . . . . . . . . . . . . . . . . . . . . . . . 4.1.6 Notch Impact Toughness . . . . . . . . . . . . . . 4.1.7 Rotating Bending Fatigue. . . . . . . . . . . . . . 4.2 Wear Resistance . . . . . . . . . . . . . . . . . . . . . . . . . 4.2.1 Abrasive Wear . . . . . . . . . . . . . . . . . . . . . 4.2.2 Impact Wear. . . . . . . . . . . . . . . . . . . . . . . 4.2.3 Wear by Cavitation . . . . . . . . . . . . . . . . . . 4.3 Corrosion Resistance . . . . . . . . . . . . . . . . . . . . . . 4.3.1 Submersion Tests . . . . . . . . . . . . . . . . . . . 4.3.2 Current Density/Potential Tests. . . . . . . . . . 4.3.3 Tests on Intercrystalline Corrosion . . . . . . . 4.4 Magnetic Properties . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . . . . .

85 85 85 88 88 91 93 93 94 95 96 96 99 99 100 102 106 106 108

5

Manufacture . . . . . . . . . . . . . . . . 5.1 Melting and Casting. . . . . . . . 5.1.1 Ingots . . . . . . . . . . . . 5.1.2 Centrifugal Castings . . 5.1.3 Sand Castings . . . . . . . 5.1.4 Refractories . . . . . . . . 5.2 Hot Working. . . . . . . . . . . . . 5.3 Heat Treatment . . . . . . . . . . . 5.3.1 Solution Annealing . . . 5.3.2 Interrupted Quenching . 5.3.3 Continuous Quenching. 5.3.4 Aging . . . . . . . . . . . . 5.4 Cold Drawing . . . . . . . . . . . . 5.5 Welding . . . . . . . . . . . . . . . . 5.6 Machining . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . .

. . . . . . . . . . . . . . . .

111 112 112 113 113 114 115 116 116 118 120 123 123 124 125 126

6

Assessment . . . . . . . . . . . . . . . . 6.1 From Structure to Properties . 6.1.1 Mechanical Properties 6.1.2 Wear Behaviour . . . . 6.1.3 Corrosion Resistance . 6.1.4 Nonmagnetic State . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

. . . . . .

127 127 128 130 130 131

. . . . . .

Contents

6.2

From Manufacture to Application . . . . . . . . . 6.2.1 Constitution and Hot Manufacture. . . . 6.2.2 Workhardening and Cold Manufacture 6.2.3 Application . . . . . . . . . . . . . . . . . . . 6.3 Pros and Cons of HIS . . . . . . . . . . . . . . . . . 6.3.1 Pros. . . . . . . . . . . . . . . . . . . . . . . . . 6.3.2 Cons . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

ix

. . . . . . . .

. . . . . . . .

. . . . . . . .

. . . . . . . .

. . . . . . . .

. . . . . . . .

. . . . . . . .

. . . . . . . .

. . . . . . . .

. . . . . . . .

. . . . . . . .

. . . . . . . .

. . . . . . . .

131 132 133 134 137 137 138 138

Appendix A: Tables. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

141

Appendix B: Figures . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

151

About the Authors. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

163

Index . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

165

Chapter 1

Introduction

Standard stainless austenitic CrNi (Mo) steels are used in a wide range of applications because of their high corrosion resistance and ductility. Their low interstitial content enhances weldability but lowers the yield strength. In contrast to sheet material, castings and forgings do not as much rely on weldability but would profit from more strength. In these cases solid solution strengthening by interstitial atoms of carbon and nitrogen is a promising way of raising strength without losing too much ductility. However, the solubility of interstitial elements depends on the concentration of substitutional elements as e.g. chromium, nickel, manganese and molybdenum. Of these, Cr and Mo are required for corrosion resistance, Ni and Mn for austenite stability. Therefore a proper balance of substitutional elements has to be established to enhance the concentration of interstitials and thus the strength of austenitic steels. In principle one may use C or N or C ? N but C alone did not meet the requirements.

1.1 High Nitrogen Steels Chromium reduces the activity of carbon in steel thus increasing its solubility. However, the maximum solubility at the border to carbide precipitation, i.e. the range of homogeneous austenite, is reduced. Nickel enhances the activity of carbon and lowers its solubility in the lattice. The effect of chromium and nickel on interstitial nitrogen is similar to that on interstitial carbon, except for a higher solubility of the former. In contrast, volatile nitrogen escapes from the melt as N2 gas. This situation is expressed by a solubility of N [ C in austenite but N \\ C in the melt and comprises the basic dilemma of stainless high nitrogen steels (HNS) [1]: Nitrogen offers more solubility in austenite and therefore more strength but requires special manufacturing routes to introduce high contents to the steel. Of these melting and solidification under N2 pressure (pressure metallurgy) and H. Berns et al., High Interstitial Stainless Austenitic Steels, Engineering Materials, DOI: 10.1007/978-3-642-33701-7_1, Ó Springer-Verlag Berlin Heidelberg 2013

1

2

1 Introduction

solid state nitriding of steel powder followed by hot compaction (powder metallurgy) have been industrially exploited. But compared to standard ingot metallurgy at normal pressure of air, pressure and powder metallurgy considerably increase the costs. Instead of special manufacturing processes suitable alloying may be used to enhance the nitrogen solubility of the melt. It aims at an exchange of nickel by manganese, because the latter raises the nitrogen solubility in contrast to the former. About twice as much manganese as nickel is required to fend off d-ferrite which is expressed e.g. by the Schaeffler diagram [2]. Steel Cr18Mn18N0.55 (in mass %) is e.g. used for retaining rings on electric generator shafts, steel Mn23Cr21Ni2N0.85 for drill collars. The 0.2 % proof strength of these two HNS grades, molten at normal pressure of air, was raised to about 400 and 600 MPa, respectively, which amounts to two or three times the level encountered in standard steel Cr18Ni10. The increase from 0.55 to 0.85 mass % N in these two HNS is brought about by a higher content of the substitutional elements chromium and manganese which spur the nitrogen solubility. However, the fracture elongation in tensile tests is lowered and the ductile to brittle transition temperature (DBTT) in notch impact tests is raised pointing to some embrittlement.

1.2 High Interstitial Steels To avoid this deficiency it was proposed to saturate a lean austenitic CrMn steel with interstitial nitrogen at normal pressure and add interstitial carbon to further raise the strength. This concept of high interstitial steels (HIS) with C ? N was presented in 2002 [3] and steel Mn17Cr15N0.43C0.39 showed the following mechanical properties: proof strength Rp0.2 = 494 MPa, true fracture strength R = 2635 MPa, elongation A = 78 % [4, 5]. This remarkable combination of strength and ductility was explained by an increase in the concentration ne of free electrons in austenite via combined alloying with C ? N, compared to alloying with C or N alone. A high ne enhances the ductile metallic character of interatomic bonding. From previous work on the austenite of martensitic stainless steel it was already known [1, 6] that—in the order of alloying with C, N or C ? N—short range atomic ordering is promoted which stabilizes the austenitic phase and raises the interstitial solubility. The significance of the C/N ratio for the precipitation of M2N nitrides or M23C6 carbides and for welding of austenitic CrMn steels was mentioned in [3]. The stepwise addition of up to 0.4 mass % C to steels with (mass %) 19Mn, 17Cr, 3Ni and 0.4–0.55 N raised strength and ductility but hardly the DBTT [7–9]. The corrosion resistance was improved as well, in that the passivation and repassivation was enhanced and the resistance to pitting corrosion strengthened. These beneficial effects assigned to carbon may have been based on C ? N,

1.2 High Interstitial Steels

3

though. It was pointed out that the temperature of beginning precipitation increased with the carbon content leading to embrittlement at lower temperatures. The strengthening part of this precipitation was used in steels for exhaust valves, as e.g. Cr21Mn9Ni4C0.53N0.42 which came into use already in 1952 [10]. They are probably the earliest HIS, but were not meant to dwell on homogenous austenite. A most recent development started from high manganese TWIP steels which rely on twinning induced plasticity [11]. Chromium was added to provide a moderate corrosion resistance and C ? N to enhance the strength. Steel Mn25Cr12C0.32N0.45 for instance arrived at Rp 0.2 = 443 MPa, R = 1635 MPa and A = 99.7 % [12, 13]. This unique combination of strength and high ductility resulted in a specific fracture energy Ws = 751 J/cm3 which is probably the highest ever measured at room temperature [14]. At a stacking fault energy (SFE) of 31 mJ/cm3 twinning is the major strengthening mechanism. The carbide/nitride precipitation was studied in great detail by experiment and simulation as well as the implications for welding [13].

1.3 Aim All development starts from the respective state of art. For high strength austenitic steels it was e.g. compiled in the proceedings of several HNS conferences and a book on HNS [1]. In addition the previous work on HIS of moderate interstitial content was taken into account. Starting from this basis our aim is to explore the potential of HIS and develop advanced grades centering on the following objectives: (a) Strive for a really high interstitial content of C ? N, e.g. between 0.8 and 1.1 mass %, to boost the strength of homogeneous austenite. (b) Balance the effects of C ? N on strength, embrittling precipitation and corrosion, paying special attention to the C/N ratio. (c) Rely on melting and solidification at normal pressure of air to avoid costly pressure or powder metallurgy. (d) Deal with manufacturing on an industrial scale, e.g. with melting, casting, forging, heat treatment. (e) Select HIS for special applications. In all, it is the key aim of this book to come up with manufacturable and applicable new steels with a unique combination of properties, as strength, toughness, wear and corrosion resistance as well as non-magnetisability.

4

1 Introduction

Fig. 1.1 Multiscale approach of investigating high-interstitial austenitic steels from the electron structure to components. Examples are: ab initio calculations, CESR = electron spin resonance, Mössbauer spectroscopy, TEM = transmission electron microscopy, LOM = light optical microscopy, EDX = energy dispersive X-ray analysis and EBSD = electron backscattered diffraction in a scanning electron microscope (SEM), tensile tests

1.4 Procedure To reach this engineering goal the competence of metal physics and materials technology is combined. The development starts with a survey on the constitution of the Fe–Cr–Mn–C– N alloy system to come up with suitable HIS grades. This task is carried out by thermodynamic simulations which not only give the type, amount and composition of phases in dependence of the alloy composition, temperature and nitrogen pressure, but also allow to predict process parameters for melting, casting, forging and heat treatment. Next, the structure of selected HIS grades is studied in depth, starting from the electron structure of austenite and the atomic nano-scale distribution of the elements involved. Transmission electron microscopy (TEM) is e.g. employed to investigate structural changes provoked by mechanical loading. Scanning electron microscopy (SEM) and light optical microscopy (LOM) help to analyse e.g. the progress of fracture, wear and corrosion. High carbon, high nitrogen or low interstitial austenitic reference steels are included for comparison. The attempt is made to describe the structure of HIS from the electron scale up to the scale of components (Fig. 1.1). Key properties of selected HIS grades are measured and traced back to structural features of austenite as e.g. concentration of free electrons, short range atomic order, stacking fault energy, transformation or twinning induced plasticity (TRIP and TWIP), grain boundary precipitation. This knowledge on properties and structural background is to provide a sound basis for applications.

References

5

References 1. Gavriljuk VG, Berns H (1999) High Nitrogen steel. Springer, Berlin 2. Schaeffler AL (1949) Constitutional diagram for stainless steel weld metal. Metal progress 56:680–680B 3. Shanina BD, Gavriljuk VG, Berns H, Schmalt F (2002) Concept of a new high-strength austenitic stainless steel. Steel Res 73:105–113 4. Schmalt F, Berns H, Gavriljuk VG (2004) Mechanical properties of a stainless austenitic CrMnCN steel, Steel Grips 2. Suppl High Nitrogen Steels 2004:437–446 5. Schmalt F (2004) Nutzung der Löslichkeit von C ? N in nichtrostenden Stählen, doctoral thesis Ruhr University Bochum, see also Fortschr. Ber. (2005) VDI 5-702, VDI Verlag, Düsseldorf 6. Gavriljuk VG, Berns H (1999) Precipitates in tempered stainless martensitic steels alloyed with nitrogen, carbon or both, Trans Tech Publ. Zürich Mat Sci Forum 318–320:71–80 7. Bernauer J, Speidel MO (2003) Effects of carbon in high-nitrogen corrosion-resistant austenitic steels. In: Speidel MO, Kowanda C, Diener M (eds) Proceeding HNS 2003, vdf Hochschulverlag AG, ETH Zürich, pp 159–168 8. Bernauer J, Saller G, Speidel MOl (2004) Combined influence of carbon and nitrogen on the mechanical and corrosion properties of Cr-Mn steel grades, Steel Grips 2, Suppl High Nitrogen Steels, pp 529–537 9. Bernauer J (2004) Einfluss von Kohlenstoff als Legierungselement in stickstofflegierten Chrom—Mangan Stählen, doctoral thesis ETH Zürich, No. 15457 10. Müller R, Weintz R (1998) Ventilwerkstoffe für Verbrennungsmotoren, Materialwiss. u. Werkstofftechn 29:97–130 11. Bonaziz O, Allain S, Scott CD, Cugy P, Barbier D (2011) High manganese austenitic twinning induced plasticity steel: a review of the microstructure properties relationships, Current Opinions in Sol. State Mat Sci 15:141–168 12. Mujica Roncery L, Weber S, Theisen W (2010) Development of Mn-Cr-(C-N) corrosion resistant twinning induced plasticity steel: Thermodynamic and diffusion calculations, production, and characterization. Metall Mat Trans 41A(10):2471–2479 13. Mujica Roncery L (2010) Development of high-strength corrosion-resistant austenitic TWIP Steels with C ? N, doctoral thesis, Ruhr University Bochum 14. Berns H, Gavriljuk VG (2007) Steel of highest fracture energy, Key engineering materials, vols 345–346. Trans Tech Publications, pp 421–424

Chapter 2

Constitution

2.1 General Remarks The constitution of high interstitial steels describes the state of atomic order in thermodynamic equilibrium. It depends on the three variables of state: concentration of alloying elements in iron, temperature and pressure. A region that is in the same state of order is known as a phase. Nitrogen is a volatile element and therefore the gas phase and especially the partial pressure of N2 has to be taken into account. The liquid phase appears during melting, solidification and welding. Of the solid state phases austenite, d-ferrite, carbides, nitrides and sigma phase are to be expected. To cope with such a complex alloy system the commercial software program THERMO-CALCTM, version R with TCFE4 database [1] was used to calculate the constitution of multi-component HIS. It depends on experimental data and theoretical models covering a wide range of steel compositions by minimizing the Gibb’s free energy [2]. The program provides phase diagrams as isothermal or isoplethal sections through a system. Also the mole, mass or volume fraction of phases in a given steel may be plotted over the temperature. In addition the chemical composition of each phase is available. In the high temperature range from solidification to solution annealing, which are of practical importance, less deviation from the calculated equilibrium is to be expected than at lower temperatures. However during quenching the kinetics of precipitation are of interest only. It is shown in the next chapter that the atoms in homogeneous austenite are not necessarily distributed evenly, but that microsegregation (mm range) or short range atomic decomposition (clustering, nm range) may cause a chemical inhomogeneity which is not covered by the calculations. Tramp elements are not considered to reduce the computing time except for manufactured grades. Limited experimental verification was in good agreement with the simulation. Therefore the calculated phase diagrams are taken as a reasonable guideline to reveal tendencies. H. Berns et al., High Interstitial Stainless Austenitic Steels, Engineering Materials, DOI: 10.1007/978-3-642-33701-7_2, Ó Springer-Verlag Berlin Heidelberg 2013

7

8

2 Constitution

Fig. 2.1 Isoplethal phase diagram of Fe-18Cr-18Mn in dependence of (a) Carbon content. b Nitrogen content. c C ? N content at constant C/N = 0.6. d C content at constant N = 0.6 mass %, i.e. increasing C/N. Shaded area = homogeneous austenite

2.2 Variation of Interstitial Content To start with, the alloy system Fe-18Cr-18Mn is studied in respect to additions of C, N or C ? N, because the commercial grade Cr18Mn18N0.55 has been successfully manufactured and applied. In fact, nitrogen opens up a wide field of homogeneous austenite (shaded area in Fig. 2.1b) while carbon provides austenite only in combination with carbide M23C6 and/or ferrite (Fig. 2.1a). This is to say that a steel alloyed with 18 mass % Cr to make it stainless and with the same content of Mn to stabilize austenite cannot be strengthened by interstitial carbon without precipitation of carbides which consume chromium and impair toughness. In contrast the high solubility of nitrogen in austenite during solution annealing (Fig. 2.1b) offers intensive strengthening within the target range of 0.8–1.1 mass % N without detrimental precipitation of M2N nitrides. However, the N2 isobars reveal that a partial pressure pN2 = 0.8 bar nitrogen in air is not sufficient

2.2 Variation of Interstitial Content

9

to dissolve 0.8–1.1 mass % N in the melt. At lower contents solidification passes through a regime of ferrite whose solubility for nitrogen is much lower than that of austenite. The resulting degassing may cause boiling or foaming of the solidifying melt or pores in the solid material. Pressure metallurgy is a means of raising the nitrogen content in the melt to a level which assures a fully austenitic solidification. It is important to note that pN2 governs the uptake of nitrogen to the equilibrium content, but that the formation of bubbles at the beginning of degassing depends on the total pressure including the ferrostatic part of the melt which somewhat eases the problem in respect to the calculations. The concept of alloying with C ? N is now applied to the basic composition of steel Cr18Mn18N0.55 to avoid pressure metallurgy (Fig. 2.1c). The C ? N content is plotted along the abscissa at a selected C/N ratio in mass %. The atomic ratio is smaller by [1-(12/14)] 100 & 14 %. The shaded phase field of homogeneous austenite is sufficiently large to dissolve 0.8–1.1 mass % C ? N at a solution anneal temperature TSA of e.g. 1100 °C. As in Fig. 2.1b the L-F-A triple point remains just below 0.9 mass % of interstitials but the corresponding pN2 is lowered from 1.5 to 0.8 bar and a fully austenitic solidification is to be expected (Fig. 2.1c). This means that the partial replacement of N by C is as effective in fending off d-ferrite but comes with less volatility of the remaining nitrogen. The solidus temperature TS is lowered by carbon, but the temperature of beginning precipitation TP is hardly changed within the interstitial range of interest. The type of precipitate turns from M2N to M23C6 though, which is followed by M2N at lower temperatures. The addition of carbon to commercial steel Cr18Mn18N0.55 will reduce the content of d-ferrite during solidification and allow of more nitrogen in the melt. Therefore Fig. 2.1d starts from 0.6 mass % N and directly demonstrates the effect of carbon implying an increase of the C/N ratio. The L-F-A triple point is located at about 0.25 mass % C and the corresponding temperatures TS and TP are about 1320 and 950 °C respectively. At higher carbon contents the range of homogeneous austenite between TS and TP is narrowed and with it the interval for hot working and solution annealing. At the example of the commercial grade Fig. 2.1c and d clearly demonstrate that the C/N concept is suited to increase the interstitial content by melting at normal pressure of air, but that there is a limit to the optimal carbon content.

2.3 Effect of C/N Ratio An isothermal section through the Fe-18Cr-18Mn-C–N system at TSA = 1100 °C outlines the shaded target area of homogeneous austenite (Fig. 2.2). It is encased by ferrite to the left, M2N to the right and M23C6 above. The carbon content of austenite grows from (mass %) 0.3 at 0.2 N to almost 0.8 at 1.1 N. This underlines the beneficial effect of jointly alloying C ? N in respect to the intended increase of the interstitial content in austenite. It may be raised further by a higher TSA,

10

2 Constitution

Fig. 2.2 Isothermal phase diagram at 1100 °C of Fe-18Cr-18Mn-C–N. The dotted lines represent different C/N ratios. Shaded area = homogeneous austenite

Fig. 2.3 Isoplethal phase diagram of Fe-18Cr-18Mn in dependence of the C/N ratio at C ? N = 1 mass %. At (C/N)op the shaded phase field of homogeneous austenite extends to the lowest temperature Top

because the phase field of austenite is expanded at the expense of precipitates, but gives way to ferrite on the left. The dotted lines in Fig. 2.2 represent C/N ratios. At C/N = 1 the respective line cuts the austenitic phase field at the low interstitial end, while at 0.7 it touches the point of highest interstitial solubility. At 0.5 it runs in parallel to the A/A ? M23C6 boarder. The results of this isothermal plot suggest C/N \ 0.7. In an isoplethal section through the Fe-18Cr-18Mn system the C/N ratio is varied by plotting C and N in opposite directions along the abscissa at a constant content of C ? N = 1 mass % (Fig. 2.3). It is evident that the phase field of homogenous austenite extends to the lowest temperature Top at (C/N)op = 0.41 at which M2N and M23C6 start to precipitate simultaneously. The index ‘‘op’’ refers to optimal conditions in respect to

2.3 Effect of C/N Ratio

11

Fig. 2.4 Isothermal phase diagram at 1100 °C of Fe–Cr-Mn-0.3C-0.6 N. Shaded area = phase field of homogeneous austenite with C18 mass % Cr, C ? N = 0.9, C/N = 0.5

retarding the begin of precipitation during quenching from TSA. To the left of (C/N)op M2N starts to precipitate at TP [ Top and to right this holds true for M23C6. In the middle part of the C/N range a fully austenitic solidification prevails.

2.4 Variation of Substitutional Content The basic substitutional elements are chromium and manganese. Molybdenum and copper are of interest in respect to corrosion resistance. Tramp elements can affect the constitution.

2.4.1 Chromium and Manganese The influence of these elements on the constitution is demonstrated by an isothermal section at TSA = 1100 °C and a C ? N content within the target range (Fig. 2.4). The plot suggests a wide phase field of homogeneous austenite of which more than half does not apply, if a minimum content of 18 mass % Cr is chosen to promote corrosion resistance. The lower the manganese content the closer the alloys come to superheated stainless tool steel with retained austenite and little ductility. In addition a partially ferritic solidification and a loss of nitrogen is to be expected. The higher the manganese (and chromium) content, the lower the concentration of free electrons [3]. Therefore the Mn content in Fig. 2.5a is varied only moderately in the range of 18 ± 5 mass %. Unexpectedly ferrite is stabilised by Mn at 1100 °C and M23C6 as well. However, the austenitic phase field grows at the expense of M2N. It shrinks in respect to M23C6 by a higher C/N ratio (Fig. 2.5b) and by a lower temperature (Fig. 2.5c) which may entail a shift of the

12

2 Constitution

Fig. 2.5 Isothermal phase diagrams showing the effect of Cr and C ? N content and of (a) Mn content. b C/N ratio. c Temperature. Shaded phase field = homogeneous austenite, dotted line = C/N = 0.7 at 1050 °C, (+) = target composition

target point (+) at (mass %) 18 Cr and 1 (C ? N) from the A to the A ? M23C6 phase field. As expected, Cr stabilises ferrite and C ? N are required to obtain austenite which loses ground to M23C6 and the more so the higher C/N (Fig. 2.5b). After this view on the constitution at solution anneal temperature TSA in Figs. 2.4 and 2.5 the situation at higher and lower temperatures is of interest. At the example of two steels it is demonstrated that an increase from 13 to 23 mass % Mn lowers the temperature range of solidification but prevents ferrite and raises the temperature of beginning N2 gas evolution (Fig. 2.6). At temperatures below TSA manganese enhances the precipitation of M23C6 at the expense of M2N as already visible in Fig. 2.5a. The promotion of r-phase by Mn is confined to such a low range of temperature that it is likely to be subdued during quenching. The

2.4 Variation of Substitutional Content

13

Fig. 2.6 Phase fraction of two steels with (mass %) 18Cr and 1 (C ? N) at C/N = 0.6 but different Mn content in dependence of temperature, shaded area = homogeneous austenite

Fig. 2.7 Phase fraction of steel Cr18Mn18(C ? N)1 at two different C/N ratios in dependence of temperature, shaded area = homogeneous austenite

jump from 13 to 23 mass % Mn narrows the regime of homogeneous austenite on either side (Fig. 2.6). Thus 18 mass % Mn seem to be a good compromise between high TS, low TP and suppression of ferrite as well as gas. The respective steel Cr18Mn18(C ? N)1 is now analysed as to the influence of the C/N ratio (Fig. 2.7). To stay below (C/N)op (Fig. 2.3) would mean to give away interstitial solubility and strength. A mole fraction of 1 would correspond to C/N = 0.857 which according to Fig. 2.3 would require TSA [ 1100 °C. As for the higher Mn content in Fig. 2.6, the higher C/N ratio narrows the range of homogeneous austenite on both sides (Fig. 2.7). The results suggest not to exceed these limits of the C/N range.

14

2 Constitution

Fig. 2.8 Isoplethal phase diagram of Fe-20Cr-18Mn0.6N–C, shaded area = homogeneous austenite

Fig. 2.9 Isoplethal phase diagram of Fe-18Cr-18Mn0.6N-0.25C-Mo

In Fig. 2.4 chromium contents below 18 mass % were already excluded because of corrosion resistance. Contents above this level are likely to improve this property but promote ferrite and M23C6. This becomes immediately evident if Fig. 2.8 is compared with Fig. 2.1d. The shaded phase field of austenite with 20 mass % Cr is reduced but still allows steels in the upper interstitial target range.

2.4.2 Molybdenum and Copper Alloying stainless steels with molybdenum is a common measure to impede pitting corrosion [4]. As this element is a carbide former and a ferrite stabilizer, the question is what content is permitted in austenitic CrMnCN steels. The influence

2.4 Variation of Substitutional Content

15

Fig. 2.10 Isoplethal phase diagrams of Fe-18Cr-18Mn-2Mo in dependence of a C ? N content at C/N = 0.6. b C/N ratio at C ? N = 1, the pairs of (C/N)op and Top are marked by (+) for different Mo contents Fig. 2.11 Isoplethal phase diagram of Fe-18Cr-18Mn0.6N-0.25C-Cu

of Mo on the constitution is depicted in Fig. 2.9 at the example of steel Cr18Mn18N0.6C0.25. The gas phase is shifted to higher temperatures as Mo lowers the activity of nitrogen in the melt. Ferrite is stabilised to lower temperatures but—up to & 2 mass % Mo—not below &1300 °C which allows hot working in the range of homogeneous austenite. The temperature TP of beginning precipitation (M23C6 followed by M2N) is hardly raised and at 2 mass % Mo stays just below 1000 °C. The temperature of r precipitation is raised, though. The addition of B2 mass % Mo to the above steel appears to be feasible. Starting from the previous example, the influence of C ? N and C/N is investigated next (Fig. 2.10). Compared to Fig. 2.1c the austenitic phase field in

16

2 Constitution

Fig. 2.12 Isoplethal phase diagram of Fe-18Cr-18Mn-2Cu in dependence of a C ? N content at C/N = 0.6. b C/N ratio at C ? N = 1, the pairs of (C/N)op and Top are marked by (+) for different Cu contents

Fig. 2.10a is reduced by ferrite to the left and liquid above. In respect to Fig. 2.3 Top and (C/N)op in Fig. 2.10b are changed only moderately by up to 4 mass % Mo. Stainless steels are alloyed with copper to reduce general corrosion e.g. in nonoxidising acid solution [4]. In contrast to molybdenum, copper stabilises austenite and raises the activity of interstitials. This is reflected in Fig. 2.11 indicating a steep rise of TP. At CuG & 1.8 mass % the evolution of N2 gas ends the range of homogeneous austenite. At this Cu level the range of 0.8–1.1 mass % C ? N leads to an austenitic solidification and a reasonable TSA (Fig. 2.12a). The pairs of (C/N)op and Top at 1 mass % C ? N increase considerably with the Cu content (Fig. 2.12b). While the addition of 2 mass % Mo or Cu seem to be feasible, joint alloying of both elements to this level leads to a dramatic shrinkage of the austenitic phase field which would make it difficult to process such a steel.

2.4.3 Tramp Elements Small quantities of the strong carbide and nitride formers vanadium, niobium and titanium may result in MX precipitates. In view of the high interstitial content of HIS these precipitates hardly dissolve at TSA and would represent an additional phase. Silicon raises the activity of C and N and thereby promotes precipitation. This is reflected for Fe-18Cr-18Mn-1(C ? N) by a shift of the pairs (C/N)op and Top from 0.408 and 971 °C at zero Si to 0.486 and 1026 °C at 0.5 mass % Si and further to

2.4 Variation of Substitutional Content

17

0.560 and 1075 °C at 1 mass % Si. To avoid a detrimental increase of Top by up to 100 °C it is recommend to not fully exploit the range of Si B 1 % given e.g. in EN10088 for austenitic steels, but to keep its content as low as possible.

2.5 Selection of Steels The phase field of homogeneous austenite is the target area of HIS to be reached by solution annealing and preserved by quenching. An austenitic solidification is desirable to transfer nitrogen from the melt to the austenite without degassing. To keep the concentration of free electrons high and with it the ductile metallic character of interatomic bonding a reduction of substitutional alloy content, namely of Cr, Mn, Mo, would be helpful (see Sect. 3.1.3). However, the experience with HIS Mn17Cr15N0.43C0.39 and general knowledge on stainless steel speak for 18 mass % Cr. As shown above, 18 mass % Mn are a reasonable match to avoid ferrite during solidification. The interstitial content is aimed at C0.8 mass % C ? N to boost strength. In view of experience with HNS and the calculations above the envisaged upper limit of 1.1 mass % C ? N is confirmed to avoid problems during hot working and heat treatment. As to the C/N ratio, one has to start from the soluble content of nitrogen in alloys with 18 mass % of Cr and Mn each at normal pressure of air. It slightly depends on the carbon content, but 0.6 mass % N seems to be a fair value to start with. Combining the C ? N content and the C/N ratio in mass % we arrive at C ? N = N [(C/N) ? 1]. Inserting N = 0.6 and (C/N)op = 0.41 the C ? N content is 0.85. Higher interstitial contents have to rely on more carbon which entails C/N [ (C/N)op and TP [ Top. Based on these constitutional considerations three steels with 18 mass % of Cr and Mn each were selected: one at the lower end of the C ? N target range, one at the upper end and one in the middle. These three new HIS were molten and hot worked on an industrial scale and designated according to their C ? N content, times 100, i.e. CN85, CN96 and CN107 (Table 2.1). To these Mo and Cu were added by remelting a smaller batch. A series of HIS with 0.65–1.15 mass % C ? N was produced as castings. In accordance with European standards the designation is preceeded by ‘‘G’’. A few reference steels are listed of which CrNi represents a standard low interstitial grade, MnC a high carbon Hadfield steel and CrMnN a high nitrogen steel. The grades MnCr82 and MnCr70 are reference HIS of lower chromium and interstitial content. For the investigation of structure (Chap. 3) and properties (Chap. 4) premachined specimens were solution annealed, quenched and machined to final size. Details of HIS manufacture are given in Chap. 5.

0.001 0.002 \0.001 \0.001 \0.001 \0.001 0.002 0.001 \ 0.001 \0.001 \0.001 0.022 0.014 – 0.007 –

18.3 18.2 18.8 17.9 18.3 17.6 19.9 20.2 20.1 19.9 18.3 18.7 0.2 21.0 14.7 12.0

18.5 18.9 18.9 19.0 18.6 19.3 18.0 18.0 18.1 18.0 18.4 1.9 12.1 23.1 17.2 25.4

Mo 0.04 0.06 0.07 0.96 0.94 0.04 0.04 0.04 0.04 0.06 0.02 – – 0.2 \0.02 –

Cu – –(a) –(a) – – 1.90 0.398 0.316 0.306 0.268 –(a) – – – – –

Ni 0.26 0.34 0.40 0.32 0.38 0.28 0.47 0.39 0.40 0.46 0.10 9.04 0.1 1.5 \0.07 –

C?N 0.85 0.96 1.07 0.94 1.03 0.96 0.65 0.88 0.98 1.15 0.85 0.05 1.20 0.92 0.82 0.77

0.44 0.56 0.85 0.52 0.78 0.62 0.05 0.35 0.69 0.80 0.43 0.08 132 0.05 0.90 0.71

C/N

No. 1–6 = Hot worked steel, 7–10 = Centrifugal castings, 11 = Sand casting, a section of which was hot worked ? CN859, 12–16 = Hot worked reference steels (a) V \ 0.07

0.018 0.021 0.026 0.025 0.027 0.023 0.017 0.017 0.018 0.018 0.017 0.020 0.090 – 0.044 –

Si

0.59 0.614 0.578 0.620 0.582 0.594 0.616 0.654 0.583 0.641 0.596 0.050 0.009 0.880 0.431 0.447

0.26 0.30 0.43 0.23 0.42 0.18 0.33 0.28 0.14 0.33 0.54 0.57 0.49 0.30 0.48 –

0.26 0.344 0.489 0.324 0.452 0.370 0.033 0.228 0.400 0.512 0.256 0.004 1.190 0.040 0.387 0.319

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16

CN85 CN96 CN107 CN94Mo1 CN103Mo1 CN96Cu2 GCN65 GCN88 GCN98 GCN115 GCN85 CrNi MnC CrMnN MnCr82 MnCr77

Table 2.1 Chemical composition in mass % of the austenitic steels investigated No. designation C N P S Cr Mn

18 2 Constitution

References

19

References 1. Software System and Users Guide (2008) Thermo-Calc Software AB SE- 11347 Stockholm 2. Saunders N (1995) Phase diagram calculation for high-temperature structural materials. Phil Trans Royal Soc London A351:543–561 3. Shanina BD, Gavriljuk VG, Konchitz AA, Kolesnik SP (1998) The influence of substitutional atoms upon the electron structure of the iron-based transition metal alloys. J Phys Condensed Matter 10:1825–1838 4. Heimann W, Oppenheim R, Wessling W (1993) Stainless steels. In: Steel, vol 2. Springer, Berlin, pp 382–422

Chapter 3

Structure

It is generally accepted among metal scientists and engineers that ‘‘structure’’ stands for crystal lattice, lattice defects and their distribution as well as for grain size. In the solid solutions, the type of solute atoms and their distribution, as well as precipitates are taken into account. The aim of this chapter is to show that, in fact, the structure of metals and alloys starts from localized or free electrons. Under external force and resultant straining, the atoms are being shifted from their positions, and mechanical response, plastic deformation or brittle fracture, depend on the character of interatomic bonds. In comparison with the nuclei, the response of electrons is quicker by many orders of magnitude. The closed electron shell, so-called ‘‘ion core’’, can be excluded from the consideration because it does not take part in chemical reactions and, under straining, can be only slightly deformed, i.e. polarized. Only the external, i.e. valence electrons are responsible for chemical bonds and deformation behaviour. According to a modern approach, they reveal an ‘‘itinerant’’ behaviour, which means that the same valence electrons are sometimes free and sometimes localized at the atomic sites. However, some permanent part of valence electrons is always free, and the free/localized electron ratio is the controlling factor of the nature of metals, particularly in respect to their behaviour under external action. The prevailing localized valence electrons form covalent bonds between the atoms in the crystal lattice, which causes brittleness because even a slight shift of the atoms under shear stress in the slip plane leads to the breaking of interatomic bonds. This is, e.g., the case of the transition metals and alloys of group V and VI in the periodic table (V, Nb, Ta, Cr, Mo, W). Free electrons are responsible for the metallic character of interatomic bonds, and the higher their fraction is the more ductile are metals and alloys. In relation to phase transformations, valence electrons are responsible for the height of the energy barrier which has to be overcome by the atoms constituting a new crystal lattice, either during its nucleation or during their jumps through the interface between matrix and new phase. Again, the stronger the H. Berns et al., High Interstitial Stainless Austenitic Steels, Engineering Materials, DOI: 10.1007/978-3-642-33701-7_3, Ó Springer-Verlag Berlin Heidelberg 2013

21

22

3 Structure

covalent bonds between the atoms are, the higher this energy barrier is. For this reason, the transition metals of group V and VI do not reveal any polymorphic transformations. Moreover, as will be shown in this chapter, the control of interatomic bonds and the free/localized electron ratio affects short-range atomic order in multicomponent solid solutions and, for this reason, their thermodynamic stability. It is a privilege of metallurgists that iron belongs to metals with a rather high part of free electrons. Alloying it with the elements located to the right of iron in the periodic table (Ni, Co, Cu) increases the concentration of free electrons thereby enhancing the metallic character of interatomic bonds and assisting ductility. Elements to the left of iron (Mn, Cr, Mo, V etc.) act in the opposite direction (see about details [1]). The alloyed steel is a complicated engineering material and the knowledge of a fundamental correlation between the interatomic bonds on the one hand and thermodynamic stability as well as mechanical and chemical properties on the other should be useful for their deliberate and successful design. The current chapter is an attempt to approach this ambitious task.

3.1 As-Quenched The initial microstructure of homogeneous austenite is achieved by solution annealing and quenching in water. At first this structure will be discussed followed by an investigation on the effect of straining presented in Sect. 3.2.

3.1.1 Electron Structure: Calculated and Measured The proposed concept of alloying the austenitic steels with carbon ? nitrogen is based on the ab initio calculations of the electron structure where only the type of the crystal lattice and the charge of the nuclei of constituting atoms are set as starting points. The interatomic distances are obtained from the minimum of the calculated cohesive energy and all further properties are calculated without any assumptions. The construction of the calculated atomic configuration consisting of 32 substitutional and 2 interstitial atoms is shown in Fig. 3.1. The corresponding chemical composition amounts to (atom%) 58.8 Fe, 23.5 Mn, 11.8 Cr and 5.9 of interstitials. This composition was chosen for the calculations as a compromise between the requirements to keep up the permanent translation of calculated configuration in three orthogonal directions and to minimize the total number of calculated atoms in order to reduce the calculation time. The method of full potential linearized augmented plane waves (FLAPW) based on the density functional theory [2, 3] and the program package Wien2k [4] were used for the calculations. Of the electron properties which can be obtained using

3.1 As-Quenched

23

Fig. 3.1 Atomic configuration Fe20Mn8Cr4 with C2, N2 or C1N1 interstitial atoms chosen for ab initio calculations of the electron structure. The chosen symmetry of the configuration allows its translation along three orthogonal directions

Fig. 3.2 Density of electron states in the valence electron band of the calculated solid solutions. The change of the DOS in the vicinity of the Fermi level is shown in the upper insert

these calculations, the density of electron states, DOS, at the Fermi level, EF, is particularly important because it determines the interatomic bonds, namely the concentration of free electrons and the thermodynamic stability of structures. According to Fermi statistics for free electrons, their distribution on the energy scale is so that only electrons at the Fermi level can change their energy, i.e. be really free and contribute to heat capacity, conductivity etc.

24

3 Structure

Figure 3.2 shows how nitrogen, carbon or carbon ? nitrogen affect the density of electron states in the upper energy band of the CrMn austenitic steel. The electron structure of the iron is described as 1s22s22p63s23p63d64s2, where 1–4 are the main quantum numbers which determine the energy of the corresponding electron band, whereas s, p and d characterize the space symmetry of electrons, namely a spherical symmetry for s-electrons and a leaf symmetry for pand d-electrons. The two latter differ in their angular moments. The electron energy states 1s22s22p63s23p6 form a closed shell, the ion core, which remains unchanged at any chemical reactions, whereas 3d64 s2 electrons belong to the valence band and reveal a so-called itinerant behaviour. As mentioned above, sometimes they are free, sometimes localized. However, any time a definite constant part of these electrons are free. The highest valence electron energy corresponds to the Fermi level, of which the width is proportional to kT, where k is the Boltzmann constant. The Fermi level in Fig. 3.2 is located at the co-ordinate origin. All energy states located below the Fermi level are occupied. Only the electrons within the kT interval can change their energy. Just these electrons determine the electron capacity, conductivity and thermodynamic stability of phases. Alloying affects the distribution of valence electrons on their energy states. The addition of new elements does not just shift the Fermi level along the rigid band. Each new chemical composition creates its own valence electron energy band. One can see in Fig. 3.2 that interstitial elements create so-called bound states at the bottom of the valence electron band of the FeCrMn substitutional solid solution, shift the electron levels on the energy scale and change their amplitude. As follows from the insert in the upper left corner of Fig. 3.2, interstitial elements change the population of states at the Fermi level at which electrons can change their energy and, therefore, contribute to a change in the properties of materials. Alloying with carbon decreases the DOS, whereas nitrogen increases it. The partial substitution of carbon by nitrogen leads to a further increase of DOS. The most significant consequence of the effect of interstitial elements on the DOS at the Fermi level presented in Fig. 3.2 is the expected change in the concentration of free electrons, i.e. in the character of interatomic bonds. Judging on the results presented in Fig. 3.2, one can predict that carbon should decrease the concentration of free electrons, whereas nitrogen and, particularly, carbon ? nitrogen should increase it. A qualitative presentation of this effect is given in Figs. 3.3 and 3.4. A striking difference in the distribution of valence (free) electrons in the vicinity of carbon and nitrogen atoms is observed (compare Figs. 3.3a and b), which suggests that nitrogen atoms migrate through the crystal lattice being surrounded by clouds of free electrons, whereas carbon atoms are expected to have a shortage of electrons in their vicinity as compared with free atoms. It is relevant to note that these theoretical results confirm the old experimental observations by Seith et al. of the electrotransfer in austenitic steels with carbon [5] or nitrogen [6] according to which the carbon atoms are positively charged, whereas the nitrogen ones carry a negative electric charge.

3.1 As-Quenched

25

Fig. 3.3 Spatial distribution of the valence electron density el/a.u.3 along the (110) plane in the fcc lattice of solid solutions. a Fe20Mn8Cr4C2. b Fe20Mn8Cr4N2 and c Fe20Mn8Cr4C1N1. 1 a.u. = 0.529 Å

Fig. 3.4 Projection of the spatial distribution of the valence electron density on the (100) plane. The electron density in the interatomic space increases in the sequence of a Fe20Mn8Cr4C2 ? b Fe20Mn8Cr4N2 ? c Fe20Mn8Cr4C1N1. The plane (100) is chosen in order to demonstrate how the alloying with C ? N increases the density of free electrons even in the vicinity of the carbon atoms (compare Figs. a and c). The nitrogen atoms are not present on the (100) plane in the C1N1 composition

The spatial valence electron distribution around C and N atoms in case of alloying with C ? N (Fig. 3.3c) repeats main features of steels with C or N. However, the two-dimensional projections of the spatial valence electron distribution allow to estimate the electron distribution within the interatomic space, and it is clearly seen from Fig. 3.4 that the electron density in the space between

26

3 Structure

Fig. 3.5 A scheme for CESR measurements. a Free electron energy levels split under an applied magnetic field H. b Precession of the electron spin along the field and in opposite direction (the ground and upper energy levels in Fig. 3.5a, respectively). c Absorbed microwave energy P spent for transfer of electrons from the ground to the upper energy level. For convenience, it is presented as a derivative of the applied field, because this visualizes the asymmetry of the signal, a feature of the signal from free electrons caused by their migration for the spin relaxation time

the atomic sites is highest for the steel alloyed with C ? N (Fig 3.4c). And, of course, it is higher in the N steel in comparison with the C steel. This suggests an increase in the concentration of free electrons in austenitic steels with nitrogen or carbon ? nitrogen. It is particularly important that, in case of alloying with carbon ? nitrogen, the spatial distribution of free electrons is more homogeneous and their density increases also in the vicinity of carbon atoms (compare Figs. 3.4a and c). The occurrence of excessive free electrons around the nitrogen atoms allows to interpret some unordinary mechanical properties of austenitic nitrogen steels, namely quasi-cleavage at low temperatures and high strain rates, which will be discussed in Sect. 3.2.3.1. Experimental studies confirm the results of theoretical calculations. The concentration of free electrons was measured by conduction electron spin resonance, CESR, [7–9], the essence of which is shown in Fig. 3.5.

3.1 As-Quenched Table 3.1 Effect of C, N and C ? N (mass %) on the concentration of free electrons ne in steel Cr13Mn18

Table 3.2 Concentration of free electrons ne in new HIS

27 Composition

ne, 1022

Cr13Mn18N0.7 Cr13Mn18C0.25N0.25 Cr13Mn18C0.4N0.4

1.1 2.9 3.7

-3 cm

Steel

ne, 1022 cm-3

CN107 CN96 CN85

2.0 2.9 2.8

Under an applied magnetic field (Fig. 3.5a), the degenerated energy level of free electrons with magnetic quantum number ‘ is split into two energy states with spin -‘ (spin orientation along the applied field) and +‘ (spin orientation in the opposite direction). The microwave energy spent for a transfer of electrons from the ground to the upper energy level, i.e. for overturning the spin into the direction along the field (Fig. 3.5b), is proportional to the concentration of free electrons, which is calculated using a reference sample with a known spin concentration. Three CESR signals in Fig. 3.5c belong to the austenitic nitrogen steel Cr13Mn18N0.7 and two steels of the same basic composition where nitrogen is partly replaced by carbon. Thus, the combined alloying with nitrogen ? carbon enhances the signal from free electrons, i.e. increases the concentration of free electrons in the austenitic steel as predicted by theoretical calculations. The quantitative data of free electron concentration are presented in Table 3.1. Based on the described results, a concept of alloying steels with carbon ? nitrogen was developed in [10–18]. Three austenitic steels with (mass %) 18Cr, 18Mn and (i) 0.489C ? 0.578N (CN107), (ii) 0.344C ? 0.614N (CN96) and (iii) 0.26C ? 0.59N (CN85), see Table 2.1, were chosen for detailed studies. The obtained data of the free electron concentration are presented in Table 3.2. The increase of the chromium content from steel Cr13Mn18C0.4N0.4 to steel CN85 needed for the improvement of corrosion properties, leads to some decrease in the concentration of free electrons. Nevertheless, it remarkably exceeds that of the nitrogen steel and, of course, of austenitic carbon steels. Finally, the measured concentration of free electrons in austenitic steels with carbon, nitrogen and carbon ? nitrogen is presented in Fig. 3.6. One can see that carbon does not really change the concentration of free electrons in austenitic steels. In other words, carbon supplies its valence electrons to the energy levels below the Fermi level. Alloying with nitrogen increases it up to some critical value. One can conclude that some optimum should exist for the nitrogen content in austenitic steels. The arrows in Fig. 3.6 show how the concentration of free

28

3 Structure

Fig. 3.6 Effect of carbon, nitrogen and carbon ? nitrogen content on the concentration of free electrons in austenitic steels

electrons increases if, at the same content of substitutional alloying elements, a part of nitrogen is replaced by carbon. The combined alloying with C ? N remarkably increases the concentration of free electrons and shifts its maximum towards higher interstitial contents, which, in comparison with austenitic nitrogen steels, can be used to increase the content of interstitials without pressure metallurgy (see Sect. 5.1).

3.1.2 Atomic Distribution So-called ideal solid solutions, where the atoms do not interact with each other and the gain in the Gibbs free energy is obtained only due to the increased entropy, do not really exist in nature. A different electron constitution of the host and solute metal atoms predetermines different atomic interactions. Short-range atomic order designates any deviation from the statistical atomic distribution. The prevailing bonds of the host metal atoms with the solute atoms, M–S, are characterized by the term ‘‘short-range atomic ordering’’, whereas the favour for M–M and S–S bonds means ‘‘short-range atomic decomposition’’. The steel is a multicomponent solid solution, and the segregation of alloying elements, well known to metallurgists, is the utmost result of short-range decomposition. High temperature treatments are used in order to reach the more or less homogeneous state of steels and, in such a way, provide their stability to precipitation phenomena in the course of technological operations. Nevertheless, clusters, i.e. accumulations of one kind of atoms, exist even in the liquid alloys,

3.1 As-Quenched

29

which is well known from the studies of rapid quenching from the liquid state (see, e.g. [19] for a FeSiB alloy). This is why, except for some special cases, the aim of steel designers is to enhance the tendency to short-range atomic ordering, in order to reach the highest possible thermodynamic stability of phases. Because of the interaction of solute atoms with dislocations and a tendency to precipitation, the effect of atomic distribution on mechanical and corrosion properties is not less important. Among the available experimental technique, Mössbauer spectroscopy is suitable for studies of short-range atomic order in solid solutions because the change in the nearest atomic neighbourhood of iron atoms strongly affects the main parameters of spectra, namely, the hyperfine field at the atomic nuclei, which causes the splitting of a single line of paramagnetic austenite into six lines of ferromagnetic martensite proportional to the atomic magnetic moment (so-called Zeemann splitting), the isomer shift (the shift of the spectrum gravity center caused by a change of the electron density at the nuclei) and the quadrupol interaction (e.g. the splitting of the single line of paramagnetic austenite into the doublet, if the nearest solute atoms cause a local deviation of the crystal lattice from its cubic symmetry). For the paramagnetic solid solutions like austenitic steels, the Mössbauer spectrum consists of a single line and doublets of which the quadrupol splitting is proportional to the local distortions of the crystal lattice. The study of interstitial solid solutions is most informative because the interstitial atoms cause higher distortions in comparison with the substitutional ones. For this reason, we start with the binary Fe–C and Fe–N austenites. Thereafter, the ternary Fe–C–N solid solution will be analyzed, which allows to identify important features of the atomic distribution due to combined alloying with C ? N as a consequence of a change in the electron structure. As the atomic distribution in the austenite is inherited by the martensitic phase because of the diffusionless martensitic transformation, the Zeemann sextet of the martensite was used for characterization of the atomic distribution in austenitic steels. This approach is particularly informative because the interstitial as well as substitutional alloying elements decrease the hyperfine field at the iron nuclei (i.e. the extent of Zeemann splitting) and this decrease is proportional to the number of solute atoms as nearest neighbours of the iron ones. Mössbauer spectra of binary Fe–C and Fe–N solid solutions are presented in Fig. 3.7 (see also [20, 21]). The samples were prepared using the saturation of pure iron with carbon or nitrogen in CH4 ? H2 or NH3 ? H2 mixtures at 1150 and 700 °C, respectively. The carbon and nitrogen concentration were determined by X-ray diffraction. The atomic configurations corresponding to the components in the spectra are shown to the left. The Fe–C spectrum consists of a single line belonging to iron atoms Fe0 having no carbon atoms as nearest neighbours. The doublet comes from iron atoms Fe1 with one carbon atom as nearest neighbour and Fe290 with two carbon atoms in neighbouring nearest interstitial sites. These two different configurations cause the

30

3 Structure

Fig. 3.7 Mössbauer spectra of binary austenitic solid solutions, atom %. a Fe-9.1 C. b Fe-9.3 N and c–f corresponding Fe–C(N) atomic configurations

same electric field gradient in the crystal lattice (only its sign is different), which is displayed by the same quadrupol splitting of the spectrum. Such 90°-pairs are never met in Fe–N austenite. Instead, in addition to the Fe1 component, the spectrum contains a doublet from the Fe2180 configuration caused by nitrogen atoms occupying interstitial sites within the second coordination sphere. Such a dumbbell-like configuration is an element of the ordered Fe4N c0 -phase. In order to obtain the values of C–C and N–N interaction energies consistent with the fractions of atomic configurations derived from Mössbauer spectra, a modelling of these solid solutions was carried out using the Monte Carlo method (Fig. 3.8). W1 and W2 are the energies of interaction between any two interstitial atoms in the first and second coordination spheres, respectively, if one of them is located in the coordinate origin. The areas marked as C–C and N–N correspond to the values of C–C and N–N interactions which are consistent with the fractions Fe1 (d), Fe290 (e) and Fe2180 (f) atoms in Fig. 3.7. It is seen that the carbon distribution in austenitic steels is characterized by a soft repulsion between C atoms in nearest interstitial sites (a small W1 for the C–C

3.1 As-Quenched

31

Fig. 3.8 Areas of C–C and N–N atomic interactions (marked with gray colour) within the first and second coordination spheres in the sublattice of interstitial sites corresponding to the fractions of Fe1 (d) Fe2–90 (e) and Fe2-180 (f) iron atoms obtained from Mössbauer spectra in Fig. 3.7

area), so that, along with single carbon atoms, some fraction of carbon pairs Fe2–90° should exist. However, a hard C–C repulsion (large W2) is revealed for carbon atoms as neighbours in the second coordination sphere of the interstitial sublattice. This means that dumbbell-like C–Fe–C configurations Fe2–180° are not met in the austenitic carbon steels, which makes the existence of an ordered Fe4C type structure impossible. In contrast, the N–N repulsion in the first coordination sphere is so hard (large W1) that nitrogen atoms cannot be nearest neighbours in the austenitic lattice, whereas the soft repulsion in the second coordination sphere (small W2) allows N–Fe–N pairs which are clearly identified in the Mössbauer spectra of austenitic nitrogen steels. This is why the ordered c0 -phase Fe4N exists, and, based on the quoted studies [20, 21], one can state that carbon atoms in austenitic steels are prone to form clusters, whereas the distribution of nitrogen atoms is characterized by short-range atomic ordering. For studies of atomic distribution in the solid solutions alloyed with C ? N, a ternary alloy of iron with mass % 0.93C and 0.91N was used [22]. The samples were prepared by Dr. Rawers, Albany Research Center, USA, and obtained by a special technique of melting an iron-carbon steel in a hot isostatic pressing (HIP) furnace under a nitrogen gas pressure of 160 MPa (see [23]). According to the Xray diffraction measurements, the alloy consisted of 60 austenite and 40 % martensite. The C ? N content was 1.57 mass in martensite and 2.40 mass % in austenite, estimated from crystal lattice dilatation. A Mössbauer spectrum of this alloy is presented in Fig. 3.9. The sextets represent the ferromagnetic martensite, whereas the central part of a single line and a doublet come from the paramagnetic austenite. Two sextets Fea0 and Fea1 belong to atomic configurations of iron atoms with no or one interstitial atom as nearest neighbour, respectively. The following features distinguish this spectrum in comparison with those of Fe–C and Fe–N austenite and martensite. The austenitic part of the spectrum contains a single line Fec0 of the iron atoms having no interstitials in the nearest neighbourhood and the doublet Fec1 of the iron atoms with one interstitial atom as nearest neighbour. In comparison with the spectrum of Fe–N austenite (see

32

3 Structure

Fig. 3.9 Mössbauer spectrum of iron with (mass %) 0.93C and 0.91N Table 3.3 Abundance of different atomic configurations in Fe-0.93C–0.91N alloy, according to Mössbauer data Phase (%) Austenite Martensite Configuration Spectrum mode

Fec0 single line

Fec1 ? Fec2-90° doublet

Fea0 sextet 1

Fea1 sextet 2

Abundance in spectrum Normalized abundance

20 53

18 47

47 86

8 14

Fea4 sextet 3 (clusters) 7 –

Figs. 3.7, 3.8), the component from the N–Fe–N dumbbell-like configurations (Fec2180 atoms) is not present. In comparison with the spectrum of Fe–C martensite (see e.g. [24]), the component Fea290 with two carbon atoms as nearest neighbours is also absent. The occurrence of the low intensive sextet Fea4 suggests that some clusters of interstitial atoms exist in the martensite, which can be a consequence of the rather slow cooling of the 5 kg ingot. The obtained results evidence a distribution of interstitial atoms in the studied Fe–C–N solid solution that is characterized by a tendency to the hard repulsion within both the first and second coordination spheres in the sublattice of interstitial sites. The abundances of different atomic sites in relation to the whole alloy are presented in Table 3.3 and are normalized to 100 % of the austenite or martensite. In order to clarify the distribution of interstitial atoms responsible for the observed abundances of different configurations of iron atoms and interstitial ones, a Monte Carlo computer simulation was used. The detailed procedure is described in [20, 21]. Results obtained for the austenitic and martensitic phases are presented in Fig. 3.10.

3.1 As-Quenched

33

Fig. 3.10 The areas of i–i interaction in two coordination spheres satisfying Mössbauer data of alloy Fe-0.93C-0.91N. a Austenitic phase, b martensitic phase

For the austenitic phase, the abundances of configurations Fec1 obtained from Mössbauer studies can be reached only in a narrow range of concentrations, i.e. Ni/NFe & 0.085, which corresponds to 1.9572.21 mass %. Ni and NFe are the number of interstitial and iron atoms involved. It can be seen that in case of Ni/NFe & 0.098, the data look like the Fe–C austenite, and Ni/NFe & 0.09 leads to the combination of the interaction energy profiles found for C–C and N–N interactions in binary Fe–C and Fe–N austenites, respectively (see Fig. 3.8). At smaller concentrations of interstitials only a strong repulsion in two coordination spheres can lead to the agreement with Mössbauer data, and apparently it is this situation that prevails in the studied alloy. As a result of the simulation made for the martensitic phase, only the concentration ratio Ni/NFe & 0.064 provides the agreement with Mössbauer data of repulsion between interstitial atoms within the two coordination spheres. At the higher contents of interstitials, an attraction within the second coordination sphere should occur, which is, however, not observed in the Mössbauer spectrum, possibly because of some clustering in martensite resulting from too slow cooling of the 5 kg ingot after melting and nitrogen saturation (see sextet 3, correspondingly atomic configuration Fea4 in Fig. 3.9 and Table 3.3). Summing up, one can state that the values of interaction energies between interstitial atoms obtained from the Monte Carlo simulation of the Fe–C-N solid solution correspond to such a distribution of carbon and nitrogen atoms that is not prone to form clusters within the first two coordination spheres. This result is quite different from the distribution of carbon atoms in binary Fe–C austenite and martensite where interstitials can occupy neighbouring interstitial sites (see Fig. 3.7e for the austenite) and the effective C–C repulsion exists in the second coordination sphere only. Compared to the data of binary Fe–N alloys where nitrogen atoms can occupy neighbouring sites in the second coordination sphere, these 1808 N–N configurations (Fig. 3.7f) are not revealed in the ternary Fe–C–N alloy.

34

3 Structure

Fig. 3.11 Mössbauer spectra of steel Cr15Mo1 alloyed with (mass %) a 0.6C, b 0.62N and c 0.29C ? 0.35N after solution treatment at 1100°C followed by quenching in water

Such a strong i–i repulsion within the first two coordination spheres should provide a higher thermodynamic stability of the iron-based Fe–C-N solid solutions, which is consistent with the results of the calculated phase equilibrium (see Chap. 2, Constitution). Experimental evidence of an increased stability of ironbased C ? N solutions to phase transformation was obtained by Mössbauer spectroscopy. The steel Cr15Mo1 alloyed with (mass %) 0.6C, 0.62N or 0.35C ? 0.29N was used in order to compare the stability to martensitic transformation. Mössbauer spectra are shown in Fig. 3.11. The single component belongs to austenite and the sextet to martensite. A striking difference between fractions of the retained austenite in these steels after quenching demonstrates the non-additivity of carbon and nitrogen effects in austenitic steels. The physical nature of this phenomenon lies in the correlation between the electron structure, short-range atomic order and thermodynamic stability of solid solutions. Like binary Fe–C, Fe–N and ternary Fe–C–N austenites, the alloying of austenitic steels with nitrogen and particularly with carbon ? nitrogen increases the concentration of free electrons, which promotes short-range atomic ordering and, as consequence, increases the thermodynamic stability of solid solutions. It is remarkable that the same effect is observed in the tempered martensite. Dilatometric curves of tempering are presented in Fig. 3.12 for all three steels. The first, second and third transformation during tempering, which correspond to the

3.1 As-Quenched

35

Fig. 3.12 Dilatometry of asquenched Cr15Mo1C0.6, Cr15Mo1N0.62 and Cr15Mo1C0.29N0.35 martensites during tempering at a heating rate of 1.5 K/min

precipitation of an intermediate e-carbide, the decomposition of the retained austenite and the transformation e ? h to cementite, are clearly identified in carbon martensite. Because of the increased stability of the retained austenite in the alloyed steel, the second transformation is shifted to temperatures above that of cementite precipitation. The precipitation of M7C3 carbide occurs at temperatures above that of austenite decomposition. The dilatometric curve of the nitrogen steel is rather smooth, and the effect of nitride precipitation is shifted to higher temperatures extending over a broad temperature range. The retained austenite is not decomposed during heating but is transformed into martensite in the course of subsequent cooling. No tempering effects can be identified in C ? N martensite by dilatometry. As follows from Mössbauer studies, this difference in the tempering behaviour is caused by different atomic distributions (Fig. 3.13). Tempering at 650 °C removes the interstitial elements from the solid solution. Therefore, the spectra in Fig. 3.13 are controlled mainly by the distribution of chromium atoms and, to an insignificant extent, molybdenum atoms. The component Fe0 belongs to the iron atoms with no substitutional atom as nearest neighbour and it decreases if carbon is substituted by nitrogen or nitrogen by carbon ? nitrogen. This is clear evidence that the chromium distribution is more homogeneous in the C ? N martensite and it remains so even after tempering at rather high temperatures.

36

3 Structure

Fig. 3.13 Lines V6 (nuclear transition +1/2 ? +3/2) in Mössbauer spectra of steel Cr15Mo1 alloyed with carbon, nitrogen or carbon ? nitrogen after quenching from 1100 and tempering at 650 °C for 2 h

3.1.3 Chemical Nanoscale Homogeneity In the previous section, Mössbauer spectroscopy was employed to analyse the atomic distribution in alloyed steels and to detect changes from short-range ordering to short-range decomposition, i.e. clustering. Such changes come with a reduction of the chemical homogeneity, i.e. a higher concentration of alloying atoms in clusters. It is the aim of this section to investigate their size which requires methods different from Mössbauer spectroscopy. It will be shown that alloying with C ? N improves the chemical homogeneity on a nanoscale (see also [25]). Three steels CN85, CN96 and CN107 together with the Hadfield steel MnC were chosen for comparative studies (see their composition in Table 2.1). The following experimental techniques were used: conduction electron spin resonance (CESR), magnetic measurements and transmission electron microscopy (TEM) of stacking fault energy (SFE). As shown in Sect. 3.1.1, the CESR allows to measure the concentration of free electrons in the paramagnetic objects. At the same time, one can investigate the interaction between free and localized electrons, which provides quantitative data about the fractions of single substitutional atoms and their clusters. This is made possible by measuring (i) the temperature dependence of the magnetic susceptibility derived from the CESR spectra and (ii) the g-factor which stands for the splitting of the electron energy levels under an applied magnetic field (Fig. 3.5) and is different for different chemicals elements. Three electron subsystems contribute to the total magnetic susceptibility of austenitic steels: free electrons, single atoms of transition metals (d-atoms) and superparamagnetic clusters.

3.1 As-Quenched

37

The magnetic susceptibility of free electrons vs0, the so-called Pauli susceptibility, does not depend on the temperature. The magnetic susceptibility of isolated localised d-electrons vd0 changes with the temperature according to the CurieWeiss law: vd0 ¼ C1 =ðT  hp Þ;

ð3:1Þ

where C1 and hp are the Curie constant and the paramagnetic Curie temperature, respectively. The subsystems of free s-electrons and of isolated localised d-electrons can exchange their electrons and the contributions of their interaction to the magnetic susceptibility are described as vs ¼ vs0 ð1 þ a1 vd0 Þ and vd ¼ vd0 ð1 þ a1 vs0 Þ;

ð3:2Þ

where a1 is the parameter of exchange interaction between the free electrons and isolated localised d-electrons (isolated atoms of d-elements). It is seen that the ratio of magnetic susceptibilities of s- and d-electron subsystems v-1 = vs/vd has r to be a linear function of temperature. The superparamagnetic clusters are formed in austenitic steels by the atoms having magnetic moments (the atoms of transition metals with a not-filled electron d-shell like Cr, Ni, Mn etc.). If they are collected in a cluster, its total magnetic moment M is proportional to the number of atoms in the cluster. These moments have an occasional orientation in the crystal lattice, however they can be polarized by the applied external magnetic field. The clusters of d-atoms contribute to the whole susceptibility and change the relative parts of s- and d-contributions. The changed value of vs0 is written as v0s0 ¼ vs0 =ð1  a2 vd2 ðT ÞÞ;

ð3:3Þ

where a2 is the parameter of exchange interaction between the free electrons and clusters, vd2 is the magnetic susceptibility of the superparamagnetic cluster system which obeys the Langevin law (see [26]): vd2 ¼ C2  Lðh=TÞ; C ¼ vd2 at T ¼ 1 K;

ð3:4Þ

Lðh=TÞ ¼ cothðh=TÞ  T=h; h ¼ MH=kB B; where h is the Curie temperature of clusters (proportional to the number of atoms in the cluster), H is an external magnetic field, B is a constant. The temperature dependence of the total relative magnetic susceptibility v1 ¼ r vs =vd is described by the equation 1 1 v1 r ðT Þ ¼ a1 vs0 þ vs0 vd1 þ a2 vd2 vs0 vd1 :

ð3:5Þ

. In other words, the CESR technique allows to estimate the distribution of transition metal atoms in solid solutions (see [7–9] about details).

38

3 Structure

Fig. 3.14 Spectrum of conduction electron spin resonance in steel CN85. Two components are marked with the short dash and dash dot lines

Measuring the temperature dependence of magnetization provides one with the information about single d-atoms (the Curie-Weiss law) and superparamagnetic clusters (the Langevin law) in addition to that given by the CESR measurements. Both techniques allow to estimate the chemical homogeneity of solid solutions on a scale of smaller than 5 nm. We also used measurements of the stacking fault energy in a non-standard manner to estimate the short-range decomposition in solid solutions. The SFE in any solid solution depends only on its chemical composition. Some statements about the dependence of SFE on the grain size or dislocation density do not take into account the effect of grain boundary segregation on the change of composition in the grain bulk and the change of dislocation splitting with increasing dislocation density due to elastic stresses of neighbouring dislocations. In fact, the SFE is mainly controlled by the electron structure, namely by the density of electron states at the Fermi level (see e.g. [27]). For measurements of SFE, the triple node technique was used (see about details [28]). Thus, these three techniques are complementary for studies of the atomic distribution in the solid solutions.

3.1.3.1 Free and Localized Electrons: Electron Spin Resonance Of the four steels studied, the steel MnC steel is the only one for which CESR was not observed, which indicates a very low concentration of free electrons. As example, Fig. 3.14 demonstrates the CESR spectrum of steel CN85, recorded at T = 130 K. It consists of two main components. The wide spectral line belongs to free electrons, whereas the second spectral line is a narrow signal typical for localized paramagnetic centres, i.e. for electrons localized at the atoms.

3.1 As-Quenched

39

Fig. 3.15 a Temperature behaviour of CESR spectra in steel CN85 above 85 K and b Ferromagnetic resonance in CN85 at temperatures B85 K

The spectra of this steel measured in two temperature regions are shown in Fig. 3.15. The spectrum at T = 320 K (see Fig. 3.15a) was recorded simultaneously with the CESR signal of the reference sample containing 8 9 1014 spins, which is seen at a magnetic field of H = 0.16 T. Starting from 85 K, a reversible phase transformation is observed from paramagnetic to some new magnetic state within a temperature range of about 3 K (Fig. 3.15b). The broad absorption signal is very intensive and looks like a ferromagnetic absorption. At temperatures below 83 K, the signal is not remarkably changed, the broad spectral line of free electrons disappears, the narrow line of the localized paramagnetic centres decreases and can be observed in the background of this broad signal. This transformation can indicate the transition from the paramagnetic to a ferromagnetic state or to blocked magnetic moments of superparamagnetic clusters. The temperature dependence of the signal integral intensities is shown in

40

3 Structure

Fig. 3.16 a Temperature dependence of the integral intensities for the CESR signals from free (Ie/I0) and localized (Iloc/I0) electrons and b Resonance magnetic fields Hres,e and Hres,loc in steel CN85. The dashed lines represent the calculated values

Fig. 3.16a in comparison with that of the reference sample. Ie, Iloc are the resonance signal integral intensities of free electrons and those localized at the paramagnetic atoms. I0 means the integral intensity of the CESR signal of the reference sample reduced to T = 1 K. Its temperature dependence is described by the Curie–Weiss law, Iref = I0/T. As expected, the integral intensity of signal Ie of free electrons does not depend on temperature. A small contribution of the localized spin absorption to the signal of free electrons occurs due to the exchange interaction between these two subsystems. Based on the obtained data, the following concentration of free electrons, ne, was obtained for steels CN107, CN96 and CN85, respectively: ne = 2.0 9 1022 cm-3, 2.9 9 1022 cm-3, and 2.8 9 1022 (see Table 3.2). The concentration of superparamagnetic clusters, i.e. clusters of substitutional atoms of alloying elements, is found from the temperature dependence of the resonance fields presented in Fig. 3.16b.

3.1 As-Quenched

41

They cause local magnetic fields on the free electrons. In this case, the temperature dependence of Hres,e of free electrons is described by the temperature function of the local magnetic field caused by superparamagnetic clusters, i.e. by the Langevin function L = coth(h/T)-T/h: ð0Þ Hres;e ¼ Hres;e  B  ðcothðh=TÞ  T=hÞ:

ð3:6Þ

Here, H(0) res,e is the resonance magnetic field for the CESR in the Fe-based alloys in the absence of any internal local magnetic fields; B = 4pMcl is the maximum local magnetic field caused by a system of superparamagnetic clusters with magnetic moment M; h is the energy of an individual superparamagnetic cluster in temperature units of which the magnetic moment under external magnetic field H is M: h ¼ MH=kB ¼ 2lB HN=kB ;

ð3:7Þ

where lB is the Bohr magneton, N is the number of spins in the cluster, kB is the Boltzmann constant. The function (3.6) is shown in Fig. 3.16b by the dashed line with the following values: H(0) res,e = 340 mT; B = 180 mT; h = 200 K. By substituting h from (3.7), the number of magnetic atoms in the cluster was estimated to be about 500 atoms, which corresponds to an average cluster size of about 3 nm. Using the obtained B = 180 mT and M = 0.9210-17 erg1/2cm3/2, one can estimate the concentration of superparamagnetic clusters ncl = B/(4pM) = 1.55  1019 cm-3. Thus, based on the obtained CESR spectra and their analysis, one can conclude that all studied steels have a high concentration of free electrons with its maximum in steels CN96 and CN85. At the same time, along with the localized electrons, i.e. single atoms of substitutional elements, clusters exist in the studied steels.

3.1.3.2 Clusters: Magnetic Measurements Aiming at a clarification of the unusual behavior of the CESR measurements below 85 K and the estimation of the size of magnetic non-homogeneities, we measured the static magnetization of steels CN85, CN107 and MnC within a wide range of magnetic fields. Results are given in Fig. 3.17a for steel CN85 at different temperatures and for steels CN107 and MnC in Fig. 3.17b, c at a temperature of 300 K. A feature of these data is the absence of magnetic saturation and the paramagnetic character of magnetization, which means that the transition at 85 K is not a paramagneticferromagnetic one. In Fig. 3.17a one can see a negligible temperature dependence of the magnetic moment. Versus magnetic field H, it was satisfactorily described for all steels by Eq. (3.8) as a combination of a linear M(H) function and the Langevin function M = const(cothan(MclH/kT)-kT/MclH), which is typical for paramagnetism caused by superparamagnetic clusters (see in detail [25]). For this reason, the curves in Fig. 3.17 were fitted using the function:

42

3 Structure

Fig. 3.17 Magnetic moment versus magnetic field: a steel CN85 at different temperatures. b fitted curve of magnetic moment at 300 K for steel CN107 and c for steel MnC

Table 3.4 Parameters of the magnetization curves in Eq. (3.8) Steel T, K M0, emu/g CN107 CN85

MnC

300 300 240 100 50 4 300

0.12 0.22 0.23 0.23 0.22 0.24 0.01

H0, T

H1, T

0.08 0.2 0.2 0.2 0.15 0.2 0.01

7 6.7 6.5 7.1 7.1 7.2 5.0

M ¼ M0  ðcotanhðH=H0 Þ  H0 =HÞ þ H=H1

ð3:8Þ

with parameters M0, H0, H1 presented in Table 3.4. The constant M0 is equal to glBncl/q, where ncl is the concentration of clusters in the sample, g is factor of spectroscopic splitting, lB is Bohr magneton, q is the density of the alloy. In the Langevin function (cotanh(H/H0)-H0/H), H0 is determined as kT/ (glBns), where ns is the number of paramagnetic spins in a cluster. It is supposed that all clusters have approximately the same size. The linear function versus H can be presented by the Curie-Weiss paramagnetic susceptibility, as well as by the Van Vleck atomic paramagnetism. In case of C–W susceptibility, H1 is obtained

3.1 As-Quenched

43

Table 3.5 Concentration of superparamagnetic clusters, ncl, and average number of paramagnetic atoms per cluster ns Steel ncl, cm-3 ns CN107 CN85 MnC

4.6  1019 8.4 1019 3.5  1018

2570 1012 20000

from the Curie-Weiss law as 4kT/g2l2BnPC, where nPC is the number of the local paramagnetic centers, i.e. that of substitutional solute atoms. The estimation of nPC shows that the value of H1 = 7 T in Table 3.4 could be obtained if the concentration of paramagnetic centers nPC is equal to about 3.4  1022 at T = 300 K and is by one order larger at T = 4 K, which is impossible. Thus, H1 does not depend on temperature and, therefore, the C–W paramagnetism is not justified by the temperature dependence of H1 up to magnetic fields H = 8 T. The only kind of the paramagnetic system with a linear function M(H) and not dependent of T is the Van Vleck paramagnetism [26]. All the iron atoms take part in the formation of the Van Vleck paramagnetic susceptibility. At smaller fields, the C–W paramagnetic susceptibility becomes actual. Using the above determinations, we find the concentration of superparamagnetic clusters ncl and average number of paramagnetic atoms in clusters ns for the steels CN107, CN85 and MnC (Table 3.5). The size of superparamagnetic clusters can be estimated using the number of paramagnetic atoms in the cluster ns. The concentration of clusters in the CN steels is consistent with that obtained for steel CN85 using CESR, although their size is a bit higher. According to the obtained data, steel CN107 is characterized by a low concentration of clusters ncl and a size of about 5 nm, whereas ncl in steel CN85 is higher, however, the clusters have a smaller size of about 3.5 nm. This result suggests that chemical homogeneity in the studied CrMnCN steels is improved with decreasing C/N ratio. The carbon steel MnC is characterized by the smallest concentration of superparamagnetic clusters which have the highest size of about 8 nm. In [25], using small angle scattering of neutrons, the size of chemical clusters in the CrMnCN steels was determined at about 10 nm with an error of about 30 %. The point is that, in comparison with magnetization measurements, the resolution of this technique is not sufficient to study clusters with a size smaller than 5 nm. A much higher neutron scattering in steel MnC was studied with sufficient precision, and the measured size of chemical clusters was about 30 nm. This discrepancy can be evidence of a larger size of chemical clusters in comparison with superparamagnetic ones and needs a special comment. It should be noted that the formation of clusters from substitutional solute atoms can have a different effect on the polarization of the magnetic moments leading to the superparamagnetism and on the enhanced neutron scattering which depends only on the different neutron scattering amplitudes of the nuclei.

44

3 Structure

Fig. 3.18 Frequency distribution of measured dislocation node radii in austenitic steels. a CN96 and b MnC

3.1.3.3 Stacking Fault Energy as a Tool to Estimate the Chemical Homogeneity In addition to the experiments described above, we used the measurement of SFE to obtain information on the chemical inhomogeneity of the studied steels. Such measurements are also useful for testing a possible correlation between the electron structure and SFE in view of the data in [27] where an inverse correlation was observed between the SFE and the density of electron states at the Fermi level. The SFE c was estimated using the following equation (see [28]): c ¼ 0:26 lb2p =ri :

ð3:9Þ

where l is the shear modulus and bp is the Burgers vector of a partial dislocation, ri the radius of a circle inscribed in the dislocation triple node. As an example, the results of measuring the dislocation splitting are shown for steel CN96 and MnC in Fig. 3.18 and the obtained dislocation node radii ri and SFE for all studied steels are presented in Table 3.6. Two maxima of dislocation splitting (see Fig. 3.18) give evidence of shortrange decomposition in the studied steels. The fitting of the experimental data was carried out using a Gauss distribution, which is natural for the statistical scattering of the experimental data. One can see that the homogeneity of the solid solutions is improved with decreasing content of carbon in the CN steels, consequently with decreasing C/N ratio, and it is significantly worse in the carbon steel MnC where the gap between the measured SFE values is the highest (Table 3.6). The obtained values of SFE are presented in Fig. 3.19 along with the data on the concentration of free electrons which are proportional to the density of electron states at the Fermi level D(EF). Qualitatively, the obtained results are consistent with those in [29, 30], where the short-range decomposition in neutron- and electron-irradiated austenitic CrNi steels was studied by local X-ray emission spectroscopy. It was shown that, with a modulation wave of about several microns, the areas enriched in iron and

3.1 As-Quenched

45

Table 3.6 Dislocation node radius ri and SFE r2, nm SFE1, mJ/m2 Steel r1, nm

SFE2, mJ/m2

DSFE, mJ/m2

CN107 CN96 CN85 MnC

43 27 36 41

16 11 7 27

7.3 11.5 10.0 6.4

10.2 16.0 12.0 10.6

59 38 43 68

Fig. 3.19 Stacking fault energy, SFE, and concentration of free electrons ne in CrMnCN steels versus C ? N content in atom %

chromium alternate with those enriched in nickel. Since the electron irradiation does not change the atomic interactions and only, due to enhanced diffusion, assists the thermodynamic equilibrium, it is clear that the atomic distribution in austenitic steels is far from that in ideal solid solutions. Both SFE1 and SFE2 are inversely proportional to the concentration of conduction electrons, which is consistent with the tendency observed in [27] for pure metals. It is relevant to note that the same inverse relation between SFE and D(EF) was observed in [31] for austenitic CrMn steels with nitrogen, whereas a direct proportionality occurred in the same steels if additionally alloyed with nickel.

3.1.3.4 Effect of C/N Ratio on the Chemical Homogeneity The main question is how the C/N ratio in the austenitic CrMnCN steels affects their chemical homogeneity. The following analysis can be made based on the measurements of electron spin resonance, magnetization and stacking fault energy. In comparison with nitrogen steels studied in [7–9], the ones with carbon ? nitrogen are characterized by an increase in the concentration of free electrons. The data presented in Table 3.2 show that the concentration of free electrons generally increases with a decreasing C/N ratio. Also there is no remarkable difference between steels CN96 and CN85, which points to some optimum of the C/N ratio. In addition the data of CESR (Sect. 3.1.3.1) and magnetization (Sect. 3.1.3.2) show that, with decreasing C/N ratio, the density of clusters increases, whereas their size decreases. In other words, the atomic

46

3 Structure

distribution becomes more homogeneous, which suggests a tendency to shortrange atomic ordering. This result is consistent with earlier data on the correlation between the metallic character of interatomic bonds and the atomic distribution in austenitic steels with nitrogen and carbon [32] and allows to assume that an appropriate C/N ratio can provide the highest thermodynamic stability of austenitic steels. A correlation between atomic interaction and atomic distribution is also confirmed by the markedly increased size of clusters in Hadfield steel MnC (see the data of magnetization in Sect. 3.1.3.2) and the absence of the CESR signal in this steel, which is a sign of an extremely low concentration of free electrons. As a result, the Hadfield steel reveals a rather low stability in relation to precipitation at low temperature heating. In view of the decreased concentration of free electrons, i.e. the prevailing covalent character of interatomic bonds in the Hadfield steel, a special comment can be given in relation to its good impact toughness, though it is smaller compared with austenitic steels containing nitrogen or carbon ? nitrogen (see [15]). The stacking fault energy of about 50 mJ/m2 [33] is not favourable for strain-induced c ? e transformation. At the same time, the carbon enhanced covalent bonds locally increase the shear modulus within the carbon clouds around dislocation, which increases the line tension of dislocations and, correspondingly, decreases their mobility. These two factors should create appropriate conditions for the extremely intensive twinning in the Hadfield steel by cold working. In fact, the Hadfield steel is a good example of the TWIP effect, which is the real reason for its tough mechanical behaviour. As shown in [33], such an intensive twinning is not observed in austenitic steels with nitrogen or carbon ? nitrogen. Finally, let us discuss the applicability of SFE measurements to estimate the chemical homogeneity of solid solutions. The distance between the SFE maxima DSFE in Table 3.6, decreases with decreasing C/N ratio, which is consistent with the data of chemical inhomogeneities obtained by other experimental methods. Summing up one should state that the studied austenitic CrMnCN steels are characterized by a high concentration of free electrons, which evidences the enhanced metallic character of interatomic bonds. A feature of the atomic distribution in the austenitic solid solutions is the existence of clusters rich in Cr and Mn, which can be detected by their superparamagnetism or the enhanced small angle scattering of neutrons on the nuclei of atoms in the chemical inhomogeneities. The size of clusters decreases with a decreasing C/N ratio, whereas their number increases, which suggests that the chemical homogeneity of steels increases. The same conclusion follows from the measurements of the stacking fault energy. A clear gap in the frequency curves, showing the distribution of dislocation splitting, corresponds to the chemical inhomogeneity of solid solutions on a nano-scale. The difference in the measured values of the stacking fault energy DSFE corresponds to the short-range decomposition which decreases with decreasing C/N ratio.

3.2 Structural Change by Loading

47

3.2 Structural Change by Loading 3.2.1 Tensile Straining The aim of this section is to analyse the influence of carbon, nitrogen and carbon ? nitrogen on the change of structure in austenitic steels caused by cold work and, based on the obtained data, discuss the controlling mechanism of cold work hardening. The enhanced strengthening by plastic deformation is a general feature of ironbased interstitial solid solutions. In relation to austenitic steels, a number of studies were devoted to Hadfield steel MnC (e.g. [34–50], and, nevertheless, the nature of its extremely high cold work hardening remains a controversial issue. In early publications, it was attributed to strain-induced c ? e [34–37] or c ? a [38, 39] martensitic transformations. These observations were later explained also as a result of segregation [39], precipitation [40] or decarburization [41]. Perhaps for the first time, the idea of twinning as a reason for increasing the cold work hardening of Hadfield steel was expressed in [42]. Another mechanism based on dynamic strain ageing, namely the break-up of pinned dislocations from carbon clouds and the restored pinning due to accelerated pipe diffusion of carbon atoms along the dislocations was proposed in [43], where both the enhanced cold work hardening and the serrated flow were explained in such a way. However, it should be noted in this relation that the enhanced migration of atoms by pipe diffusion occurs only for substitutional solutes because only the migration of vacancies is accelerated along the dislocation pipes, and that is not a migration mechanism for interstitials. On the contrary, in comparison with the bulk, the interstitial atoms lose their mobility in the vicinity of dislocation cores because the elastic term practically disappears from the gradient of chemical potential (see, e.g., experimental data [44] about the retarded carbon diffusion in previously cold worked iron and other metals). In further studies of Hadfield type steels the formation of twins or slip depending on the crystallographic orientation was confirmed [45], whereas the decisive role of dynamic strain ageing was denied [46]. In studies of austenitic high nitrogen CrMnNi steels, the cold work hardening was attributed from the very beginning to the nitrogen-caused planar slip at lower strains and deformation twinning at higher strains [47–49]. At the same time, the increase in dislocation density during cold working was considered as a factor preventing brittle fracture [50]: the dislocations shield stresses created by intersecting twins where the cracks are thought to be nucleated. Alloying with nickel assists the dislocation mode of plastic flow and shifts twinning to higher strains. With decreasing deformation temperature or increasing nitrogen and manganese contents, the onset of twinning is shifted to lower strains, whereas the dislocation density decreases in between the twins. A particular role in cold work hardening of nitrogen steels was ascribed to secondary twinning, i.e. the formation of twins within the space between the

48

3 Structure

primary twins [51]. The multiple twin system was analysed in terms of an increase in the efficiency of twins as hardening obstacles for gliding dislocations and formation of new twins. At the same time, no remarkable twinning was found in other studies of cold worked austenitic nitrogen steels. For example Kubota et al. [52] examined the structure of SUS316L steels (i.e. CrNi ones) with 0.02 and 0.56 mass % nitrogen and Cr18Mn18 type steels with 0.51 and 0.84 mass % nitrogen after tension tests and observed an increase in work hardening with increasing nitrogen content and no deformation twinning but planar dislocation arrays forming Lomer-Cottrell barriers at their intersections. In a similar way, Saller et al. [53] demonstrated the microstructures of strained CrMn austenitic steels alloyed with 0.3 or 0.9 mass % nitrogen, where planar slip was the main structural feature in case of high nitrogen content, whereas twins and e-martensite appeared occasionally at rather high strains. This discrepancy between the observations made by different authors could be caused by a different crystallographic texture. One of the factors controlling twinning in austenitic steels, as well as strain-induced c ? e formation, is the stacking fault energy: the lower the SFE the smaller is the stress for twin formation. According to [54, 55], the actual stress on the leading partials can be higher or lower than that on the trailing ones depending on the crystal orientation in relation to the tensile axis (see more about details in the next Section). Correspondingly, the dislocation splitting becomes larger or smaller, which should affect twinning or c ? e transformation. As shown by Lee et al. [56], this results in the orientation dependence of twinning during tensile deformation of high nitrogen austenitic steel Cr18Mn18Mo2N0.9. In relation to the tensile axis, two twinning systems, primary and conjugating, were observed in \111[ grains, whereas only one twinning system was activated in \110[ grains and no deformation twinning occurred in \100[ grains.

3.2.1.1 Yielding The yield strength belongs to the most important characteristics of engineering materials used by designers of parts or constructions which are exploited under loading. In absence of precipitates, it is determined by the following structural factors: (i) the pinning of dislocation sources by clouds of solute atoms, (ii) the potential relief of crystal lattice created by interactions of host atoms resulting in the so-called Peierls barriers, the height of which increases with decreasing temperature (see, e.g. [57]), (iii) the interaction between gliding dislocations and ‘‘dislocation forest’’ consisting of the dislocations intersecting the slip plane, (iv) the effect of solute atoms which create fields of elastic stresses to be overcome by dislocations, (v) the grain size which limits the length of the slip plane and, thereby, the number of dislocations in their plane ensemble, called pile-up. The latter is particularly important because the shear stress sl acting on the leading

3.2 Structural Change by Loading

49

dislocation in the pile-up increases with the number of dislocations n in the pile-up as sl = ns, where s is the applied stress. The dislocation segments in hot forged and solution annealed steels are usually surrounded by clouds of interstitial atoms, which requires an additional stress for slip activation. If so, the start of yielding can be accompanied by a sudden decrease of loading, ‘‘yielding tooth’’, which in early studies was attributed to the breaking of dislocations from interstitial clouds [58]. However, as shown later by experiment [59, 60], as well as in the theoretical estimations [61], the stress needed to break the pinned dislocations from interstitial clouds is so high that it exceeds even the ultimate strength and, therefore, the emission of new fresh dislocations by dislocation sources should occur before unpinning. In turn, the yielding tooth can result from dynamics of dislocation multiplication in a solid solution where dislocations are blocked by solute atoms. In contrast to bcc materials, the Peierls relief in metals with an fcc crystal lattice, which is just the case of austenitic steels, is shallow and cannot really affect dislocation slip. The same applies to a possible asymmetry of dislocation nuclei, which in bcc crystals results from a small splitting of dislocations in two or three atomic planes and is proposed instead of the Peierls relief to be responsible for a striking increase of the yield strength with decreasing temperature [62]. Nitrogen in austenitic steels remarkably enhances the temperature dependence of the yield strength, which is not typical of materials with an fcc crystal lattice. This feature stimulated an idea supported by theoretical calculations [63], according to which nitrogen in austenitic steels causes an asymmetry of dislocation nuclei and, for this reason, makes the temperature behaviour of the yield strength similar to that of bcc materials. However, the experimental data [64, 65] have clearly shown that the temperature behaviour of the yield strength in austenitic nitrogen steels is controlled by the intersection of dislocations, as predicted by Seeger’s theory for fcc crystals [66]. Therefore, some other mechanisms should exist for the favourable nitrogen effect on the yielding of austenitic steels at low temperatures (see Sect. 3.2.2). In interstitial solid solutions, moving dislocations overcome elastic stresses created by interstitial atoms, which increases the term r0 in the Hall–Petch equation for the yield strength: ry ¼ r0 þ ky d1=2 ;

ð3:10Þ

where d is the grain size and the Hall–Petch coefficient ky is an ‘‘unblocking constant’’ characterizing the stress needed for the activation of dislocation sources in a neighbouring grain in order to transfer deformation through the grain boundary (see, e.g., [67] for details). Both r0 and ky depend on the affinity of interstitial atoms to dislocations, which is characterized by the binding enthalpy. In contrast to a-iron where the binding enthalpy of carbon and nitrogen atoms with dislocations is practically the same, it was shown [68] that austenitic steels are distinguished by a significantly stronger affinity of nitrogen atoms to dislocations. This feature has its nature in a higher

50

3 Structure

elastic interaction between nitrogen atoms and dislocations in comparison with that of carbon atoms in iron austenite, as well as in the additional chemical interaction due to dislocation splitting. Judging from the concentration dependence of the a-iron lattice parameters [69], there is no remarkable difference between carbon and nitrogen atoms in their elastic contribution to the dislocation-interstitial interaction, which is dominant in bcc metals. On the contrary, nitrogen atoms in iron austenite cause higher lattice distortions in comparison with carbon ones. Moreover, their chemical interaction with dislocations stems from a different solubility of interstitial atoms in the fcc solid solution and in the stacking faults which in fcc crystals have a hcp structure. Nitrogen can be dissolved in hcp iron in a wide concentration range (see, e.g., the binary Fe–N diagram in [70]), whereas the carbon solubility in hcp iron is limited. The oversaturation of stacking faults with carbon is clearly illustrated by precipitation of cementite in plates of the e-martensite transformed in the course of fatigue tests of steel CN96 (see Sect. 3.2.4.2). Thus, so-called Suzuki’s atmospheres of carbon and nitrogen atoms at split dislocations add to the term r0 which generally consists of contributions from the lattice friction stress (Peierls relief), the stress needed to overcome distortions created by solute atoms and the intersecting dislocations penetrating the slip plane. The interaction between gliding dislocations and dislocation forest is enhanced due to their splitting in austenitic steels because the stress needed for the formation of a constriction at the point of intersection is the higher the larger the splitting is. Along with this, the carbon and particularly nitrogen atoms effectively contribute to the increase of the coefficient ky enhancing thereby the grain boundary strengthening of austenitic steels, which was extensively studied in [71– 75] and interpreted in [67]. Based on findings mentioned above, one can predict that both carbon and nitrogen in austenitic steels increase the yield strength and the effect of nitrogen is expected to be more pronounced. Results of mechanical tests of C ? N austenitic steels will be presented in Sect. 4.1.

3.2.1.2 Effect of Stacking Fault Energy It is natural to suppose that a different mechanical behaviour is controlled by a different substructure formed during straining of austenitic steels with carbon, nitrogen and carbon ? nitrogen. The stacking fault energy SFE is responsible for dislocation slip, as well as for twinning or formation of e during plastic deformation. Usually, an SFE smaller than 100 mJ/m2 is considered to be low (see, e.g., [76]). Dislocation slip prevails at high SFE, which makes the dislocations narrow and eases a change of the slip plane. Some middle SFE value assists twinning along with dislocation slip. Strain-induced c ? e transformation occurs at sufficiently low SFE.

3.2 Structural Change by Loading

51

Fig. 3.20 Effective splitting of leading (1/6½211) and trailing (1/6½112) partial dislocations under applied shear stress in the fcc lattice

It is worth noting that the SFE values should be measured under condition of extremely low density of dislocations, which is achieved due to solution treatment. This ‘‘true’’ stacking fault energy depends only on the chemical composition of solid solutions and determines their thermodynamic stability. In cold worked steels, a so-called ‘‘effective’’ stacking fault energy (see [54, 55]) is controlled by an increase or decrease of dislocation splitting depending on the orientation of the Burgers vectors of the leading and trailing partial dislocations in relation to the acting applied stress or to the retained stresses after deformation (Fig. 3.20). This is why the change of crystallographic texture by deformation and the different crystallographic orientation of grains require the measurement of a number of triple node radii and statistical analysis of the experimental data as done in Sect. 3.1.3.3 (see Fig. 3.18). In the studied steels, the SFE was measured to be equal to 52 in steel MnC, 36 in steel CrMnN and 40 in steel CN96 with an accuracy of about ± 3 mJ/m2. The value obtained for steel MnC is much higher than follows from the thermodynamically estimated 22 mJ/m2 in [77] and correlates better with the well known effect of increasing carbon on the SFE of austenite (see e.g. [78, 79]). Within the error of measurements, the SFE values of steels CrMnN and CN96 are the same, which suggests a nearly similar deformation behaviour.

3.2.1.3 Substructure of Cold Worked Steels Structural and phase transformations caused by cold working the carbon steel MnC, nitrogen steel CrMnN and carbon ? nitrogen steel CN96 of Table 2.1 are going to be analyzed in the attempt to find features responsible for their mechanical behaviour. For this purpose, tensile specimens were strained to a pre-selected elongation between 10 and 50 %. The respective true stress r/true strain u curves were derived from engineering stress/strain curves up to necking (see Fig. 4.1). The exponent and coefficient of Ludwik’s Eq. (3.11) were taken from the final slope of these curves (Table 3.7).

52

3 Structure

Table 3.7 Exponents and coefficients of the Ludwik Eq. (3.11) after pre-selected elongation Elongation e [%] 10 20 30 50 (true strain u) (0.09) (0.18) (0.26) (0.41) MnC CrMnN CN96 a

n K n K n K

0.29 1195 0.28 1700 0.28 1585

0.49 1770 0.41 2185 0.41 2060

0.65 2250 0.43 2270 0.46 2315

0.77 2600 0.41a 2200a 0.53 2430

taken at e = 40 %, i.e. below Au

r ¼ K/n :

ð3:11Þ

As follows from the obtained data, the intensive cold work hardening of Hadfield steel MnC is supported by Ludwik’s exponent n which tops the other steels at the high end of elongation. While n of MnC and CN96 increases from 10 to 50 % elongation, the cold work hardening of CrMnN was slightly reduced at 40 % leading to the lowest uniform elongation Au of the three steels investigated (Table A1). At the elongation of 50 %, K changes proportional to the content of interstitial atoms, be it carbon or nitrogen. The change in the structure of three selected steels is analysed. It is expedient to start from a general picture which can be obtained using the optical microscopy even if it could not be decisive in clarifying the difference between mechanisms of cold work hardening in the studied steels. As an example, the sequence of structural change with increasing tensile strain is presented in Fig. 3.21 for steel CN96 (see also [80]). The strain of 10 % causes separate slip lines in some grains (Fig. 3.21a). Even at 30 % of strain, not all the grains are involved in slip (Fig. 3.21b), possibly because of the difference in their orientation to the applied stress. At the same time, a sign of the slip on secondary planes can be already seen. At 50 % of strain, which is the limit of the uniform elongation, the density of slip lines increases significantly and the slip becomes more homogeneous (Fig. 3.21c). Nevertheless, the deformation degree is different for different grains. An important feature of the microstructure after 50 % strain is also the curvature of slip lines, which is clearly due to the interdependence in the deformation of the neighbouring grains having a different crystallographic orientation. Such a character of the slip lines also indicates that twinning and c ? e transformation are not preferential as the deformation mode of this steel. Transmission electron microscopy allows one to distinguish further features in structure of the three studied steels which can be responsible for their different cold work hardening and plasticity limit (see about details [33]). An intensive twinning is observed in steel MnC at a rather small elongation of 10 % (Fig. 3.22a). Remarkable is the large thickness of the twin plates. Twinning continues at higher strains and, additionally, the dislocation density significantly

3.2 Structural Change by Loading

53

Fig. 3.21 Microstructure of steel CN96 after tension tests with deformation degree. a 10 %. b 30 %. c 50 %

Fig. 3.22 Structure of steel MnC after a 10 % elongation with zone axis [112], twin zone axis ½112and twinning plane ð 1 11Þ. b 30 % elongation with zone axis [110], twin zone axis ½ 1 10 and twinning plane ð111Þ

54

3 Structure

Fig. 3.23 a Planar slip and b tangled dislocations in steel CrMnN after 10 % elongation

increases (Fig. 3.22b). The dislocations are tangled and reveal a weak tendency to form a cellular structure. No traces of a or e martensite are observed. Planar dislocation slip (Fig. 3.23a) and a tangled dislocation structure (Fig. 3.23b) are developed in the range of small elongations of the nitrogen steel CrMnN. The nitrogen-induced short-range atomic ordering is the obvious reason for the formation of planar dislocation arrays. An increase in elongation up to 30 % results in planar gliding on intersected slip planes (Fig. 3.24a). At the same time an essential twinning appears in this strain range (Fig. 3.24b). Moreover, two systems of twinning are found, one of which, namely {113} \120[ , is uncommon. Possibly, new twinning systems are due to conditions of retarded strain in the highly strengthened fcc lattice. The occurrence of intensive twinning is at variance with the data of [53] where the authors observed very rare twinning in the same steel after tensile deformation. Intensive twinning, including that on intersecting planes, occurs also if steel CrMnN is elongated by 50 % (Fig. 3.25). Again, an uncommon twinning plane is observed. A small tensile deformation of steel CN96 causes planar slip like that in the pure nitrogen steel CrMnN (Fig. 3.26). Twinning is involved in plastic deformation at higher elongation (Fig. 3.27a). At an elongation of 50 %, twinning continues and one can observe intersecting twinning systems (Fig. 3.27b). Planar slip continues to be an essential feature of the substructure up to the elongation of 50 % (Fig. 3.28), which is a bit different from the observations in steel CrMnN where it was not met. In both deformed steels, CrMnN and CN96, one can rarely find strain-induced e martensite. As an example, e plates in steel CN96 subjected to an elongation of 50 % are shown in Fig. 3.29. As follows from the diffraction analysis, the Nishiyama-Wassermann orientation relation occurs between the fcc and the hcp crystal lattices.

3.2 Structural Change by Loading

55

Fig. 3.24 a Planar slip on intersecting planes and b twinning on intersecting twinning systems: with zone axis [231], twin zone axis ½ 2 31, twinning plane ð1 11Þ and twin zone axis ½ 1 20,  twinning plane ð113Þ in steel CrMnN after 30 % elongation

Fig. 3.25 Twins in steel CrMnN after 50 % elongation with zone axis ½013 and two systems of twinning: with twin zone axis ½013, twinning plane (131), and twin zone axis ½35 4, twinning plane (111)

Fig. 3.26 Planar slip in steel CN96 after 10 % elongation

56

3 Structure

Fig. 3.27 a Twins in steel CN96 after 30 and b twinning at intersecting planes after 50 % elongation Fig. 3.28 Planar slip in steel CN96 after 50 % elongation

3.2.1.4 Mechanisms of Cold Work Hardening It is natural to discuss possible mechanisms of cold work hardening in austenitic steels starting from Seeger‘s theory for the yield stress and cold work hardening in fcc crystals of low stacking fault energy [66, 81]. According to Seeger, in absence of twinning or strain-induced fcc ? hcp transformation, the yield stress and cold work hardening of fcc crystals of low stacking fault energy is controlled by gliding dislocations cutting through the dislocation forest. In this case, the splitting of dislocations, i.e. the stacking fault energy, plays an important role because the jog formation at dislocation intersections needs a preliminary constriction of split dislocations. If alloying with interstitial elements does not lead to a remarkable dislocation splitting, they contribute mainly to solid solution hardening, i.e. to the yield stress, whereas the flow stress reveals an insignificant dependence on the content of interstitials.

3.2 Structural Change by Loading

57

Fig. 3.29 a Plates of e martensite in steel CN96 after 50 % elongation. b Diffraction pattern from the area 1 in a with fcc zone axis [111]. c Diffraction pattern from area 2 in a with fcc zone axis ½111 and hcp zone axis ½ 3 61 which are in Nishiyama-Wassermann orientation relation. d Dark field in the light of e reflection marked with a circle in c

Such a conclusion is confirmed by mechanical tests of the high nickel austenitic nitrogen steel Cr18Ni16Mn10 (Fig. 3.30). Nickel significantly increases the SFE in austenitic steels [79], making dislocations narrow, whereas nitrogen in this steel, increased from 0.06 to 0.5 mass %, changes the SFE non-monotonously between 43 and 65 mJ/m2 (see [31]), which does not cause a remarkable dislocation

58

3 Structure

Fig. 3.30 Effect of nitrogen on microhardness of steel Cr18Ni16Mn10 after solution treatment (e = 0 %) and its increment due to cold rolling [70]

splitting in comparison with the nitrogen-free steel. This is why nitrogen contributes mainly to solid solution hardening of this steel (see the curve at e = 0 %), whereas cold work hardening is not changed with an increasing degree of deformation. A similar result was obtained in [82], where the authors did not find any effect of nitrogen on the hardness of the cold rolled steel Cr18Ni16Mo2. An increase in the chromium content can compensate this Ni effect and recover the nitrogen-induced cold work hardening in high nickel austenitic steels (see e.g. [83]), because chromium decreases the SFE in iron austenite [79]. The effects of carbon, nitrogen and carbon ? nitrogen on the substructure under conditions where twinning and/or c ? e transformation accompany the plastic flow are to be compared. The extremely high cold work hardening in Hadfield steel MnC is obviously connected with intensive twinning which starts at small strains (see Fig. 3.22). Twinning occurs in spite of the rather high stacking fault energy of 52 mJ/m2, i.e. in spite of the reduced dislocation splitting. This is thought to be related to a strong elastic interaction between carbon atoms and dislocations. Because of the high concentration of interstitial atoms, their elastic fields hinder dislocation slip and, as an alternative mode of plastic deformation, promote twinning even in case of comparatively narrow dislocations. It is also worth noting that the twins in steel MnC are unusually thick, which additionally enhances cold work hardening. As mentioned in Sect. 3.1.3.4, Hadfield steel is a good example of the TWIP effect. At higher strains, along with twinning, a significant increase in the density of tangled dislocations without any inclination to localized slip should effectively contribute to cold work hardening. It is worth noting that the chaotic distribution of dislocations contributes to cold work hardening much more than a cellular dislocation structure. The cold work hardening is not as high as in steels CrMnN and CN96 at small elongations up to 10 %. The only feature of the substructure is planar slip on one of acting slip systems (see Figs. 3.23 and 3.26). This planar slip is usually attributed to short-range atomic ordering based on the studies [54, 55]. However, the role of the nitrogen-induced planar slip in cold work hardening is sometimes

3.2 Structural Change by Loading

59

exaggerated (e.g. [53]), at least at moderate strains where planar slip bands do not intersect each other. To discuss this topic again, the example of high nickel austenitic nitrogen steels is used. As mentioned above, nitrogen does not contribute to cold work hardening if the nickel content is high. At the same time, pronounced planar slip occurs during deformation (see, e.g., data on planar slip in steel Cr26Ni32N0.36 [84]). Thus, it is hardly possible to attribute cold work hardening in steels CrMnN and CN96 to planar slip, at least at moderate strains. Only the intersection of planar slip systems at higher strains should contribute to cold work hardening due to formation of Lomer-Cottrell barriers which effectively block the slip in active {111} planes. In this relation one should remark that, at large strains, slip planarity on intersecting slip systems is met statistically more often in steel CN96 than in CrMnN, which can provide a higher cold work hardening in steel CN96. This can be also attributed to a stronger short-range atomic ordering in carbon ? nitrogen austenitic steels as compared with the nitrogen steel (see [22]). A remarkable twinning, although not as intensive as in steel MnC, along with a moderate c ? e transformation is observed at elongations of 50 %, and one cannot really find a clear difference in the structure evolution in both steels alloyed with nitrogen or nitrogen ? carbon except for some prolonged planar slip in steel CN96. A similar change of structure during plastic deformation is possibly controlled by nearly the same stacking fault energy. It seems that additional arguments could be obtained by measuring the binding enthalpy between dislocations and interstitials in both steels. As mentioned earlier, according to [68], the pinning of dislocations by nitrogen atoms in austenitic steels is stronger than that by carbon atoms. A higher affinity of nitrogen atoms to dislocations is one of the reasons for some additional stress needed for cutting the gliding dislocations through a dislocation forest pinned by interstitial atoms. It is so far unknown how strong the pinning of dislocations by carbon ? nitrogen atoms is. One can also predict that the enhanced metallic character of interatomic bonds, which is brought about by the higher concentration of free electrons in nitrogen and particularly in carbon ? nitrogen steels (see Sect. 3.1), has to retard the opening of cracks during plastic deformation and assist a higher fracture elongation of N and C ? N steels.

3.2.2 Effect of Subzero Temperature In contrast to austenitic carbon steels, a feature of the nitrogen ones is an enhanced temperature dependence of the yield and flow stress. It was shown in [65, 85–88] that, while at room temperature the difference between the tensile strength of austenitic carbon and nitrogen steels does not really exceed the error of measurements, it increases with decreasing temperature. According to [88], nitrogen at

60

3 Structure

Fig. 3.31 Temperature dependence of the yield strength in steel Cr18Ni16Mn10 not purposefully alloyed with carbon or nitrogen and alloyed with carbon or nitrogen [65]. The data for steels MnCr82, CN85, CN96, CN107 and CrMnN stem from [80]

4 K is more effective strengthener of austenitic steels, by a factor of 1.4, than carbon. Based on these studies, the idea was proposed [89] that nitrogen assists a bcclike behaviour of the flow stress of austenitic steels with temperature, which is strongly different from the behaviour of fcc substitutional solid solutions, and theoretical calculations [63, 90] were performed, according to which the alloying of austenitic steels by nitrogen causes a splitting of the core of screw dislocations on two or more non-parallel planes, exactly as it takes place in the crystals with a bcc lattice. In contrast to these calculations, Obst and Nyilas [64, 76, 91] obtained by precise measurements that a three-stage temperature dependence of the flow stress occurs in austenitic nitrogen steels in accordance with Seeger’s theory [66, 81] for fcc crystals of low stacking fault energy (i.e. lower than 100 mJ/m2). These three stages include the intersection of dislocation forest by edge dislocations (stage 1), screw dislocations (stage 2) and emission of vacancies which is needed for the slip of a jog formed in case of two intersecting screw dislocations (stage 3). According to [65], nitrogen in austenitic steels affects the third stage (see Fig. 3.31). As the constriction of split dislocations is a prerequisite for jog formation, the effect of nitrogen on the low temperature strengthening has its origin in the nitrogen-enhanced temperature dependence of the stacking fault energy: the more split the dislocations are, the higher is the stress needed for their constriction. As shown in [65], nitrogen enhances a decrease of the stacking fault energy with decreasing temperature in consistency with the temperature dependence of the density of electron states D(EF), i.e. the concentration of free electrons, and the inverse correlation between SFE and D(EF). Results of steels listed in Table 2.1 are also presented in Fig. 3.31. Quantitative data are given in Table 3.7 and A6. From these it follows that, at higher contents of interstitials, nitrogen and carbon ? nitrogen increase the yield strength and its temperature dependence within the range of the 2nd stage, where the intersection of gliding screw dislocations with the dislocation forest controls their slip, as well

3.2 Structural Change by Loading

61

as at the 3rd stage, where the emission of vacancies is needed for slip of a jog formed due to the intersection of screw dislocations. This is obviously due to the C- and (C ? N)-caused temperature dependence of the stacking fault energy which is enhanced by nitrogen and should be even more enhanced by carbon ? nitrogen (see effect of C, N and C ? N on the concentration of free electrons in Fig. 3.6). One can see that the uniform as well as the total elongation remains high at 173 K, which evidences a good low temperature plasticity of the CrMnCN steels (Table A6). The coefficients K and n in the Ludwik Eq. (3.11) are increased with decreasing temperature. As follows from the data of [64] for the austenitic nitrogen steel Cr25Ni15N0.35, the coefficient n can be even decreased within the temperature range of stage 3 in Fig. 3.31. Considering an increasing dislocation density, the constriction of the split dislocations should require a lower applied stress because of stresses caused by neighbouring dislocations. Thus, a high ductility accompanies the strengthening of austenitic (C ? N) steels at temperatures as low as -100 °C (Sect. 4.1.2). As will be shown in Sect. 4.1.6, this is also true for the impact toughness except for the phenomenon of ductile-to-brittle transition, which is very unusual for materials with an fcc crystal lattice and will be discussed in the following Section.

3.2.3 Effect of Strain Rate 3.2.3.1 Rapid Tensile Tests Tensile tests of steel MnCr82 alloyed with C ? N (Table 2.1) were carried out at an initial strain rate between 3.3 9 10-4 and 3.3 9 10-2 [92]. As the rate was raised, the yield strength increased, the ultimate tensile strength and elongation decreased, whereas the reduction of area remained nearly unchanged. The strain rate dependence of plastic flow and fracture of austenitic nitrogen steels was studied by Tomota et al. in more detail [93]. They tested CrMn and CrNi steels varying the strain rate by 3 orders of magnitude from 2.7 9 10-3 s-1 to 2.7 9 10-1 s-1. The authors changed the nitrogen content (mass %) from 0.513 to 0.844 in a CrMn steel and substituted Mn by Ni in a steel with 0.556N. Raising the strain rate at RT increases the yield strength and reduces the elongation and reduction of area. The ultimate strength behaves non-monotonous, depending on the nitrogen content. At 77 K, no plastic deformation and cleavage fracture occurred in steel Cr19Mn19N0.844 having the highest nitrogen content, whereas some elongation was obtained in steel Cr17Mn19N0.513, which is generally expected. However, surprising is the absence of necking and a mixture of dimple and cleavage-like fracture in steel Cr17Mn19N0.513 at the small strain rate of 2.7 9 10-3 s-1, whereas necking and ductile dimple fracture occurred at the higher strain rate of 2.7 9 10-1 s-1. It is remarkable that the same effect (necking and dimple fracture)

62

3 Structure

is obtained at the strain rate of 2.7 9 10-3 s-1 in steel Cr17Mn1Ni13N0.556, where, at the same chromium content, manganese is substituted by nickel. Tomota et al. [93] interpreted their results of steel Cr17Mn19N0.513 in terms of ‘‘work softening’’ caused by an abrupt increase of local temperature in slip bands due to rapid loading. Another possible explanation can be based on the effect of nitrogen on the atomic interactions, as discussed in Sect. 3.1. Increasing the concentration of free electrons by nitrogen enhances the metallic character of interatomic bonds, which should assist ductile fracture. However, as follows from Fig. 3.6, the effect of nitrogen on the concentration of free electrons in austenitic nitrogen steels is nonmonotonous. Reaching a maximum at about 0.5 mass % of nitrogen, the concentration of free electrons decreases with a further increase in the nitrogen content and is rather small in steel Cr18Mn20N0.88 the composition of which is close to that of steel Cr19Mn19N0.844. The favourable effect of nickel can be attributed to the nickel-caused increase of the concentration of free electrons in iron austenite (see [1]), which increases plasticity at RT, as well as at 77 K, and provides a ductile fracture. Taking into account this correlation between the concentration of free electrons and mechanical properties, one can foresee a favourable effect of alloying CrMn steels with C ? N (see Fig. 3.6), which will be shown in Sect. 4.1. A further increase in the strain rate up to that of impact loading leads to the phenomenon of a ductile-to-brittle transition, DBT, which is typical for metallic materials with a bcc crystal lattice but is unique for those with an fcc one. Tobler and Meyn [94] were probably the first to observe the quasi-cleavage of austenitic nitrogen steels during impact tests. It was a special feature that the fracture occurred along the crystallographic {111} planes, i.e. along the close– packed slip planes, whereas the cleavage in the bcc crystals proceeds in {001} planes which are not close-packed. This cleavage-like fracture was observed in notch impact as well as in tensile tests performed at subzero temperatures. In contrast, the toughness of austenitic carbon steels gradually decreases with decreasing temperature and increasing strain rate. The following features characterise this kind of brittle fracture: (i) alloying elements have a strong effect (nickel decreases, whereas nitrogen, manganese and chromium assist cleavage, see [95–97]; (ii) the grain size has a negligible effect [98], which is quite different from the behaviour of bcc materials, where cleavage is markedly enhanced by increasing the grain size. A number of hypotheses were proposed for interpretation of this unusual phenomenon. Tobler and Meyn [94] suggested a critical role of planar slip resulting in an accumulation of shear strain in the narrow bands, their weakening by accumulated dislocations and final separation. According to Müllner et al. [99], the intersection of deformation twins is a reason for the nucleation of cracks. Nitrogen, as well as manganese, shift the onset of twinning to lower strains where the density of dislocations is low and a critical stress for brittle fracture can be achieved before the crack tip blunts due to dislocation slip. This mechanism is at

3.2 Structural Change by Loading

63

clear variance with the data about the absence of twinning and the high density of dislocations just under the fracture surface, as obtained by Tomota et al. [100]. Moreover, these authors have clearly demonstrated the localization of plastic flow, the nitrogen-enhanced slip on the {111} planes (‘‘slipping-off’’, according to [100]), the nucleation of submicron cracks in these planes and their fusion on nonactive {111} slip planes, which creates cleavage-like facets on the fracture surface. Based on the data about the effect of nitrogen on the electron structure and short-range atomic order, we can add the following details to the interpretation of DBT in austenitic nitrogen steels proposed by Tomota et al. First, the localization of dislocation slip is obviously related with nitrogencaused short-range atomic ordering. It is worth noting that this ordering is not simply related with the preferential Cr–N atomic bonds, as is often claimed. Due to the increase in the concentration of free electrons, nitrogen changes the distribution of substitutional solutes assisting preferential bonds between atoms of different kind. In contrast, carbon enhances covalent interatomic bonds, which assists clustering of one kind of atoms. Second, the nitrogen-caused ‘‘slipping-off’’, according to the terminology proposed in [100], is caused by the increase in the concentration of free electrons resulting in a decreased shear modulus l and, consequently, line tension of dislocations C & lb2, thereby increasing their mobility. Additionally, the nitrogenenhanced Snoek-Köster relaxation and strain dependence of damping can be mentioned as an experimental evidence of nitrogen-enhanced mobility of dislocations (see e.g. [101]). Third, the impact loading is accompanied by an ‘‘abnormal mass transfer’’ (e.g. [102]), of which the transport of solute atoms by dislocations is a possible mechanism [103, 104]. If so, the nitrogen atoms should accompany dislocations in the course of impact loading and locally enhance the metallic character of interatomic bonds within the nitrogen clouds around the dislocations. As a result, the shear modulus l is locally decreased in the vicinity of dislocations. Their line tension decreases and their mobility increases. Moreover, in case of dislocation slip accompanied by a transfer of nitrogen atoms, the distance d between the dislocations in their plain arrays (pileups) decreases as d & (plb)/16(1-m)ns, where l is the shear modulus, b is the Burgers vector, m is the Poisson coefficient, n is a number of dislocations in the pileup, s is an applied stress. Consequently, at the same applied stress s, the number of dislocations n in the pileup increases and, correspondingly, the stress sl on the leading dislocation in the pileup, sl = ns, increases, which assists an earlier opening of microcracks because of slip being blocked by Lomer-Cottrell barriers formed by dislocations at their intersection (see Fig. 3.32). In this relation, it is relevant to note that interstitial atoms affect the mobility of dislocations in two ways (see [101]). If they are essentially immobile, they always pin the dislocations. If they are sufficiently mobile to follow the dislocations in the course of deformation, the mobility of dislocations depends on the effect of interstitials on the electron structure, i.e. the mobility increases with an increasing concentration of free electrons and decreases in the opposite case.

64

3 Structure

Fig. 3.32 Pileups of gliding dislocations in two intersecting (111) planes in the fcc crystal lattice. Each dislocation is split, i.e. consists of a stacking fault encased between two partial dislocations (not shown in the Figure). An immobile sessile dislocation is formed due to a reaction between two gliding dislocations at the point of their intersection. This sessile dislocation and two neighbouring split gliding dislocations form a Lomer-Cottrell barrier. It is also immobile and blocks dislocation slip on both intersecting planes

This is why alloying of austenitic steels with nitrogen increases their impact toughness at RT and down to a temperature which still provides conditions for relaxation of stresses caused by the Lomer-Cottrell barriers at the intersection of localized slip planes. Therefore, a pseudo-brittle fracture of austenitic nitrogen steels under impact loading occurs at temperatures at which the stresses caused by the nitrogen-intensified and nitrogen-localized slip cannot be relaxed. This is not the case for austenitic carbon steels, where dislocations are not split, the slip is not localized and the mobility of dislocations is not increased. For this reason, the impact toughness of austenitic carbon steels is much smaller at RT and decreases gradually with decreasing temperature. As alloying with carbon ? nitrogen enhances the concentration of free electrons and, consequently, the atomic ordering, it would be reasonable to expect a further enhancement of the pseudo-brittleness of austenitic C ? N steels. But at a given interstitial content the DBTT of C ? N steels is lower than that of N steels (see [15, 105], and Table A11). The same effect is achieved by alloying austenitic nitrogen steels with nickel (see e.g. [93]). We can suppose that the reason for this is the significant increase in the concentration of free electrons, either by carbon ? nitrogen [10, 25] or by nickel [1]. This increase raises the mobility of dislocations so much, that conditions for the relaxation of stresses due to dislocation slip are created before submicron cracks can open along the intersecting slip planes.

3.2 Structural Change by Loading

65

Fig. 3.33 Localised slip in a radial-axial metallographic section of CN96 at the entrance side near the hole, LOM. a Steps of deformation and ASB. b Band of localized slip. c Cracks in such a band

3.2.3.2 Ballistic Impact The joint addition of C ? N to stainless austenitic CrMn steels improves the specific tensile fracture energy Ws (Table A1) composed of strength and ductility to the highest level measured so far (see also [14]). The aim of this Section is to find out if the high energy consumption measured in standard tensile tests can be retained in case of high velocity impact loading and to analyse structural changes by the ballistic impact (see about details [106]). Ballistic tests were carried out with steel CN96 and soft core standard ammunition 7.62 9 51 WK (NATO Level 1, STANAG 4569) at standard velocity VZ. = 835 m/s. The firing system was of type Steyr (20 m) in which the plane of a 10 mm thick disc was positioned perpendicular to the trajectory. All specimens were perforated and thus did not withstand the bullet. This implies that Vz was above the ballistic limit which marks the transition from penetration to perforation. The ballistic limit of the nitrogen steel Cr19Mn11N0.37 was measured at 535 m/s (see Sect. 6.2.3). Compared to this HNS the new HIS failed to meet the expectations. Light optical (LOM) and TEM were carried out to analyse the reason for this weak performance.

66

3 Structure

Fig. 3.34 Plates of e-martensite in the hardly deformed matrix, radial-axial section taken at a distance of about 10 mm from the hole

Steps of deformation were found on the entrance side of the bullet near the hole. The microstructure in between was marked by slip lines pointing to even deformation (Fig. 3.33a). The steps extended into adiabatic shear bands (ASB) in which a large shear displacement occurred while the material in between was much less deformed. In some ASB the shear was so high that cracks occurred (Fig. 3.33c). The phenomenon of ASB by joint localisation of slip and heat has been described for ballistic loading before (see e.g. [107]). The heat generated by slip has no time to dissipate (adiabatic), thus heating the slipped region which locally lowers the yield point and gives rise to more slip just there. At higher magnification an ASB of about 20 lm in width branches into thinner ones which are only a few lm wide (Fig. 3.33b). The inhomogeneous deformation is also visible on the exit side near the bullet hole where a net of deformation steps encases areas of less but even deformation. At half-thickness of the disk the hardness reached about 500 HV0.5 at the wall of the hole. The hardness penetration in radial direction was almost 10 mm. At a radial distance of 10 mm from the hole no signs of deformation were visible by LOM. However, occasional plates of e-martensite were revealed by TEM (Fig. 3.34). A high density of twins was observed at a distance of 6 mm from the hole (Fig. 3.35). About 3 mm away from the hole, e-martensite occurred together with an increased density of dislocations, Fig. 3.36. The dark field image of the emartensite is obtained in the light of reflection (210)e which is overlapped with ( 323)f reflection of f-Fe2N nitride (see the key diagram in Fig. 3.36d). The precipitates of f-Fe2N nitride are also lighted in Fig. 3.36c along with dislocations. At a distance of only 0.1 mm from the hole, localised slip without precipitates appeared (Fig. 3.37). This suggests that 3 mm away from the hole the temperature was locally high enough to provoke nitride precipitation while at a distance of 0.1 mm it was locally raised to the solution range.

3.2 Structural Change by Loading

67

Fig. 3.35 TEM structure in a radial-axial section taken at a distance of about 6 mm from the hole. a High density of twins in the austenite. b Diffraction pattern, zone axis [110]c

Fig. 3.36 TEM structure in a radial-axial section at a distance of about 3 mm from the hole. a Two systems of e-martensite and small f-Fe2N precipitates. b Diffraction pattern showing reflections of e-martensite, zone axis [121]e, and f-nitride, zone axis [133]f. c Dark field image of e-plates in the light of reflection ( 210)e and Fe2N nitride in reflection ( 323)f. d Key diagram of electron diffraction

68

3 Structure

Fig. 3.37 TEM structure in a tangential–axial section at a distance of about 0.1 mm from the hole, zone axis (111)c. Localized slip on intersecting planes in directions \110[ , no precipitates

Planar glide, twinning, e-martensite and an increase of the dislocation density were also observed after slow tensile straining of HIS steels (see Sect. 3.2.1.3). The really new deformation feature is the formation of adiabatic shear bands [107] combined with the precipitation of nitrides. Because of this localised slip, steel CN96 fails before the high fracture energy Ws, available at low strain rates, is spent. In Sect. 6.2.3 we discuss if an application seems feasible in spite of the perforation encountered.

3.2.4 Effect of Cyclic Loading The previous sections dealt with single loading and the effect of testing temperature, strain rate and stress state (notch) on structural changes. In the following the effect of cyclic loading is studied at room temperature. At first the structure is described after repeated impacts which are characterised by predominantly compressive stresses and a high velocity. Next, slow push/pull tests are carried out to investigate their effect on structure. 3.2.4.1 Repeated Impact The aim of this Section is to describe structural changes after repeated impact loading with the attempt to clarify the controlling mechanism of the wear resistance. The tests of impact wear are described in Sect. 4.2.2. The experimental

3.2 Structural Change by Loading

69

Fig. 3.38 Fragment of X-ray diffraction of steel CN96 after impact treatment and after electropolishing of the impact-treated surface. CuKa radiation

technique included the use of mineral particles of greywacke with a hardness of 760 HV0.1 and a grid size of 11–8 mm, which hit a plane steel surface perpendicularly at a velocity of 25 m/s. The mechanical response of steels to impact loading is usually described in terms of cold work hardening, however different reasons for strengthening are discussed. A short review of proposed mechanisms for cold work hardening was given in Sect. 3.2.1. Recent observations point also to an important role of impactinduced partial amorphisation and formation of a nanocrystalline structure in a thin surface layer [108, 109]. The surface of high interstitial steels was studied after impact tests by X-ray diffraction, back scattering Mössbauer spectroscopy and TEM (see about details [110]). Figure 3.38 shows the X-ray diffraction obtained from the very surface of the impact-treated steel CN96 and after subsequently removing a surface layer of 10 lm by electropolishing. It is seen that reflections of silicon oxide SiO2 stemming from greywacke debris disappear by electropolishing. The austenite remains the only phase and there is no sign of strain-induced c ? a transformation which is often claimed in studies of Hadfield type steel exposed to impact surface hardening. In case of a-phase, the {110} martensitic doublet would be located near the [111] austenitic reflection. Nevertheless, the Mössbauer spectrum obtained from the surface layer of about 150 nm using the conversion electron mode consists of a single line of paramagnetic austenite and Zeemann sextet of a ferromagnetic phase (Fig. 3.39). The surface ferromagnetism after impact treatment is not new and usually interpreted as a sign of a strain-induced martensitic c ? a transformation. Traces of Zeemann’s sextets are observed even at the depth of about 10 lm (Fig. 3.40), whereas, after further electropolishing, the spectrum consists only of the paramagnetic austenitic component. It is remarkable that the relative emission of electrons from the impact-treated surface is rather small (Fig. 3.39) because it is subdued by the implanted SiO2

70

3 Structure

Fig. 3.39 Mössbauer spectrum obtained from the impact-treated surface of steel CN96 in the mode of conversion electrons. Fec is an austenitic component. Fe0 and Fe1 belong to Fe atoms in the martensite having no and one foreign atom as nearest neighbours

Fig. 3.40 As Fig. 3.39 but 10 lm below the impacttreated surface

Fig. 3.41 Mössbauer spectrum obtained from the impact-treated surface of Hadfield steel MnC in the mode of conversion electrons. Fecl component belongs to iron atoms located in clusters enriched by Mn and C

debris but increases after polishing (Fig. 3.40). Qualitatively the same spectra were obtained from the impact-treated surface of steel MnC (Fig. 3.41) and steel CN107. The comparison of Mössbauer results with the data of X-ray diffraction suggests that the surface ferromagnetism of the studied impact treated steels is not concerned with a-martensite.

3.2 Structural Change by Loading

71

Fig. 3.42 Structure of the surface of steel CN96 after impact treatment. a Amorphous state. b Twin structure of austenite

Fig. 3.43 Structure of the impact-treated steels. a CN96 at a distance of 5 lm. b CN107, 15 lm under the surface

TEM observations of the surface layer after impact treatment of CN steels and Hadfield steel revealed a complicated picture which is characterized by three types of substructure: an amorphous state, nanocrystals in an amorphous matrix and a strongly deformed twinned structure of austenite. Figures 3.42–3.44 demonstrate examples of the observed structures in steels CN96 and CN107. The very surface layer consists of an amorphous structure (Fig. 3.42a) with rare islands of heavily twinned austenite (Fig. 3.42b). The amorphous structure is observed in steel CN96 at a distance of 5 lm below the surface along with nanocrystals witnessed by separate reflections at the Debye ring (Fig. 3.43a). In contrast, a preferentially amorphous state occurs in steel CN107 at a distance of about 15 lm (Fig. 3.43b). It can be concluded from this comparison that the higher

72

3 Structure

Fig. 3.44 Structure of impact-treated steels at a distance of 45 lm under the surface. a CN96 and b CN107

carbon content assists the formation of an amorphous structure at a larger distance from the surface. The same tendency is observed after removing a layer of 45 lm: the structure is characterized by dislocation slip bands and rather high density of tangled dislocations in steel CN96 (Fig. 3.44a), whereas amorphous areas still can be found in CN107 (Fig. 3.44b). In consistency with X-ray diffraction data, TEM studies also confirm the absence of a-martensite. Therefore, the ferromagnetism of the impact-treated surface layer is obviously related with its amorphous state, as in amorphous ironbased ribbons (see e.g. [19]). A physical reason for the formation of an amorphous structure in austenitic steels with carbon or carbon ? nitrogen under heavy impact treatment needs a special discussion because it does not occur in similarly treated austenitic steels free of interstitials (see e.g. [111]). This phenomenon can be explained in terms of an increased concentration of vacancies in the interstitial solid solution. Smirnov [112] was the first to theoretically predict that the dissolution of interstitial atoms in metals increases the thermodynamic equilibrium concentration of vacancies. Later on, the formation of superabundant vacancies in Ni, Pd and Fe hydrides under a high hydrogen pressure has been proven by Fukai et al. [113, 114]. An increase in the concentration of vacancies caused by dissolution of hydrogen in a stainless austenitic steel was observed in [115]. Even earlier, McLellan [116] has shown that the enthalpy of vacancy formation is decreased in the vicinity of carbon atoms in the iron austenite. The increased concentration of vacancies eases the formation of an amorphous state. From this consideration, it is also understandable why the alloys intended for obtaining the amorphous ribbons by rapid quenching from the liquid state usually contain much interstitial boron, e.g. Fe80Si4B14.

3.2 Structural Change by Loading

73

Fig. 3.45 Cross-head displacement and specimen temperature in tension/compression tests of CN96 at ra = ±450 MPa in dependence of the number of cycles. Specimen 1 was run for N = 1000, specimen 2 for N = 2300 and specimen 3 until fracture at Nf = 3800. The Vickers hardness HV30 is given in brackets: 1, 2 were measured on the specimen surface, 3 on the fatigue fracture face

Thus, based on the obtained results, one can conclude that the mechanism of surface hardening by impact treatment of CN steels, as well as of Hadfield steel, is based on the formation of a complex amorphous, nanocrystalline and thin-twinned austenitic structure and is not related with a martensitic transformation in the surface.

3.2.4.2 Push/Pull Fatigue Fatigue tests allow to investigate what fraction of the tensile strength may be retained after cyclic loading. Several criteria are used to describe fatigue properties: fatigue limit, fatigue life, fatigue crack growth rate. A detailed analysis of fatigue curves measured in the coordinates of stress amplitude/number of cycles is given in Sect. 4.1.7. Described here is the substructure of steel CN96 developed at different stages of the fatigue curve obtained in tension/compression tests. A sample of steel CN96 was run at a stress amplitude of ra = ±450 MPa until its fracture at Nf = 3800 (Fig. 3.45). The states 1–3, corresponding to different parts of the fatigue curve, were taken for TEM studies of the structure at the surface and in the centre of the tested specimens. The initial decrease in the displacement of the cross-heads indicates cyclic hardening. In the final part the displacement increases because of crack initiation and growth. This observation of cyclic hardening in the CrMnCN steel is at some variance with the data of Sun et al. [117] for austenitic CrMnN steels, where the absence of hardening was shown in the whole range of fatigue life. Generally, cyclic softening prevails if cross slip of dislocations is prevented, which assists the

74

3 Structure

Fig. 3.46 Structure of steel CN96 after 1000 cycles (state 1 in Fig. 3.45, the sample is taken from the specimen surface). a TEM image. b Electron diffraction with reflections of austenite and orthorhombic nitride f-Fe2N; zone axes ½0 11c and ½5 4 21

formation of planar dislocation arrays and reversible gliding of dislocations within narrow slip bands [118, 119]. Such a dislocation structure is just typical for the nitrogen and nitrogen ? carbon austenitic CrMn steels. As the fatigue of austenitic nitrogen CrNi steels is concerned, it depends on the value of strain amplitudes. At low strain amplitudes, planar slip and, correspondingly, softening occurs within the whole fatigue life, whereas, if the strain amplitude increases, planar slip with hardening occurs within 20–100 cycles and thereafter is substituted by dislocation subcell structures and softening [120–122]. Thus, despite a similar dislocation distribution, the effect of alloying C ? N on the fatigue of austenitic CrMn steels is a bit different from that of N. The following structural change accompanies the fatigue tests of which results are presented in Fig. 3.46. The initial structure of the tested steel CN96 is characterized by a fully austenitic state and a moderate density of dislocations of about 1010 cm-2. After 1000 cycles (state 1 in Fig. 3.45), the density of dislocations was increased to between 3 9 1010 and 5 9 1010 cm-2 in the specimen centre and between 5 9 1010 and 7 9 1010 cm-2 at the surface. Planar slip is clearly observed. Twinning was also occasionally met in one or two systems, correspondingly. The most distinctive feature is the precipitation of fine iron nitride f-Fe2N particles, as identified by diffraction (Fig. 3.46b). After 2300 cycles (state 2 in Fig. 3.45), the structure becomes more or less homogeneous in the cross section of the sample. The density of dislocations increases up to 1011 cm-2. Along with nitride precipitates in the austenite, new observations are the c ? e transformation in one (Fig. 3.47) or two (Fig. 3.48) intersecting systems and the precipitation of cementite within the e-martensite plates.

3.2 Structural Change by Loading

75

Fig. 3.47 Structure of steel CN96 after 2300 cycles at ra = ±450 MPa (state 2 in Fig. 3.45). a TEM image of austenite with nitride precipitates and a plate of e-martensite containing cementite precipitates. b Reflections of h-cementite and e-martensite in the diffraction pattern obtained from the e-plate; zone axes ½2 1 2e and ½ 2 54h

As expected (see, e.g., [122]), a-martensite is formed at the intersection of eplates and here the cementite precipitates of 20–50 nm in size are larger than in the e-martensite plates. Precipitates of f–Fe2N nitride, 2–3 nm in size, are dispersed in the austenite. After fracture at 3800 cycles (state 3), some areas depleted of dislocations appear in the structure along with those of a high density in the range of 1011 cm-2. Again f–Fe2N nitrides are observed in the austenite (Fig. 3.49). The structure presented above after tension/compression tests resembles in many features that developed during unidirectional loading: dislocations, twins and e-plates, all increasing their concentration as the cyclic straining proceeds. But at the same time, the precipitation of Fe2N in austenite and of Fe3C in emartensite and in a-islands transformed at the intersection of e-plates are unique features of the CrMnCN steel after cyclic loading. It is natural that, because of a smaller carbon solubility in a bcc lattice, the cementite precipitates in the a-phase are coarser than in e. Of course, these a-islands are too small and scarce to affect the fatigue life. It is a curious question, why cementite is precipitated at the e-martensite plates, whereas the nitride precipitates are nearly evenly distributed in the austenitic bulk. One should note in this relation that carbon atoms have a strong affinity to grain and sub-boundaries in austenite and ferrite [123–125]. Therefore, they are preferentially segregated at stacking faults which are prerequisites for the formation of e-plates in the deformed austenite. In contrast, nitrogen atoms tend to segregate much less at interfaces [123, 124, 126]. However, causing larger distortions in the

76

3 Structure

Fig. 3.48 State 2 continued. a TEM image of intersecting e-plates formed in different crystallographic systems. b Diffraction pattern from the intersection site showing austenite, a-martensite and cementite; zone axes [115]a and ½2 5 4h . c Cementite particles in the a- and ephases shown in a dark field image obtained by reflection ½ 20 1h

austenite and having a larger enthalpy of binding to dislocations, they are more able to form clouds around dislocations [68]. For this reason, they precipitate preferentially within the austenitic solid solution and, due to their small size, add to the cyclic strength. Why Fe3N is precipitated during aging after unidirectional strain (see e.g. [127]), but Fe2N during cyclic loading, remains unclarified. Cyclic recovery may enhance the annihilation of dislocations of opposite Burgers vector merging their interstitial clouds, thus raising the local concentration of nitrogen atoms, which would favour Fe2N. After accumulating a critical cyclic strain, cracks are initiated which ‘‘soften’’ the specimen and lead to an increase of the displacement until failure, while the hardness still increases. However, the dislocation structure breaks down locally (Fig. 3.49) indicating the start of cyclic softening of the material.

3.2 Structural Change by Loading

77

Fig. 3.49 Structure of steel CN96 fractured after 3800 tension/compression cycles at ra = ±450 MPa (specimen 3 in Fig. 3.45). a TEM image showing an uneven distribution of dislocations and precipitates of f-nitride Fe2N. b Diffraction pattern; zone axes [112]c and [-11.5 9 -4]f

Based on work by Feltner and Laird [128], Lukas and Klesnil [129] reviewed the microstructural features of substitutional fcc Cu-, Al- and Fe-alloys after cycling in dependence of the stacking fault energy c and the cycles to failure Nf. They defined the areas A, B and C in their plot of c over Nf. As to the structure, steel CN96 belongs to area C comprising planar slip, twinning and e-martensite. However, its stacking fault energy of 40 mJ/m2 is too high to fit into the scheme. The high interstitial content of this new alloy affects both, c and Nf. In his study on the cyclic deformation response of substitutional Cu-alloys, Wang [130] proposed a criterion for the transition from wavy to planar slip based on the interrelation of c and the concentration of free electrons per atom e/a. The present high interstitial steels belong to the planar slip type, but in contrast to the respective Cu-alloys their c is inversely proportional to e/a (see Fig. 3.19). Both results seem to indicate a difference between interstitial and substitutional fcc solid solutions in respect to their response to cyclic loading.

References 1. Shanina BD, Gavriljuk VG, Konchitz AA, Kolesnik SP (1998) The influence of substitutional atoms upon the electron structure of the iron-based transition metal alloys. J Phys Condens Matter 10:1825–1838 2. Hohenberg H, Kohn W (1964) Inhomogeneous electron gas. Phys Rev B 136(3):864–871 3. Kohn W, Sham LJ (1965) Self-consistent equations including exchange and correlation effects. Phys Rev A 140(4):1133–1138

78

3 Structure

4. Blaha P, Schwarz K, Madsen GKH, Kvasnicka D, Luitz J (2001) Wien2 k, an augmented plane wave ? local orbitals program for calculating crystal properties. Karlheinz Schwarz Techn. Universität, Wien, Austria, ISBN 3-9501031-1-2 } 5. Seith W, Kubaschewski O (1935) Die elektrolytische Uberführung von Kohlenstoff in festen Stahl. Zs Electrochem 41(7b):551–558 } 6. Seith W, Daur Th (1938) Elektrische Uberführung in festen Metallegierungen. Zs Electrochem 44(4):256–260 7. Gavriljuk VG, Shanina BD, Baran NP, Maximenko VM (1993) Electron-spin-resonance study of electron properties in nitrogen and carbon austenite. Phys Rev B 48(5):3224–3231 8. Shanina BD, Kolesnik SP, Konchitz AA, Gavriljuk VG, Smouk SY, Tarasenko AV (1994) The influence of nitrogen on the paramagnetic properties of the multicomponent d-element iron-based alloy. Solid State Commn 90(2):109–113 9. Shanina BD, Gavriljuk VG, Konchitz AA, Kolesnik SP, Tarasenko AV (1995) Exchange interaction between electron subsystems in iron-based F.C.C. alloys doped by nitrogen or carbon. Physica Stat Sol (a) 149:711–722 10. Shanina B, Gavriljuk V, Berns H, Schmalt F (2002) Concept of a new high-strength lowcost stainless steel. Steel Res 73(3):105–113 11. Gavriljuk V, Rawers J, Shanina B, Berns H (2003) Nitrogen and Carbon in Austenitic and Martensitic steels: atomic interactions and structural stability. Mater Sci Forum 426– 432:943–950 12. Shanina BD, Gavriljuk VG (2004) Effect of carbon and nitrogen on electronic structure of steel. J Steel Relat Mater suppl ‘‘High Nitrogen Steels 2004’’, pp 45–52 13. Shanina BD, Gavriljuk VG, Berns H (2007) Atomic interactions in stainless austenitic CrMn steels alloyed with C, N or (C ? N). Mater Sci Forum 539–543:4993–4998 14. Berns Hans, Gavriljuk Valentin G (2007) Steel of highest fracture energy. Key Eng Mater 345–346:421–424 15. Berns H, Gavriljuk VG, Riedner S, Tyshchenko A (2007) High strength stainless austenitic CrMnCN steels—part I: alloy design and properties. Steel Res Intern 78(9):710–715 16. Gavriljuk VG, Razumov O, Petrov Yu, Surzhenko I, Berns H (2007) High strength stainless austenitic CrMnCN steels—part II: structural changes by repeated impacts. Steel Res Intern 78(9):716–719 17. Shanina BD, Gavriljuk VG, Berns H (2007) High strength stainless austenitic CrMnCN steels—part III: electronic properties. Steel Res Intern 78(9):720–724 18. Gavriljuk VG, Shanina BD, Berns H (2008) Ab initio development of a high-strength corrosion-resistant austenitic steel. Acta Mater 56(18):5071–5082 19. Mogilny GS, Shanina BD, Maslov VV, Nosenko VK, Shevchenko AD, Gavriljuk VG (2011) Structure and magnetic properties of rapidly quenched FeSiB ribbons. J NonCrystalline Solids 357(16–17):3237–3244 20. Sozinov AL, Balanyuk AG, Gavriljuk VG (1997) C–C interaction in iron-base austenite and interpretation of Mössbauer spectra. Acta Mater 45(1):225–232 21. Sozinov AL, Balanyuk AG, Gavriljuk VG (1999) N–N interaction and nitrogen activity in the iron base austenite. Acta Mater 47(3):927–935 22. Balanyuk AG, Gavriljuk VG, Shivanyuk VN, Tyshchenko AI, Rawers J (2000) Mössbauer study and thermodynamic modeling of Fe-C-N alloy. Acta Mater 48(15):3813–3821 23. Rawers JC (1999) High carbon-nitrogen iron alloys. J Mater Sci 34(5):941–944 24. Gavriljuk VG, Tarasenko AV, Tyshchenko AI (2000) Low temperature ageing of the freshly formed Fe-C and Fe-N martensites. Scripta Mater 43(3):233–238 25. Shanina BD, Tyshchenko AI, Glavatskyy IN, Runov VV, Petrov YuN, Berns H, Gavriljuk VG (2011) Chemical nano-scale homogeneity of austenitic CrMnCN steels in relation to electronic and magnetic properties. J Mater Sci 46(24):7725–7736 26. Vonsovsky VS (1971) Magnetism (in Russian). Nauka, Moscow 27. Noskova NI, Pavlov VA, Nemnonov SA (1965) A correlation between the stacking fault energy and structure of metals (in Russian). Physics Metals Metallogr 20(6):920–924

References

79

28. Ruff AW Jr (1970) Measurement of stacking fault energy from dislocation interactions. Metall Trans A 1(9):2391–2413 29. Garner FA, McCarthy JM (1990) Spinodal-like decomposition of Fe-Ni and Fe-Ni-Cr ‘‘invar’’ alloys during neutron or ion irradiation. In: Russel KS, Smith DF (eds) Physical metallurgy of controlled expansion invar-type alloys. TMS-AIME, Warendale, pp 187–206 30. Rotman F, Gilbon D, Lorant H, Dimitrov O (1987) Radial redistribution of elements in austenitic Fe-Cr-Ni based alloys under HVEM Irradiation. Mater Sci Forum 15–18:1409– 1414 31. Gavriljuk V, Petrov Yu, Shanina B (2006) Effect of nitrogen on the electron structure and stacking fault energy in austenitic steels. Scripta Mater 55(6):537–540 32. Gavriljuk VG, Shanina BD, Berns H (2000) On the correlation between electron structure and short range atomic order in iron-based alloys. Acta Mater 48(15):3879–3893 33. Gavriljuk VG, Tyshchenko AI, Bliznuk VV, Yakovleva IL, Riedner S, Berns H (2008) Cold work hardening of high-strength austenitic steels. Steel Res Intern 79(6):413–422 34. Hadfield RA (1925) Metallurgy and its influence on modern progress, with a survey of education and research. Chapman and Hall Ltd, London, pp 91–101 35. Hall JH (1929) Trans AIME 84:382–427 36. Krivobok VN (1929) Investigation on the microstructure of Hadfield manganese steel. Trans Amer Soc Steel Treat 15(6):893–928 37. Berns H, Klärner HF, Schmidtmann E (1967) Einfluß der Wärmebehandlung und Legierung auf die Kaltverfestigung von Manganhartstahl. Arch. Eisenhüttenwesen 38(7):547–553 38. White CH, Honeycombe RWK (1962) Structural changes during deformation of high-purity Fe-Mn-C alloys. JISI 200(4):457–466 39. Colette G, Crussard C, Kohn A, Plateau J, Pomey G, Weisz M (1957) Contribution à l’étude des transformations des austénites à 12 % Mn. Revue de Metallurgie 6:433–481 40. Janssen K, Jellinghaus W (1956) Perlit- und Martensitbildung in Mangan-Hartstahl. Arch. Eisenhüttenwesen 27(4):573–578 41. Berns H (1968) Entkohlung im Manganhartstahl Z. Wirtschaftl. Fertigung 63(9):437–441 42. Raghaven KS, Sastri SA, Marcinkowsky MJ (1969) Nature of work hardening in Hadfield steel. Trans TMS-AIME 245:1569–1575 43. Dastur YN, Leslie WC (1981) Mechanism of work hardening in Hadfield manganese steel. Metall Trans A 12(5):749–759 44. Matosyan MA, Golikov VM (1970) Zashitniye Pokritiya na Metallakh (in Russian). Kiev Naukova Dumka 3:58–59 45. Karaman I, Huseyin S, Ken G (1998) On the deformation mechanisms in single crystal hadfield manganese steels. Scripta Mater 38(6):1009–1015 46. Tsakiris V, Edmonds DV (1999) Martensite and deformation twinning in austenitic steels. Mater Sci Eng A 273–275:430–436 47. Uggowitzer PJ, Harzenmoser M (1989) Strengthening of austenitic steels by nitrogen. In: Foct J, Hendry A (eds) High nitrogen steels, HNS 88:174–79. Institute of Metals, London 48. Gavriljuk VG, Duz’ VA, Yephimenko SP (1990) Structure and mechanical properties of cold-worked high-nitrogen austenite. In: Stein G, Witulsky H (eds) High nitrogen steels, HNS 90, pp 100–103. Stahl&Eisen, Düsseldorf 49. Müllner P, Solenthaler C, Ugowitzer P, Speidel MO (1993) Martensite and deformation twinning in austenitic steels. Mater Sci Eng A 164:164–169 50. Müllner P, Solenthaler C, Ugowitzer P, Speidel MO (1994) Brittle fracture in austenitic steel. Acta Metal Mater 42(7):2211–2217 51. Müllner P, Solenthaler C, Speidel MO (1994) Second order twinning in austenitic steel. Acta Metal Mater 42(5):1727–1732 52. Kubota S, Xia Y, Tomota Y (1998) Work-hardening behavior and evolution of dislocationmicrostructures in high-nitrogen bearing austenitic steels. ISIJ Intern 38(5):474–481 53. Saller G, Spiradek-Hahn K, Sgheu C, Clemens H (2006) Microstructural evolution of Cr– Mn–N austenitic steels during cold work hardening. Mater Sci Eng A 427:246–254

80

3 Structure

54. Copley SM, Kear BH (1968) The dependence of the width of a dissociated dislocation on dislocation velocity. Acta Metal 16(2):227–231 55. Kestenbach HJ (1977) The effect of applied stress on partial dislocation separation and dislocation substructure in austenitic stainless steels. Phil Mag 36(6):1509–1515 56. Lee T-H, Oh C-S, Kim S-J, Takaki S (2007) Brittle fracture in austenitic steel. Acta Mater 55(11):3649–3662 57. Dorn John E, Raynak Stanley (1964) Nucleation of kink pairs and the Peierls mechanism of plastic deformation. TMS AIME 230(5):1052–1064 58. Cottrell AH, Bilby BA (1949) Dislocation theory of yielding and strain ageing of iron. Proc Phys Soc A 62(1):49–62 59. Keh AS (1963) In: Relation between structural and mechanical properties of metals. H. M. Stationery Office, London, p 436 60. Baird JD (1963) Strain aging of steel-A critical review. Iron Steel 36:450–457 61. Hirth JP, Lothe J (1968) Theory of dislocations, Ch. 18, McGraw-Hill Book Co., New York, pp 433–472 62. Vitek V, Kroupa F (1966) Dislocation theory of slip geometry and temperature dependence of flow stress in B.C.C. metals. Phys Stat Sol a 18(2):703–713 63. Grujicic M (1994) The core structure of (a/2) \ 110 [ screw dislocations in Fe–Ni–Cr–N– austenite. Mater Sci Eng A 183:223–232 64. Nyilas A, Obst B, Nakajima H (1993) Tensile properties, fracture and crack growth of a nitrogen strengthened new stainless steel (Fe-25Cr-15Ni-0.35N) for cryogenic use. In: Gavriljuk VG, Nadutov VM (eds) High nitrogen steels, HNS 93, Institute for Metal Physics, Kiev, pp 339–344 65. Gavriljuk VG, Sozinov AL, Foct J, Petrov YuN, Polushkin YuA (1998) Effect of nitrogen on the temperature dependence of the yield strength of austenitic steels. Acta Mater 46(3):1157–1163 66. Seeger A (1955) The generation of lattice defects by moving dislocations, and its application to the temperature dependence of the flow stress of fcc crystals. Phil Mag 46(382):1194–121 67. Gavriljuk VG, Berns H, Escher Ch, Glavatskaya NI, Sozinov AL, Petrov YuN (1999) Grain boundary strengthening in austenitic nitrogen steels. Mater Sci Eng A 271:14–21 68. Gavriljuk VG, Duz’ VA, Yephimenko SP, Kvasnevski OG (1987) Interaction between carbon/nitrogen atoms and dislocations in austenite (In Russian). Phys Metals Metallogr 64(6):1132–1135 69. Chen SR, Tang D (1990) Effect of interstitial atom concentration on lattice parameters of martensite and retained austenite in iron-carbon-nitrogen alloys. Mater Sci Forum 56– 58:201–206 70. Gavriljuk VG, Berns H (1999) High Nitrogen steels. Springer, Berlin 71. Norström LA (1977) The influence of nitrogen and grain size on yield strength in type AISI 316L austenitic stainless steel. Metal Sci 11(6):208–112 72. Degallaix S, Foct J, Hendry A (1986) Mechanical behaviour of high-nitrogen stainless steels. Mater Sci Technol 2(9):946–950 73. Varin RA, Kurjidlowski KJ (1988) The effects of nitrogen content and twin boundaries on the yield strength of various commercial heats of type 316 austenitic stainless steel. Mater Sci Eng A 101:221–226 74. Werner E (1988) Solid solution and grain size hardening of nitrogen-alloyed austenitic steels. Mater Sci Eng A 101:93–98 75. Uggowitzer PJ, Speidel MO (1990) Ultrahigh-strength austenitic steels. In: Stein G, Witulsky H (eds) High Nitrogen Steels, HNS 90. Stahl & Eisen, Düsseldorf, pp 156–160 76. Obst B, Nyilas A (1991) Ultrahigh-strength austenitic steels. Mater Sci & Eng A137:141– 150 77. Adler PH, Olson GB, Owen WS (1986) Strain hardening of Hadfield manganese steel. Metal Trans A 17(10):1725–1737

References

81

78. Petrov YuN (1978) Defects and diffusionless transformation in steel (in Russian). Naukova dumka, Kiev 79. Schramm RE, Reed RP (1975) Stacking fault energies of seven commercial austenitic stainless steels. Metall Trans A 6(7):1345–1351 80. Riedner S (2009) Höchstfeste nichtrostende austenitische CrMn-Stähle mit (C ? N). Doctor thesis. Ruhr Universität Bochum, Lehrstuhl Werkstofftechnik, Bochum 81. Seeger A (1954) The temperature dependence of the critical shear stress and of workhardening of metal crystals. Phil Mag 45(366):771–773 82. Akdut N, Keichel J, Foct J (1997) The influence of nitrogen and orientation on the rolling deformation mechanisms of austenitic single crystals. Steel Res 68(11):495–500 83. Speidel HJC, Speidel M (2003) Nickel and chromium based high nitrogen alloys. In: Speidel MO, Kowanda C, Diener M (eds) High Nitrogen steels, HNS 2003. Institute of Metallurgy Swiss Federal Institute of Technology, Zürich, pp 101–112 84. Sassen J, Carratt-Reed AJ, Owen WS (1989) Electron microscopy of austenitic Fe-Ni-Cr alloys containing nitrogen. In: Foct J, Hendry A (eds) High Nitrogen steels., HNS 88Institute of Metals, London, pp 159–162 85. Sandström R, Bergqvist H (1977) Temperature dependence of tensile properties and strengthening of nitrogen alloyed austenitic stainless steels. Scand J Metallurgy 6:156–169 86. Nilsson J-O, Thorvaldsson T (1985) The influence of nitrogen on microstructure and strength of a high-alloy austenitic stainless steel. Scand J Metallurgy 15:83–89 87. Byrnes MLG, Grujicic M, Owen WS (1987) Nitrogen strengthening of a stable austenitic stainless steel. Acta Metall 35(7):1853–1862 88. Reed RP, Simon NJ (1989) Nitrogen strengthening of austenitic stainless steels at low temperatures. In: Foct J, Hendry A (eds) High Nitrogen steels, HNS 88. Institute of Metals, London, pp 180–188 89. Grujicic M, Owen WS (1995) Models of short range order in a face-centered cubic Fe-Ni-Cr alloy with a high concentration of nitrogen. Acta Metall Mater 43(11):4201–4211 90. Grujicic M (1995) The effect of nitrogen on the structure and mobility of dislocations in FeNi-Cr austenite. J Mater Sci 30(22):5799–5807 91. Obst B (1998) Basic aspects of tensile properties. In: Seeber B (ed) Handbook of applied superconductivity, 2, F1.11, Institute of Physics Publishing, Bristol-Philadelphia, pp 969– 993 92. Schmalt F (2005) Nutzung der Löslichkeit von C ? N in nichtrostenden Stählen. Doctor thesis. Ruhr Universität Bochum, Lehrstuhl Werkstofftechnik, Bochum-Düsseldorf 93. Tomota Y, Nakano J, Xia Y, Inoue K (1998) Unusual strain rate dependence of low temperature fracture behaviour in high nitrogen bearing austenitic steels. Acta Mater 46(9):3099–3108 94. Tobler RL, Meyn D (1988) Cleavage-like fracture along slip planes in Fe-18Cr-3Ni-13Mn0.37 N austenitic stainless steel at liquid helium temperature. Metall Trans A 19(6):1626– 1631 95. Harzenmoser MAE (1990) Massive aufgestickte austenitisch-rostfreie Stähle und Duplexstähle. Doctoral thesis. Eidgenössische Technische Hochschule, Zürich 96. Ishizaki J, Orita K, Terao K (1992) The influence of chemical composition and pre-strain on impact toughness transition behaviour of 18 %MN-18 %Cr-N austenitic steels. J ISIJ (Tetsu to Hagane) 78(12):1846–1853 97. Uggowitzer PJ, Paulus N, Speidel MO (1992) Ductile-to-brittle transition in nitrogen alloyed austenitic stainles steels. In: Nordberg H, Börklund J (eds) Application of stainless steels’92. The Institute of Metals ASM Intern, Stockholm, pp 62–72 98. Tomota Y, Endo S (1990) Cleavage-like fracture at low temperatures in an 18Mn-18Cr0.5 N austenitic steel. ISIJ Intern 30(8):656–662 99. Müllner P (1997) On the ductile to brittle transition in austenitic steel. Mater Sci Eng A 234–236:94–97 100. Tomota Y, Xia Y, Inoue K (1998) Mechanism of low temperature brittle fracture in high nitrogen bearing austenitic steels. Acta Mater 46(5):1577–1587

82

3 Structure

101. Gavriljuk VG, Shivanyuk VN, Shanina BD (2005) Change in the electron structure caused by C, N and H atoms in iron and its effect on their interaction with dislocations. Acta Mater 53(19):5017–5024 102. Larikov LN, Falchenko VM, Mazanko VF, Gurevich SM, Kharchenko GI, Ignatenko AI (1975) Abnormal acceleration of diffusion during impulse loading of metals. Rep Acad Sci USSR (in Russian) 221:1073–1075 103. Pogorelov AE, Ryaboshapka KP, Zhuravlev AF (2002) Mass transfer mechanism in real crystals by pulsed laser irradiation. J Appl Phys 92(10):5766–5771 104. Karnaukhov IN, Pogorelov AE, Chernolevsky MS (2006) Non-activated mechanism of mass transfer by moving dislocations (in Russian). Metallofiz Noveishie Technol 28(6):827–835 105. Bernauer J (2004) Einfluß von Kohlenstoff als Legierungselement in stickstofflegierten Chrom-Mangan-Stählen. Doctor thesis. Eidgenössische Technische Hochschule, Zürich 106. Berns H, Riedner S, Gavriljuk V, Petrov Y, Weihrauch A (2011) Microstructural changes in high interstitial stainless austenitic steels due to ballistic impact. Mater Sci Eng A 528(13– 14):4669–4675 107. Nakkalil R (1991) Formation of adiabatic shear bands in eutectoid steels in high strain rate compression. Acta Metal Mater 39(11):2553–2563 108. Xu Y, Chen Y, Xiong J, Zhu J (2001) Acta Metal Sinica 37(2):165–170 109. Petrov YN, Gavriljuk VG, Berns H, Schmalt F (2006) Surface structure of stainless and Hadfield steel after impact wear. Wear 409(6):687–691 110. Gavriljuk VG, Tyshchenko AI, Razumov ON, Petrov YuN, Shanina BD, Berns H (2006) Corrosion-resistant analogue of Hadfield steel. Mater Sci Eng A 420(1–2):47–54 111. Prokopenko GI, Petrov YuN, Vasilyev MA, Trophimova LN, Bliznyuk VV (2008) Structural-phase transformations in austenitic steel under ultrasonic impact surface treartment (in Russian). Metal Phys Adv Technol 30(1):115–131 112. Smirnov AA (1991) Theory of vacancies in interstitial alloys (in Russian). Metal Phys 13(9):40–44 113. Fukai Y, Okuma N (1993) Evidence of copious vacancy formation in Ni and Pd under a high hydrogen pressure. Jpn J Appl Phys part 2, 32(9A):L1256–1259 114. Fukai Y (2003) Formation of superabundant vacancies in M-H alloys and some of its consequences: a review. J Alloys Compd 356–357:263–269 115. Gavriljuk VG, Bugaev VN, Petrov YuN, Tarasenko AV, Yanchitsky BZ (1996) Hydrogeninduced equilibrium vacancies in FCC iron-base alloys. Scripta Mater 34(6):903–907 116. McLellan RB (1988) The thermodynamics of interstitial-vacancy interactions in solid solutions. J Phys Chem Solids 49(10):1213–1217 117. Sun H, Diener M, Uggowitzer PJ, Speidel MO (1990). Low cycle fatigue behaviour of high nitrogen steels. In: Stein G, Witulski H (eds) High Nitrogen steels, HNS 90. Stahl & Eisen, Düsseldorf, pp 220–223 118. Grosskreutz JC (1972) Strengthening and fracture in fatigue (Approaches for achieving high fatigue strength). Metall Trans 3(5):1255–1262 119. Margolin H, Mahajan Y, Saleh Y (1976) Grain boundaries, stress gradients and fatigue crack initiation. Scripta Metall 10(12):1115–1118 120. Taillard R, Foct J (1989) Mechanisms of the action of nitrogen interstitials upon low cicle fatigue behaviour of 316 stainless steels. In: Foct J, Hendry A (eds) High Nitrogen Steels, HNS 88. The Institute of Metals, London, pp 387–391 121. Degallaix S, Dickson JI, Foct J (1989) Effect of nitrogen on fatique and creep-fatigue behaviour of austenitic stainless steels. In: Foct J, Hendry A (eds) High Nitrogen Steels, HNS 88. Institute of Metals, London, pp 380–386 122. Venables JA (1962) The martensite transformation in stainless steel. Phil Mag 7(1):35–44 123. Rudy ML, Huggins RA (1966) Grain boundary segregation and the cold work peak in iron containing carbon and nitrogen. TMS AIME 236(12):1662–1666 124. Lagerberg G, Josefsson A (1955) Influence of grain boundaries on the behaviour of carbon and nitrogen in a-iron. Acta Metal 3(5):236–244

References

83

125. Petrov YuN (1993) On the carbon distribution at structural imperfections in manganese austenite. Scripta Metall et Mater 29(11):1471–1476 126. Petrov YuN, Gavriljuk VG, Berns H, Escher Ch (1999) Nitrogen partitioning between matrix, grain boundaries and precipitates in high-alloyed austenitic steels. Scripta Mater 40(6):669–674 127. Gavriljuk VG, Stein G, Berns H (2003) Structural stability of austenitic CrMn(Mo)N steels for high-strength, corrosion-resistant retaining rings. Steel Res Int 74(7):444–452 128. Feltner CE, Laird C (1968) Factors influencing the dislocation structure in fatigue metals. TMS AIME 242(8):1253–1257 129. Lukas P, Klesnil M (1972) In: Devereux O, Owen F (eds) National Association of Corrosion Engineers. Houston, pp 118–132 130. Wang Z (2004) Cyclic deformation response of planar-slip materials and a new criterion for the wavy-to-planar-slip transition. Phil Mag 84:351–379

Chapter 4

Properties

This chapter deals with key properties relevant in service. Those important for manufacturing are discussed in Chap. 5. The position of new HIS is revealed by comparing them to reference steels given in Table 2.1. The mechanical properties are in the center of attention, because steels of higher strength are to be developed. These stainless grades require a corrosion resistance comparable to some standard steels which has to be proven by respective tests. The high austenite stability of HIS poses the question, if they are useful for nonmagnetic applications. Last, not least, a higher strength of ductile austenite offers a chance to improve the resistance to impact wear and cavitation. See [1] for more details.

4.1 Mechanical Properties Tensile tests are carried out to characterise steels under slow uniform loading within a wide range of temperature. Notch impact tests come with a multiaxial stress state and increased velocity, which in connection with subzero testing temperatures offer insight into impending embrittlement. Fatigue is responsible for many failures and respective tests help to find out what fraction of tensile strength is retained under cyclic loading.

4.1.1 Tensile Properties at Room Temperature Specimens of Ø5 mm were tested according to EN10002-1 at room temperature (RT) and a rate of 0.5 mm/min. The engineering stress–strain curves of hot worked steels (longitudinal taking) are presented in Fig. 4.1a. From these, common properties of (i) proof, ultimate and true fracture strength (Rp0.2, Rm, R), (ii) ductility in H. Berns et al., High Interstitial Stainless Austenitic Steels, Engineering Materials, DOI: 10.1007/978-3-642-33701-7_4, Ó Springer-Verlag Berlin Heidelberg 2013

85

86

4 Properties

Fig. 4.1 Stress strain curves derived from tensile tests at RT of new HIS and reference steels, (a) engineering stress r0 and elongation e, (b) true stress r and true plastic strain u up to Au, thin lines: Ludwigson fit, see Table A2

terms of uniform and fracture elongation (Au, A), reduction of area (Z) and (iii) toughness expressed by the specific fracture energy Ws (deduced from the area under the curves in Fig. 4.1a) are derived and listed in Table A1 (Appendix A). Assuming constant volume the engineering stress–strain curves are converted into true stress– strain curves within the range of uniform elongation (Fig. 4.1b). A fit of the Ludwigson equation [2] r ¼ K1  un1 þ eK2 þn2 u

ð4:1Þ

to the experimental data yielded the constants presented in Table A2. The first term represents the Ludwik equation [3], which is adjusted to smaller strains by the second term. The meaning of each constant is discussed e.g. in [2]. However, the quality of the fit depends on the starting conditions, which in turn influences the constants. Here, the procedure is used to compare the exponent n1 of cold work hardening. The accumulated tensile properties are used to compare the new HIS with reference steels (Fig. 4.2). CN96 starts e.g. from a proof strength of 600 MPa and work hardens to a true fracture strength of more than 2500 MPa within 74 % of elongation reaching a work hardening exponent of 0.92 (see Sect. 3.2.1). The HNS CrMnN

4.1 Mechanical Properties

87

Fig. 4.2 Comparing key tensile properties of new HIS (shaded area) with those of reference steels (Table A1), Rp0.2 = proof strength, R = true fracture strength, n1 = exponent of cold work hardening (Ludwigson), Ws = specific fracture energy, the steel number refers to the chemical composition in Table 2.1

slightly exceeds this strength level but comes with a distinctly lower elongation. Necking starts earlier and proceeds to a higher reduction of area (Table A1, Fig. 4.1a). The high carbon grade MnC begins at a significantly lower proof strength, work hardens intensely but fails with hardly any necking after only 46 % elongation. As expected, the low interstitial steel CrNi offers the lowest proof strength, yet work hardens steadily at n1 = 0.66 up to A = 83 %. The high manganese TWIP steel MnCr77 surpasses this ductility and together with a higher strength level and work hardening (n1 = 1) the highest fracture energy Ws is achieved. The addition of (mass %) 1 Mo or 2 Cu to the new HIS lowers the strength level only slightly but the reduction of area by about one third (Table A3). The fracture energy remains quite high, though. The interstitial content of HIS, saturated with about 0.6 mass % N at normal pressure of air, is raised by adding carbon. This increases the strength of centrifugal castings (Fig. 4.3, Table A4) while the ductility is diminished. The properties of a sand casting are given for comparison. The strength level of castings is well below that of forgings, because the grain size is considerably coarser (Table A5). In fact, some of the high proof strength of CrMnN may be due to the fine-grained structure.

88

4 Properties

Fig. 4.3 Effect of interstitial content on strength (Rp0.2, Rm) and ductility (Z = reduction of area) of centrifugal castings GCN65 to GCN115, encircled = results of sand casting GCN85 (Table 2.1 and A4)

4.1.2 Tensile Properties at Subzero Temperatures Testing temperatures TT to -100 °C cover the climatic range and some industrial cooling operations. Embrittelment is known of HNS [1] and therefore two new HIS were also tested at -196 °C. The proof strength increases as TT is lowered (Fig. 4.4, Table A6, Sect. 3.2.2). The ductility of HIS and HNS is hardly affected down to -100 °C, but then drops to one tenth in liquid nitrogen. This ductile to brittle transition is not observed for CrNi. The Ludwigson fit for CN96 in Table A2 reveals that the exponent n1 of cold work hardening is raised only marginally, but that eK2 increases considerably, which is in line with the enhancement of Rp0.2 at lower TT. The true strength R grows as well which is reflected by a higher K1. The fracture face of CN96 tested at RT is covered with ductile dimples (Fig. 4.5). At -100 °C some intercrystalline fracture appears. The fracture in liquid nitrogen is transcrystalline, i.e. cleavage-like brittle. This corresponds to the low values of A and Z in Table A6. At -100 °C the ductility accumulated before the onset of fracture counts. The brittle patches of intercrystalline fracture develop as the crack passes by. The as-cast steel GCN85 reacts to a decrease of TT in about the same way as the hot worked grade CN85x but at a lower level of strength and ductility (Table A6).

4.1.3 Tensile Properties at Elevated Temperatures Strengthening of cold worked HNS was observed by aging at 500 °C [4]. Carbonitride precipitation starts between 700 and 750 °C after 1 h as depicted in Fig. 5.9. This poses the question, if HIS may be used at temperatures up to 700 °C without severe embrittlement or a loss of austenite stability. After heating, tensile specimens were soaked for 1 h before loading and tested at a crosshead speed of 0.5 mm/min. The results of hot tensile test are listed in Table A7. The engineering stress–strain curves of CN85 are smooth at 20, 400 and 500 °C but show serrations at 300 °C accompanied by softening and strain aging at 400° C

4.1 Mechanical Properties

89

Fig. 4.4 Effect of subzero testing temperature TT on proof strength Rp0.2 and fracture elongation A of HIS CN859, CN96, CN107 (shaded area, Table A6) and reference steels CrMnN and CrNi

(Fig. 4.6). A different type of waviness appears between 550 and 650 °C, the peaks of which are farther apart than those at 300° C. According to [5] this Portevin-LeChatelier effect is caused by interaction of dislocations and interstitial atoms at 300 °C and of substitutional atoms at 550–650 °C. The curve runs smoothly again at 700 °C. The start of precipitation at C550 °C is confirmed by intercrystalline corrosion after etching of metallographic sections taken from the gauge length. However, no precipitates were found by SEM within the austenite grains nor on grain boundaries of the neck at 600 °C (Fig. 4.7a). Even at 700 °C most grain boundaries in the gauge length did not appear to be decorated (Fig. 4.7b). Here, the failure mode turned from microvoid coalescence to intergranular cracking perpendicular to the direction of stress as TT was increased from 500 to 700 °C.

90

4 Properties

Fig. 4.5 Fractography (SEM) of CN96 after tensile tests at a 22, b -100, c -196 °C Fig. 4.6 Proof strength Rp0.2 and fracture elongation A of CN85 in dependence of the testing temperature TT (Table A7)

In conclusion, precipitation seems to be responsible for a loss of ductility at TT [ 500 °C but also for a strengthening effect (Fig. 4.6). The proof strength of CN85 at 700 °C is as high as that of CrNi at RT. Cold expanded retaining rings are shrink-fitted on generator shafts where they are moderately heated in service. Table A8 answers the question what loss of strength is to be expected of HIS, cold worked by 20 elongation and aged above shrink temperature, by heating up to 150 °C. At 150 Rp0.2 is about 10 % lower than at 100 °C.

4.1 Mechanical Properties

91

Fig. 4.7 Metallographic sections after hot tensile tests (CN85, SEM), a Longitudinal section through neck showing pores (600 °C), b Transverse section from gauge length (700 °C), GB = grain boundary

4.1.4 Creep Properties The new HIS are not meant for long-time creep service, because too much embrittling precipitation has to be anticipated. However, the short-time creep behaviour up to several 100 h may be of interest for tooling applications. Steel CN85 was selected to cut down on precipitation. It was tested in the as-quenched state and also after subsequent cold working by 20 % elongation at constant creep stresses and TT of 600, 650 and 700 °C in a range of 15–420 h. According to [6] creep properties are improved most by cold deformation between 10 and 30 %. The present investigation is of tentative nature because of a limited number of specimens. At the example of tests at 650 °C and a stress of 350 MPa the effect of cold working is explained in Fig. 4.8a,b The high initial creep rate of the solution annealed state comes down to the minimum creep rate e_ min of 0.0614 %/h after a creep elongation e(_emin ) of 5.39 %. In the cold worked condition the respective values are 0.00841 %/h and 0.13 % which points to extensive strengthening by precipitation. The time t(_emin ) to reach the minimum creep rate is, however, hardly affected by cold working, which slightly raises the time to fracture tf but lowers the elongation to fracture ef by more than an order of magnitude. A metallographic section was taken from the gauge length of the solution annealed specimen that fractured after 15 h at 650 °C, 400 MPa and 9 % elongation (Fig. 4.8c). Large precipitates and creep cracks impaired thinning of specimens for TEM. The figure was taken in SEM mode while selected area diffraction (SAD) of transparent precipitates was used to identify the phases. Carbides of type M7C3 and M23C6 have precipitated along grain and twin boundaries, while a eutectoid of M2N type nitrides in austenite grows into the grains. A dispersion of fine M3C carbides is distributed throughout the austenite grains. No traces of bcc structure were found by X-ray diffraction, which implies that the austenite remains stable in spite of less solute atoms. The embrittling effect of grainboundary precipitation promotes respective creep cracks (Fig 4.8d). The dispersed iron carbides are bound to contribute to the creep resistance but are too fast growing to keep this effect up.

92

4 Properties

Fig. 4.8 Results of creep test at 650 °C, steel CN85 solution annealed (SA) or subsequently cold worked by 20 % elongation (CW), a Creep rate e_ at a creep stress of 350 MPa in dependence of creep elongation e defining the minimum creep rate e_ min , the elongation to reach it e(_emin ) and the elongation to fracture ef, b creep rate e_ at the same stress in dependence of time t defining the time to reach the minimum creep rate t(_emin ) and the time to fracture tf. c Microstructure of SA specimen within the range of uniform creep elongation after fracture at 400 MPa, SEM, d As before showing creep cracks, arrow indicating direction of creep stress, LOM

The Larson-Miller parameter PLM = T (C ? log tf) is a means to combine creep results obtained at different temperatures. For T in K and tf in h those authors demonstrated in [7] a good fit for ferritic or austenitic steels and even for nonferrous metals, if C = 20. Notwithstanding the subsequent discussion on C the stress is plotted over PLM at C = 20 in Fig. 4.9 based on the results in Table A9. In spite of the enhanced strengthening, PLM is hardly changed by cold working. The results are clearly to the right of those measured for hot work tool steel X38CrMoV5-1 (AISI H11) quenched and tempered to a tensile strength Rm = 1425 MPa [8]. The better performance of fcc HIS in comparison to bcc tool steel is based mainly on the lower rate of diffusion in close packed crystals. Nucleation and growth of precipitates contribute to the creep behaviour as well.

4.1 Mechanical Properties

93

Fig. 4.9 Interdependence of creep stress and LarsonMiller parameter PLM comparing CN85 solution annealed (filled symbols) and cold worked by 20 % elongation (open symbols) (Table A9) with hot work tool steel H11 (X38CrMoV5-1) QT, Rm = 1425 MPa at 20 °C, CN85 was tested under constant stress, H11 under constant load [8]

4.1.5 Hardness The macrohardness (ISO 6507) only moderately increases with the interstitial content of centrifugal castings (Table A10). The respective values of the hot worked HIS (270–278 HV30) are almost twice as high as of CrNi (141 HV30). The macrohardness of MnC (211 HV30) stays below that of HIS, but the microhardness on the tensile fracture face (741 HV0.1) exceeds all HIS and reference steels. Cold work hardening is also evident after cold drawing CN96 (B423 HV30). Heating to the precipitation range raises the macrohardness of CN85 if accompanied by straining (B476 HV30).

4.1.6 Notch Impact Toughness The V-notch impact bending energy KV was measured according to EN10045. The specimens of hot worked steel were taken mostly in longitudinal direction, those of as-cast steel perpendicular to the direction of transcrystallisation. A typical result of HIS is shown in Fig. 4.10. Starting from a high toughness at RT a ductile-to-brittle transition occurs as TT is lowered. The respective DBTT is taken at KV = 100 J, which is still in the range of tough constructional steels. Steels CN96 attains KV = 364 J which is the highest of all measured values (Table A11). The respective steel GCN98 is the toughest of the as-cast grades at KV = 317 J. In general the new HIS are as tough as or tougher than low-interstitial CrNi which does not show a DBTT, though, but just a gradual decrease of KV as TT is lowered. The DBTT of HIS is close to -90 °C except for the high carbon grades CN107 and GCN115 which are just above -50 °C. In comparison the DBTT of the high nitrogen grade CrMnN is up to -21 °C. At room temperature KV of CN107, CN96 and CN85 is reduced by transverse taking but unexpectedly DBTT of CN85 is reduced as well.

94

4 Properties

Fig. 4.10 ISO-V impact toughness KV in dependence of testing temperature TT comparing steel CN96 and CN107, the ductile-to-brittle transition temperature DBTT is taken at KV = 100 J (Table A11)

The appearance of the fracture face turns from ductile dimples at RT to cleavage-like brittle fracture at -196 °C (Fig. 4.11). In the transition range at -80 °C dimple and intercrystalline areas are visible. Thus, about the same features of microfractography are encountered as after tensile tests (Fig. 4.5), but the ductile to brittle transition occurs at higher TT (see Sect. 3.2.3.1).

4.1.7 Rotating Bending Fatigue Rotating bending tests were carried out at RT, a frequency of 40 Hz and constant stress amplitudes ra until fracture or 107 number of cycles N. Specimens of hourglass shape and 5 mm smallest diameter were ground in axial direction by a flap wheel composed of abrasive cloth strips and loaded by rotating four-pointbending [9]. Up to 7 specimens were run at preselected ra levels to allow a statistical evaluation of the fatigue life Nf50 and the fatigue limit rf50 at N = 107 with a probability P = 50 % following the arcsin HP procedure [10]. Before the investigation a single specimen was run at 50 Hz under the highest load and observed by a thermo-camera. The temperature stayed below 40 °C until the very final cycles to fracture when it rose to 74 °C. Polished specimens were stopped after every 5000 cycles and searched for cracks by a microscope mounted on the test rig. The number of cycles to crack initiation Ni was taken at a crack length of 0.2 mm and compared to the number of cycles to fracture Nf. In the order of GCN85, CN96 and CN96 pre-streched by 20 % elongation the fatigue limit increases from 229 to 415 to 552 MPa and, divided by the respective proof strength of 447, 600 and 1022 MPa, we obtain rf50/Rp0.2 equal to 0.47, 0.69 and 0.54 (Fig. 4.12). Short cracks were initiated at slip bands on the specimen surface which is best to be seen on the coarse grained as-cast specimens

4.1 Mechanical Properties

95

Fig. 4.11 Fractography (SEM) of CN96 after ISO-V impact test at a 22, b -80, c -196 °C

(Fig. 4.13b). They follow crystallographic planes until after two or three grains they merge into the main crack which runs perpendicularly to the bending stress up to fracture. The ratio of (Ni/Nf)100 is generally well above 50 % except for one specimen where the crack started early, i.e. after 39 % of the fatigue life (Fig. 4.13a). The results of this limited study suggest that the fatigue limit at N = 107 attains about half of the proof strength. At higher stress amplitudes the fatigue life is shortened and mostly accompanied by earlier crack initiation.

4.2 Wear Resistance It is well known that wear changes the surface layer of materials, so that further wear is governed by the surface properties. Ductile austenitic steels tend to workharden in the wear surface and debris of wear-inducing counterbodies may become inserted into the surface. These processes depend on the type of loading. Abrasion by grooving wear comes with less workhardening than impact wear by mineral particles. The implosion of bubbles during cavitation is comparable to a particle impingement. Hadfield steel MnC has been used in applications where impact wear prevails because of its capacity to workharden. The new HIS were investigated in respect to their resistance to impact wear, abrasion and cavitation because their rate of workhardening is close to that of MnC, which is not stainless, though.

96

4 Properties

Fig. 4.12 Wöhler diagram of rotating bending fatigue tests plotting the stress amplitude ra over the number of cycles N, full lines = fatigue life (50 % probability), dashed lines = fatigue limit at N = 107 (50 % probability), digits = number of specimens not broken after N = 107

4.2.1 Abrasive Wear Grooving wear by mineral particles is met e.g. in mining, processing and transportation of ore and rock. This abrasion was investigated by a pin-on-plate test. The end face of area A of a pin specimen Ø 6 9 20 mm is moved under a load of 37 N at a speed of 4.8 mm/s over fresh flint grinding paper of 80 or 220 mesh size for a length L. The mass loss Dm is measured and converted to a loss of volume by means of the density q to give the dimensionless wear resistance W-1 ab = qAL/Dm. The results of pin-on-plate test are shown in Fig. 4.14. The wear resistance increases in the order of CrNi, CN96, MnC and is higher against the finer 220 mesh abrasive. The reciprocal value Wab is the wear rate (Table 4.1).

4.2.2 Impact Wear A schematic representation of the impact wear test is given in Fig. 4.15a: Individually counted particles of greywacke, mesh size 11–8 mm, impinge perpendicularly on a sample plate at a speed of 25 m/s. In fact, two samples are mounted on a rotor in adjacent positions and the wear loss of both is summed up to give the plot in Fig. 4.15b.

4.2 Wear Resistance

97

Fig. 4.13 Stages of fatigue, a Crack initiation at Ni (filled symbols) in percent of Nf (open symbols) of one specimen per ra. The Wöhler lines are taken from Fig. 4.12 for comparison, b Crack initiation at slip bands at the example of GCN85

Fig. 4.14 Abrasive wear resistance W-1 ab against 80 or 220 mesh flint paper

98

4 Properties

Table 4.1 Wear rate by abrasion Wab, wear rate by impact wear Wimp and wear rate by cavitation Wcav of different steels after running in. The factor f stands for the improvement of the wear resistance in relation to the low interstitial standard steel CrNi f Wimp f Wcav f Wab [10-5] [lg/impact] [lg/s] CrNi MnC CN96 GCN115

3.91 3.09 3.54 -

1 1.62 1.10 -

9.21 4.75 4.67 -

1 1.94 1.97 -

4.77 2.83 0.41 0.25

1 1.69 11.6 19.6

Fig. 4.15 Impact wear, a Schematic representation of impact conditions, b Mass loss depending on the number of impacts (Table A12)

After a running-in period, in which the wear surface is formed, the wear loss increases almost linearly, so that the wear rate Wimp can be derived (Table 4.1). Greywacke debris are embedded in the wear surface of the most ductile steel MnCr77 to such an extent that mass is picked up during the first thousand impacts. The wear resistance of the new HIS is about as good as that of Hadfield steel MnC. In fact all steels of high C, N or C ? N content lie in a rather narrow scatter band of wear loss, while the low interstitial steel CrNi wears about twice as fast. In a section normal to the wear surface the micro-hardness of a cold worked layer is highest at the surface and then drops as the distance from the surface grows. The hardness penetration is deepest for the softest steel CrNi where it exceeds 1 mm (Table A12). Cold work hardening is visible in the microstructure and described in Sect. 3.2.4.1. The most interesting feature of HIS and MnC is a thin amorphous top layer followed by a nanocrystalline zone.

4.2 Wear Resistance

99

Fig. 4.16 Mass loss by cavitation in dependence of test duration (courtesy of S. Huth)

4.2.3 Wear by Cavitation A local change of pressure in fluid energy machines may cause gas bubbles which subsequently implode. This is accompanied by small jets of fluid which exert impacts on a surface and degrade it until small particles are detached. This cavitation was investigated by an ultrasonic sonotrode of Ø16 mm mounted vertically on a piezo-quartz which oscillated at a frequency of 20 kHz with an amplitude of 40 lm dipped in destilled water at RT. Adjacent to the end face of the sonotrode a specimen was mounted below at a distance of 0.5 mm. After certain intervals the mass loss was measured. The results are plotted in Fig. 4.16. After an incubation time the mass loss increases about linearly to give the wear rate Wcav (Table 4.1). The lower the rate, the longer is the incubation time during which the surface is cold worked [11]. At an equal interstitial content the mass loss rate of GCN115 is more than an order of magnitude lower than that of Hadfield steel MnC. This is based on the greater toughness of HIS reflected e.g. by the higher specific fracture energy WS (Table A1, A4). The wear of CN96 proceeds more than an order of magnitude slower than that of CrNi.

4.3 Corrosion Resistance The corrosion resistance of stainless steels depends on a thin passive layer of complex structure, simply called Cr2O3 [12]. It has to be build up by the surrounding media and is under their constant attack which is influenced by chemical composition, temperature and fluid flow. We discern between general corrosion of evenly distributed mass loss and localised corrosion. It is therefore an extensive task to characterise the corrosion resistance of a new material and we have to confine the present study to a number of test examples. They comprise submersion tests, current density-potential-tests and tests on intercrystalline corrosion.

100

4 Properties

4.3.1 Submersion Tests Ground specimens of about 10 9 10 9 20 mm were submerged according to DIN 50905 in aqueous solutions of H2SO4 or HCl at room temperature for 120 h to measure the mass loss entailed by general corrosion. Initial tests showed, that specimens, which had been stored in the office for some weeks, were fully passivated and not attacked by 10 % H2SO4. Therefore all further specimens received a grinding finish by 1000 mesh SiC paper immediately before submersion. Usually the test is interrupted every 24 h to clean and weigh the specimens and to renew the test fluid. However, the exposure to air led to the passivation of some specimens thus inhibiting further corrosion and stopping the evolution of hydrogen bubbles. Therefore other tests were run for 120 h without interruption but with a renewal of the test fluid every 24 h during which the specimens stayed submerged. The mass loss Dm [g] is related to time and area to give the corrosion rate v [g/ m2h], which may be converted to the removal w [mm/a] per surface. The required densities of the new HIS were measured at 7.6–7.67 g/cm3. Early passivation means that there is no further mass loss and v as well as w decrease as time goes on, e.g. to the technically accepted range of w B 0.1 mm/a. The hot worked steels CN96, CN107 and the related alloys with Mo or Cu were weighed after 24 h in aqueous solutions of 5 or 10 mass % H2SO4 and in 1 mass % HCl. The exposure to air caused passivation and prevented any further mass loss or bubbling until the end of the 120 h interrupted test. Only CN103Mo1 met the target of w = 0.1 mm/a after 120 h in 5 % H2SO4. The remainder of specimens would have done so only after prolonged test duration. In spite of interruptions the specimens did not immediately passivate in 3 % HCl. The mass loss ceased after the fourth weighing of the almost completely dissolved CN96 specimen and after the third of CN94Mo1. Steel CN96Cu2 continued to lose mass until end of test. As for Mo the mass loss during active corrosion is lowered by Cu (Fig. 4.17) [13]. Uninterrupted test were run with the same steels in 1 and 3 % HCl and 10 % H2SO4. In the latter all specimens started with a strong effusion of hydrogen which continued for CN96, was slowed down for CN107 and CN94Mo1, after 48 h almost stopped in case of CN96Cu2 and completely ceased for CN103Mo. The corrosion products fell off easily during handling of the specimens and dissolved again in the fluid except for CN96Cu2. After removing the coat of corrosion products the surface of this specimen appeared copper-coulored and 6 mass % Cu were measured by EDX at 7 kV. The coat contained 36 mass % Cu (EDX, 15 kV). A protective layer of Cu is apparently formed on the surface during early corrosion which impedes further mass loss. But Mo and C reduce the corrosion rate as well (Fig. 4.18). However, hydrogen is emitted in 3 % HCl throughout the test. The corrosion rates are high and the effects of alloying less striking (Table A13). The series of centrifugal castings was submitted to uninterrupted submersion tests to evaluate the effect of carbon on corrosion in aqueous solutions of 1 and 3 HCl and 10 % H2SO4 [14]. In HCl the mass loss and corrosion rate were reduced

4.3 Corrosion Resistance

101

Fig. 4.17 Effect of Mo and Cu on the mass loss Dm during interrupted 120 h submersion tests in an aqueous solution of 3 % HCl

Fig. 4.18 Effect of C, Cu and Mo on the removal w after uninterrupted 120 h submersion tests in an aqueous solution of 10 % H2SO4

as the carbon content increased (Table A14). Little effect of carbon is noticed in H2SO4, especially if the test duration is prolonged 5 times (Fig. 4.19). Hydrogen bubbles were observed at the beginning of all tests but the gas evolution had stopped in 10 % H2SO4 and 1 % HCl, before the test solution was renewed after the first 24 h. It did not start again throughout the remainder of the 120 or 600 h test duration. Bubbling continued in 3 % HCl resulting in a high mass loss. Carbon apparently changes the mechanism of dissolution from shallow pit corrosion (Fig. 4.20a, b) at a low C content (GCN65) and high mass loss (Fig. 4.19) to corrosion along crystal planes (Fig. 4.20c, d) at higher C contents (GCN88 to GCN115) and lower mass loss (Fig. 4.19). These results are consistent

102

4 Properties

Fig. 4.19 Effect of carbon on the removal w after uninterrupted submersion tests in aqueous solutions lasting 120 or 600 h

with earlier work [15]: Corrosion is impeded by the most resistant crystal planes (e.g. Fig. 4.20d) but increases, if this barrier is overcome. The resulting shallow pit corrosion (e.g. Fig. 4.20b) is less dependent on crystal structure.

4.3.2 Current Density/Potential Tests In submersion tests corrosion proceeds by chemical reactions at the resting potential. If the potential is influenced from outside, we speak of electro-chemical corrosion. In a respective test according to DIN 50918 and ASTM G5-94 the potential between the specimen electrode and a reference electrode is raised and the resulting current measured at room temperature. The reference calomel electrode of potential UC = +244.3 mV is connected to the electrolyte via a HaberLuggin capillary. The dissolution of metal atoms yields an electric current which is related to the test area and plotted as current density in dependence of the steadily increasing potential, that is transformed to the potential UH = 0 mV of the standard hydrogen electrode SHE. Two hours after final grinding with 1000 mesh SiC paper the specimens were inserted into the electrolyte which was then purged with nitrogen for 30 min to expel remainders of oxygen. After a cathodic treatment at -1 V (SHE) for 1 min, the resting potential UR was measured for 30 min. Starting from 10 mV below UR the current density/potential curve was recorded by raising the potential at a rate of 0.6 V/h. Tests were run at room temperature in 0.5 molar H2SO4, corresponding to an aqueous solution of 4.74 mass % and in an aqueous solution of 3 mass % NaCl. Typical curves are shown in Fig. 4.21 and key potentials and current densities of the steels investigated are listed in Table A15.

4.3 Corrosion Resistance

103

Fig. 4.20 Surface of centrifugally cast steels after uninterrupted 120 h submersion in (a, c, d) 1 % HCl, (b) 3 % HCl

Diluted electrolytes of H2SO4 are commonly used to study general corrosion of stainless steels. The rising potential stands for an increase of the redox potential encountered in practical corrosion systems. The higher the resting potential UR and the shorter the distance to the passivating potential UP, the smaller is the range of active corrosion. In some cases a prepassivation occured at UPP before the passivating potential UP at which the current density iP drops to the low level i0 of the passive range. It is terminated at the break down potential UB which is, however, interrupted by repassivation between URP and the final break down at UB0 . To visualise the corrosion attack, specimens were polarised at specific potentials. A typical result is depicted in Fig. 4.22. In the active range severe corrosion along crystal planes and by shallow pits becomes visible. In the passive range the grains appear but scratches from grinding are still noticeable. At the repassivating potential only a few local pits have been formed, although the corresponding current density iRP is high i.e. about a quarter of iP. This means that a major part of the current is spent on building a layer of corrosion products, which was observed earlier [16]. The chemical composition of HIS has little effect on the shape of the curves in 0.5 m H2SO4 (Table A15a). There is no influence of carbon on the resistance of centrifugal castings. Their higher Cr content raises UR and lowers i0, though. The Mo content has a similar effect on UR and distinctly lowers the passivating current. As to the reference steels, CrMnN comes with a repassivation just like HIS, while CrNi does not. The key values of the latter come quite close to some of the HIS as e.g. GCN98.

104

4 Properties

Fig. 4.21 Characteristic current density/potential curves, a In 0.5 m H2SO4, b In 3 % NaCl, type I with passive range, type II without. The indices stand for A = activating, B = break down, P = passivating, PP = prepassivating, R = resting, RP = repassivating, 0 = passive

Sea water and salted roads give rise to pitting corrosion of stainless steels by Cl- ions. A simple method to measure the liability of such a steel to this type of localised corrosion is a current density-potential test in a 3 % NaCl electrolyte. The attention is not so much on the passive range of low i0 but on the breakdown potential. The higher UB, the better is the pitting resistance (Fig. 4.21b). As the potential is raised the breakdown may not be abrupt and a criterion is necessary to define UB. The results in Table A15b are based on a lasting increase of i0 by 10 lA/cm2. This sharp criterion works well with type I curves in Fig. 4.21 but may be misleading in case of type II. This is best documented in linear plots of the curves (Fig. 4.23). In these examples steel CN85 is of type II, starts to deviate early and the respective UB is -0.19 V. However, the major breakdown occurs at UB = 0.69 V. The evaluation of UB is further aggravated by repassivation of early pits which leads to small intermittent peaks of current before final breakdown. UB of CN96 (0.2 V) is slightly lower than that of CrNi (0.26 V) but very much surpassed by UB of CN107 (0.91 V). Together with the higher Cr content the high carbon content of GCN115 leads to UB = 1.28 V. This points to a beneficial influence of carbon on the pitting resistance of stainless nitrogen steels. Steel CrMnN is an example of the latter and in spite of its high Cr content does not exceed UB = 0.74 V. The grade GCN65 is a nitrogen steel as well, but its low breakdown potential of 0.18 V is most likely connected with remainders of

4.3 Corrosion Resistance Fig. 4.22 Surface of CN859 after polarization in 0.5 m H2SO4 at RT for 1 h against SHE, a In the active range at U = -0.256 V, b In the passive range at U = 1044 V, c At the repassivating potential U = 1294 V

105

106

4 Properties

Fig. 4.23 Current density i in dependence of potential U (SHE) at RT in 3 % NaCl

d-ferrite in the as-quenched microstructure. The addition of Mo comes with an increase of UB from CN96 (0.2) to CN94Mo1 (0.44 V) but with a decrease of UB from CN107 (0.91) to CN103Mo1 (0.88 V). Two mass % Cu clearly bring UB down to the lowest value (-0.04 V) of all steels tested.

4.3.3 Tests on Intercrystalline Corrosion The high interstitial content of the new steels is liable to form precipitates along grain boundaries during cooling from the temperature range of homogeneous austenite or during reheating [17]. This may cause a depletion of Cr adjacent to the boundaries and entail intercrystalline corrosion [18]. The susceptibility to intercrystalline corrosion (IC) was investigated by Strauß tests according to EN ISO 3651-2 method A. Specimens of Ø4 9 10 mm were gas quenched from TSA, ground to strips of 4 9 1 9 10 mm, subjected to the test solution, bent by 90° and inspected for cracks by standard 10 times magnification and by SEM. As no cracks were detected, no IC had occurred. The small size of specimens was chosen to minimise the case to core difference of cooling time. In Sect. 5.3.3 the cooling time was prolonged step by step until IC was observed.

4.4 Magnetic Properties The magnetic moments of individual iron atoms in HIS perform a random motion between the possible orientations of their projections. If an external magnetic field is applied, the moments tend to align in the field direction making a Larmor precession. Temperature causes a chaotic flipping of the moments which prevents their alignment. The combined action of these processes results in the paramagnetic state of HIS. However, at technically possible field strengths, the system

4.4 Magnetic Properties

107

Table 4.2 Temperature dependence of the volume magnetic susceptibility v taken at H = 0.33 T susceptbility v T [K]

MnC (10-6)

CN107 (10-4)

CN96 (10-6)

CN85 (10-4)

300 250 200 150 100 50 4.2

2.276 2.276 2.276 2.276 2.276 2.276 2.276

0.815 0.825 0.828 0.828 0.828 0.828 0.828

0.442 0.443 0.443 0.443 0.444 0.445 0.448

2.052 2.734 3.174 3.374 3.414 3.424 3.424

Table 4.3 Relative magnetic permeability lr measured by Förster Magnetoscope 1.067-103 with a handheld sensor. In a first series the specimens were tested at room temperature after (1) solution annealing, (2) subsequent cold drawing to a reduction of 16–37 % and (3) subsequent deep freezing in liquid nitrogen for 15 min. In a second series solution annealed specimens were tested at room temperature (4) on the deformed impact wear surface and (5) after subsequent deep freezing in liquid nitrogen for 2 h 1 2 3 4 5 Steel CN85x CN96 CN107 CN96 CN96 CN96 CN96 CN94Mo1 CN103Mo1 CN96Cu2 GCN65 GCN88 GCN98 GCN115 GCN85 MnCr82 CrMnN CrNi MnC

Solution annealed

Cold drawn

1.0010 1.0008 1.0007 – – – – 1.0008 1.0009 1.0010 1.0150 1.0009 1.0009 1.0008 1.0014 1.0011 1.0010 1.0900 1.0010

– – – 16 20 29 37 – – – – – – – – – – – –

% – – – 1.0011 1.0008 1.0008 1.0009 – – – – – – – – – – – –

Cooled to -196 °C

Impacted

Cooled to -196 °C

1.0011 1.0011 1.0011 1.0011 1.0008 1.0008 1.0009 1.0008 1.0009 1.0010 1.0150 1.0009 1.0009 1.0008 1.0014 1.0011 1.0012 – –

1.0011 1.0010 1.0011 – – – – – – – – – – – – 1.0011 1.0010 1.1200 1.0025

1.0011 1.0011 1.0011 – – – – – – – – – – – – – 1.0012 – –

stays far from a total alignment of the moments and thus far from saturation. The degree of alignment along the magnetic field, i.e. the magnetic moment M of a unit volume, is linked to the applied field strength H via the volume susceptibility v, in

108

4 Properties

short M = vH. For convenience it is generally accepted to use Tesla (T) as unit for M and H, so that v is dimensionless. As pointed out in Sect. 3.1.3.2, the paramagnetism of HIS is of the Van Vleck type. A transformation to the antiferromagnetic state was not observed down to a temperature of 4 K. As shown in Table 4.2, the volume susceptibility is quite small and only slightly increases during cooling. In a technical sense this persistent paramagnetism translates to a nonmagnetic behaviour of workpieces. They are not prone to guide magnetic fields or generate heat in alternating electromagnetic fields. Applications are mentioned in Sect. 6.2.3. Deviations from a nonmagnetic behaviour may occur by small fractions of a ferromagnetic phase, e.g. remainders of d-ferrite or strain-induced a-martensite, embedded in paramagnetic austenite. The relative permeability lr, i.e. the slope of the virgin magnetization curve, was measured to make sure that no trace of ferromagnetic phase was present, neither after deep freezing nor after straining. The results in Table 4.3 confirm that lr of HIS is quite low except for GCN65 which contains remainders of d-ferrite. This applies also to CrNi which in addition picks up some a-martensite by impact deformation. The very thin amorphous surface layer on impacted HIS, although ferromagnetic (Sect. 3.2.4.1), is not revealed by this test method.

References 1. Riedner S (2010) Höchstfeste nichtrostende austenitische CrMn-Stähle mit (C ? N), doctoral thesis, Ruhr-University Bochum 2. Ludwigson DC (1971) Modified stress-strain relation for FCC metals and alloys, Metallurg. Transactions A2:2825–2828 3. Ludwik P (1909) Elemente der Technologischen Mechanik. Springer, Berlin 4. Uggowitzer PJ, Speidel MO (1991) Ultrahigh-strength Cr-Mn-N steels. In: Stainless Steels’91, Chiba conference, The Iron and Steel Institute of Japan, pp 762–770 5. Nikulin I, Kaibyshev R (2011) Deformation behavior and the Portevin-Le Chatelier effect in a modified 18Cr-8Ni stainless steel, Material Sci. and Engin. A 528:1340–1347 6. Böhm H, Schirra M (1973) Einfluss der Kaltverformung auf das Zeitstand- und Kriechverhalten einiger warmfester austenitischer Stähle. Arch. Eisenhüttenwes. 44:785–791 7. Larson FR, Miller J (1952) A time-temperature relationship for rupture and creep stresses, Trans ASME 74:765–775 8. Berns H (1975) Das Zeitstandverhalten von Warmarbeitsstählen und seine Bedeutung für die Auslegung von Blockaufnehmern und Druckgusskammern, habilitation thesis, Technical University Berlin 9. Berns H, Gavrilujk VG, Nabiran N, Petrov YuN, Riedner S, Trophimova LN (2010) Fatigue and structural changes of high interstitial stainless austenitic steels, steel research int 81(4):299–307 10. Dengel D, Dahl W (ed) (1978) Verhalten von Stahl bei schwingender Beanspruchung, Verlag Stahleisen, Düsseldorf, pp 23–46 11. Berns H, Siebert S (1996) High Nitrogen austenitic cases in stainless steels. ISIJ Int 36:927– 931 12. Kruger J (1988) Passivity of metals—a materials science perspective. Int Mat Rev 33:113– 130

References

109

13. Berns H, Riedner S, Hussong B (2010) Influence of molybdenum and copper on the corrosion resistance of high strength austenitic steels. Mater Sci Forum 638–642:2979–2985 14. Berns H, Hussong B, Riedner S, Wischnowski F (2010) Effect of carbon on stainless austenitic FeCrMnN steel castings, steel res int 81:245–251 15. Dobbelaar ALJ, Herman CME, Dewit HWJ (1992) The influence of the microstructure on the corrosion behaviour of Fe-25Cr. Corros Sci 33:779–790 16. Schwenk W (1963) Beobachtungen über die Korrosion nichtrostender Stähle in Schwefelsäure unter potentiostatischen Bedingungen, Werkstoffe u. Korrosion 14:646–654 17. Mujica Roncery L, Weber S, Theisen W (2011) Nucleation and precipitation kinetics of M23C6 and M2N in an Fe-Mn-Cr-C-N austenitic matrix and their relationship with the sensitization phenomenon. Acta Mater 59:6275–6286 18. Dayal RK, Parvathavarthini N, Raj B (2005) Influence of metallurgical variables on sensitisation kinetics in austenitic stainless steels. Int Mat Rev 50:129–155

Chapter 5

Manufacture

We discern between hot manufacturing processes as e.g. melting, casting, hot working, heat treatment and welding or cold processes as cold forming and machining, although some heat may be generated. The hot steps are accompanied by phase transformations which come the closer to thermodynamic equilibrium the higher the temperature and the longer the time of treatment are. The sequence of slow solidification, hot working and solution annealing therefore approximates to phase diagrams. On the basis of the actual chemical composition of the new HIS a set of isoplethal diagrams is calculated to explore (a) the variation of C ? N, (b) the variation of C/N and (c) the mole fraction in dependence of temperature (Appendix B, Figs. B1–B10). From these, key temperatures are derived. These are the liquidus and solidus temperatures TL and TS, which mark the range of solidification, and the temperature TP, which stands for the begin of precipitation in austenite and thus about ends the range of hot working and solution annealing. These temperatures are influenced by the carbon content (Fig. 5.1). As it increases TS is lowered more than TL which widens the range of solidification enhancing segregation and, together with a rise of TP, narrows the range of homogeneous austenite. This is the more so, if 2 mass % copper are added while the effect of 1 mass % Mo or the higher Cr content in the castings is less significant. At the example of CN85, CN96 and CN107 (Figs. B1b–B3b) it becomes clear that only CN85 is close to (C/N)op defined in Fig. 2.3. At the higher carbon contents TP is raised above Top which is marked by (+) in Fig. B1b for comparison. The higher TP, the higher is the rate of diffusion and therefore the tendency to precipitation during quenching of heavy cross-sections. The cold processes are depending on work hardening which increases as the deformation proceeds (Fig. 4.1b). The high work hardening exponents n1 in Table A2 point to high forces required for metal forming or cutting. The respective microstructural changes are described in Sect. 3.2.1.

H. Berns et al., High Interstitial Stainless Austenitic Steels, Engineering Materials, DOI: 10.1007/978-3-642-33701-7_5, Ó Springer-Verlag Berlin Heidelberg 2013

111

112

5 Manufacture

Fig. 5.1 Liquidus and solidus temperature TL, TS and begin of precipitation at TP in dependence of the carbon content derived from equilibrium phase diagrams of HIS in Figs. B1 to B10, encircled = HIS with 2 mass % Cu (Table A18)

5.1 Melting and Casting All new HIS were molten in induction furnaces the capacity of which ranged from 4 kg to 5 Mg. The final temperature in the furnace has to be higher, if a ladle is used. The melt from small furnaces is usually poured directly into the mould. So there are no general recommendations for the tapping or pouring temperature of new HIS. Nitrogen was added via nitrided ferrochromium and melting as well as pouring was done under normal pressure of air.

5.1.1 Ingots Melts of 5 Mg were produced of CN96 and CN107 (Table 2.1) at Friedrich Lohmann GmbH in D-58454 Witten and teemed to three square ingot moulds each of tapered shape that was reduced to cylinders of Ø310 mm by forging. These electrodes were electro-slag-remelted at Energietechnik in D-45143 Essen, to give ESR ingots of Ø430 mm. This remelting was done in a pressurised facility (PESR) because of availability, which was run at a pressure of 5 bar argon. A 2.5 Mg melt of CN85 was teemed at the Deutsche Nickel in D-58239 Schwerte to a conical mould of 350–315 mm diameter and 3900 mm length. The ingot was processed in a standard ESR unit under normal pressure of air to give a remelted ingot of Ø510 mm. The grades CN94Mo1, CN103Mo1 and CN96Cu2 were molten in a lab furnace of 4 kg capacity and poured into a Y-shaped steel mould. After removal of the feeder small slabs of 150 by 70 by 30 mm were ready for hot rolling.

5.1 Melting and Casting

113

Fig. 5.2 Scheil simulation of GCN88 compared to equilibrium solidification, Scheil range of phases: 1 = L, 2 = L ? F, 3 = L ? F ? A, 4 = L ? A, 5 = L ? A ? M23C6

5.1.2 Centrifugal Castings The alloys GCN65 to GCN115 (Table 2.1) were molten in a 450 kg furnace at Klaus Kuhn Foundry in D-42477 Radevormwald. Starting from GCN65 portions of the melt were successively enriched with carbon in the laddle to give a series of four steels. These were solidified by horizontal centrifugal casting. The mould of 100 mm inner diameter consisted of low carbon steel S355 and was cooled from outside. Tubes of Ø100 by Ø40 mm in size were produced at a rotational speed of 200 rpm. The cooled metallic mould raises the rate of solidification and increases the deviation from equilibrium. This is expressed by a Scheil simulation (Fig. 5.2) provided by Thermo-Calc which resembles a non-equilibrium cooling that allows of infinitely rapid diffusion in the liquid but none in the solid state except for the light elements carbon and nitrogen. Compared to equilibrium the solidus temperature is lowered by almost 100 °C. Rapid solidification is expressed in the macrostructure by the radial growth of elongated primary grains (Fig. 5.3a). Remainders of d-ferrite are encountered in GCN65, while some decorated grain boundaries are visible in GCN115 even after quenching (Fig. 5.3b).

5.1.3 Sand Castings Sand moulds were prepared by the KSB foundry in D-91257 Pegnitz and cast from an 80 kg furnace to give castings as depicted in Fig. 5.4 for further investigation. After slow cooling in the mould the grain boundaries are decorated with pseudo pearlite (Sect. 5.3.2). These M(C,N) precipitates extend close to the surface,

114

5 Manufacture

Fig. 5.3 Microstructure after centrifugal casting and solution annealing (LOM). a Section through a cast tube of GCN88. b Decorated grain boundary of GCN115

Fig. 5.4 Drawing (mm) of sand castings made of GCN85 with positions of specimens (grey) for mechanical testing. Strips s = 26 mm thick were cut from castings for further hot rolling

indicating that there is little decarburisation or loss of nitrogen after casting (Fig. 5.5). The surface itself is smooth, reflecting just the roughness of the sand mould. After solution annealing and quenching EDX line scans revealed microsegregations of Cr and Mn. The maximum content of these elements divided by the minimum content stands for the degree of segregation S and amounts to SCr = 1.27 and SMn = 1.34. The average spacing is about 50 lm [1].

5.1.4 Refractories It was shown that melting, remelting and different casting procedures of the new HIS are industrially feasible. A nitrogen content of about 0.6 mass % was reached without pressure- or powder metallurgy which, together with carbon, is a sound basis for interstitial strengthening. To reach this high nitrogen content, nickel had to be replaced by manganese which tends to interact with refractories, i.e. furnace lining, slag, sand mould and inclusions. The monolithic lining of induction furnaces used to be made of rammed fireclay-bonded sand to give mullite (3Al2O32SiO2) which is accompanied by expansion that prevents a strike of the melt to the cooling coil [2]. This acidic lining matches with an acidic slag. Early work of Körber [3, 4] suggested that Mn and—to a lesser extent—Cr are capable to reduce SiO2 and the

5.1 Melting and Casting

115

Fig. 5.5 Section normal to the as-cast surface of GCN85 (SEM). a Decorated grain boundaries. b Pseudo-pearlite growing from grainboundary

more so the higher the temperature and the lower the carbon content are. This implies that Mn is transferred from the steel to the slag and SiO2 is lost from the lining which is accompanied by a rising Si content in the melt. This corrosive wear of the quartz constituent in the lining is most pronounced at the slag line. A basic MgO lining and a respective slag would be resistant to the attack of Mn but without expansion of the rammed lining. A compromise is seen in an MgAl2O4 spinell with excess Al2O3. This refractory material comes with sufficient expansion and is called neutral between acidic and basic [5]. Melting of larger HIS volumes in an electric arc furnace may be handled in a hearth consisting of MgO bricks, which stay hot. In the foundry the interaction of high manganese melts with sand moulds is inhibited by a proper facing, as e.g. zirconia [6]. This procedure is well known from Hadfield manganese steel.

5.2 Hot Working The Ø430 mm ESR ingots of CN96 and CN107 were forged at Friedrich Lohmann to bars of Ø120 mm. The Ø510 mm ESR ingot of CN85 was rolled at Deutsche Edelstahlwerke in D-58452 Witten to 180 mm square and further reduced to bars of Ø65 mm by a radial forging machine. Sections of CN96, Ø290 9 530 mm in size, were upset forged, punched and ring-rolled on a tyre rolling mill at Bochumer Verein in D-44793 Bochum to obtain tyres with rim for rail vehicles. The laboratory slabs were hot rolled to half of their initial thickness at the Institut für angewandte Materialtechnik of the University in D-47057 Duisburg. Last not least a strip was taken from a GCN85 sand casting according to Fig. 5.4 and rolled from 26 mm initial thickness to 11 mm final thickness at the Max Planck Institute in D-40237 Düsseldorf. According to Table 2.1 this material is called CN85x and allows to study the effect of hot working on the mechanical properties by comparing CN85x in Table A1 with GCN85 in Table A4. The recrystallized fine grained microstructure of the former (Table A5) considerably

116

5 Manufacture

raises strength and ductility. However, the higher strength entails a small decrease of the notch impact toughness at room temperature and a slight increase of the DBTT (Table A11). Comparing the tensile properties of CN85x and CN85 (Table A1) does not show much effect of previous electro-slag-remelting the latter. As the testing temperature is decreased to -100 °C the ESR grade provides a higher proof strength and respectively less elongation A but more necking, i.e. a higher reduction Z (Table A6). After ESR the notch impact energy at RT is definitely raised by about 20 %, while the DBTT is slightly lowered (Table A11). The higher upper-shelf energy of ESR is consistent with a better cleanness. Smaller nonmetallic inclusions induce smaller microvoids during straining and fracture by macrovoid coalescence requires more deformation. This is more pronounced for the multiaxial stress state of the notch impact test than for the tensile test. Hot working with different types of forging and rolling steps covered a wide range of cross-sections and were done without any unexpected incident or scrap.

5.3 Heat Treatment Solution annealing followed by quenching is the main heat treatment of HIS. Aging is sometimes applied to cold worked HIS.

5.3.1 Solution Annealing Precipitates are observed in the microstructure after slow cooling from casting or hot working (Fig. 5.5). The purpose of solution annealing is to dissolve these precipitates and achieve a structure of homogeneous austenite. To this avail the solution anneal temperature TSA has to be above the equilibrium temperature of beginning precipitation TP which is shown in Fig. 5.1. An increase of TSA abbreviates the time for dissolution and allows to dissolve different chemical compositions locally provoked by microsegregation (see Sect. 5.1.3), which are not reflected by Thermo-Calc. There is, however, a limit to TSA because of grain growth which impairs the strength of hot worked steel. The coarser grain size of castings (Table A5) is less affected by TSA which in general is raised above that of hot worked steel, because a casting is not exposed to the effect of diffusion annealing during soaking at the forging temperature. Remainders of undissolved precipitates may be encountered in castings (Fig. 5.3b) but rarely in forgings. The steels investigated were solution annealed at different temperatures, quenched in water and inspected by LOM and SEM to find TSA experimentally. The results are compiled in Table A16. The equilibrium pressure of nitrogen gas pN2, calculated by Thermo-Calc for TSA of each steel, is required to keep the nitrogen content in the surface at the given level. It depends on the chemical

5.3 Heat Treatment

117

Fig. 5.6 Subsurface concentration of elements in CN96 after solution annealing at 1100 °C for 3 h in air. The results stem from glow discharge spectroscopy

composition and the temperature, but stays in all cases below the partial pressure of nitrogen in air. Therefore some up-take of nitrogen is to be expected rather than a loss. This is indeed supported by measurements depicted in Fig. 5.6. After prolonged annealing of CN96 in air, nitrogen is increased towards the surface, while some carbon is lost by oxidation. In total, the interstitial content of the subsurface zone is changed little, which is underlined by Fig. 5.7a, which shows this zone after precipitation annealing at 900 °C to generate pseudo-pearlite. It is evident that a thin layer of only 0.2 mm in depth is enriched with pearlite, respectively C ? N. The exchange of elements between the steel surface and air is inhibited by a thin coat of oxides which are rich in chromium. This is not the case, if solution annealing is carried out in an N2 atmosphere. As a result the subsurface content of nitrogen is raised considerably (Fig. 5.7b) as indicated by the higher content of precipitates near the surface which is topped by a thin nitride layer. In contrast solution annealing in vacuum entails a severe loss of N and Mn (Fig. 5.8). Again pseudo-pearlite is used as a marker. Its content is reduced from core to case followed by layers of (i) austenite plus grainboundary precipitates, (ii) ferrite plus carbides and finally (iii) plain ferrite at the surface, the Mn content of which was down to 0.5 mass % as measured by EDX [7].

118

5 Manufacture

Fig. 5.7 Microstructure (LOM) in a section normal to the surface of CN96 after annealing at 1100 °C for 3 h. a In air of about 1 bar pressure. b In nitrogen of 1 bar pressure, followed by quenching in water and subsequent annealing at 900 °C for 2 h in air to generate pseudo-pearlite

5.3.2 Interrupted Quenching The constitutional diagrams in Figs. B1 to B10 give evidence of phases as M2N nitrides, M23C6 carbides and r-phase below TP. It is the purpose of quenching from TSA to subdue their precipitation. In this context it is important to know the temperature range in which the rate of precipitation is highest. Therefore specimens were solution annealed, cooled to an isothermal temperature between 1000 and 700 °C, held for 1 h and quenched in water. The resulting content of pseudopearlite was evaluated by microscopy and plotted in Fig. 5.9. It increases with the C ? N content which raises the temperature of maximum precipitation. The lower concentration of grain boundaries in the coarse grained as-cast steel GCN85 reduces the volume of precipitates as compared to hot worked CN85x. The addition of 1 mass % Mo tends to reduce the volume of pseudo-pearlite while 2 mass % Cu have little effect (Fig. B11). Pseudo-pearlite is known from high nitrogen steels where it consists of M2N and austenite lamellae. A specimen of CN96 was intensely etched to locally remove the austenite and analyse a protruding lamella by EDX (Fig. 5.10). About half of the atoms consist of C and N and the other half of metal ones. Notwithstanding the problem of measuring light elements quantitatively the MX lamella is not an equilibrium phase (Fig. B2). Even after annealing at temperatures between 950 and 1020 °C the specimens did not fully arrive at equilibrium, but grain boundary precipitates pointed to carbides while those within grains resembled nitrides. This coincides with results found after creep at 650 °C (Fig. 4.8c) and with the general experience that carbon has a strong affinity to grain boundaries while nitrogen has not. The method of interrupting the quench by isothermal holding was also applied to notch impact specimens to measure the effect of precipitation on toughness KV and sensitisation to intercrystalline corrosion IC. Only after a short holding of \1 min the specimens stay free of IC. Within 10 min KV comes down to a fully brittle level, especially if the C ? N content is high (Fig. 5.11).

5.3 Heat Treatment Fig. 5.8 Microstructure in a section normal to the surface of CN96 after annealing at 1100 °C for 3 h in vacuum of 4  10-5 bar, quenching in pressurized N2 gas and subsequent annealing at 900 °C for 2 h in air to generate pseudo-pearlite. a Overview of the sub-surface zone (LOM). b Near surface zone enlarged (SEM). c EBSD analysis of the near surface zone

119

120

5 Manufacture

Fig. 5.9 Content of pseudopearlite after solution annealing and isothermal holding at temperature Tis for 1 h and subsequent quenching in water

5.3.3 Continuous Quenching Interrupted quenching has shown that the climax of precipitation occurs between 1000 and 700 °C. The t10/7 cooling time is therefore a suitable measure to search for the onset of precipitation at the critical cooling time tc10/7 (Fig. 5.12). Precipitation begins at grain boundaries and may be detected by microscopy or notch impact testing (Fig. 5.11). However, a most sensitive method is a Strauß test for intercrystalline corrosion according to EN ISO 3651-2 method A. Specimens of Ø4 9 10 mm were continuously gas quenched from TSA at different t10/7, ground to plates of 4 9 1 9 10 mm, subjected to the test solution, bent by 90° and inspected for cracks at 10 times magnification. After increasing t10/7 step by step tc10/7 is assigned to the transition of uncracked to cracked specimens. The results are presented in Fig. 5.13. The critical cooling time is lowered as the C ? N content or the C/N ratio is raised. The latter stems from an increase of TP above Top which is shown in Fig. B1b. The higher TP, the more the range of precipitation during quenching is shifted to faster diffusion which is indirectly reflected in Fig. 5.9. Copper appears to be an exception, because it raises TP (Fig. B6) but also tc10/7 (Fig. 5.13) [8]. The grainboundary segregation of carbon is inhibited by copper just as by nickel which may explain this result. The cooling time of a work piece depends on the quenchant and on the size and shape of the cross-section in which it grows from case to core. In the order of quenching in water, oil or air the notch impact toughness was reduced considerably and the more so the higher the C ? N content. The sensitisation to intercrystalline corrosion increased in the same order (Table 5.1). These results call for quenching in water or even intensive quenching in agitated water to reduce the steam envelop. Published data on natural cooling in water are a guideline for simple shapes as round or rectangular bars of different size. Numerical simulations are another way of finding t10/7 from case to core of a workpiece. The higher the C ? N content or the C/N ratio, the smaller is tc10/7 and

5.3 Heat Treatment

121

Fig. 5.10 Microstructure of steel CN96 after solution annealing and isothermal holding at 900 °C for 10 min. a Pseudo-pearlite after deep etching in V2A at 60 °C for 25 min (SEM). b EDX analysis of protruding lamella, see frame in a

also the cross section that stays free of precipitation to the core. However, cooling is faster at the surface which tends to prevent IC. Some loss of toughness may be acceptable in the core. In other applications IC is not to be expected. Therefore the admissible cross-section is related to the in-service conditions. During manufacture a sensitisation to pickling agents has to be avoided.

122 Fig. 5.11 Effect of solution annealing and subsequent holding in a salt bath at 900 °C for tis B 600 s on the notch impact energy KV and the sensitisation to intercrystalline corrosion IC

Fig. 5.12 Schematic representation of the t10/7 cooling time and the critical cooling time tc10/7 at which precipitation begins, TSA = solution anneal temperature, TP = equilibrium temperature of beginning precipitation, G = grain boundary decoration, D = discontinuous precipitation of pseudopearlite

Fig. 5.13 Critical cooling time tc10/7 of new HIS

5 Manufacture

5.3 Heat Treatment

123

Table 5.1 Notch impact toughness KV [J] and intercrystalline corrosion IC (+) in dependence of the t10/7 cooling time [s] during quenching from TSA (Table A16) in water, oil or air, respectively, measured in the core of 10 mm square test bars steel water oil air CN 96 CN107

KV

t10/7

IC

KV

t10/7

IC

KV

t10/7

IC

367 348

1.4 2.2

– –

278 193

5.9 6.9

; +

204 12

47.1 52.4

+ +

Fig. 5.14 Effect of cold working and aging on hardness HV30 of CN96 measured at RT. a After 44 % of cold upsetting. b After reduction of cross-section by cold-drawing

5.3.4 Aging Aging is caused by a rearrangement of interstitial atoms in cold worked austenite up to prestages of precipitation. To induce a deformed structure the height of a solution annealed and quenched cylinder Ø13 mm was cold compressed from 20 to 11.2 mm which corresponds to an upset strain of 44 %. Immediately after cold working the specimen was consecutively aged at temperatures up to 550 °C. The first peak of hardness at 400 °C in Fig. 5.14a is most likely caused by aging, the second at C550 °C by precipitation (see Sect. 4.1.3).

5.4 Cold Drawing The intensive work hardening of HIS may be used to raise the proof strength. Cold drawing of rod or wire is one way to go. Solution annealed rods of Ø18 mm were processed on a Schumag 2B drawing bench of 8 MN force at the Deutsche Edelstahlwerke in D-58089 Hagen. In several stages the reduction of cross-section

124

5 Manufacture

Fig. 5.15 Effect of reduction by cold drawing of rods on the engineering stress–straincurves of CN96, the dashed line represents the true stress

was raised to 37 % which corresponds to the elongation obtained. A plot of respective engineering stress–strain-curves reveals that the yield point is raised by proceeding cold reduction and that it follows the true stress curve (Fig. 5.15). After 37 % reduction the 0.2 % proof strength is about doubled and the ratio Rp0.2/ Rm is raised to 0.84 (Table A17). Fracture elongation and energy as well as the workhardening exponent n1 are reduced, because part of the plasticity has been spent by cold drawing. Remarkable is Z = 58 % after 37 % reduction by drawing starting from Au = 12 % which gives proof of the high workhardening capacity and ductility of HIS during necking. Shortly after drawing the hardness was measured and again after 3.5 years. Age hardening had occurred at room temperature as depicted in Fig. 5.14b.

5.5 Welding Tungsten inert gas (TIG) welding is widely applied to stainless steels. The electric arc creates a hot spot on the steel surface that entails a weld pool in which the heat is distributed by convection. As the heat source moves a solidifying mushy zone is travelling behind and a heat affected zone (HAZ) is induced in the base metal. A summary of TIG welding high nitrogen steels is given in [9] and also applies to high interstitial steels. The main problems encountered are (i) a loss of N and Mn from the weld pool, (ii) the formation of nitrogen bubbles which do not fully escape from the bath but are caught in the mushy zone forming pores, (iii) precipitation in the HAZ. Admixing a few percent of nitrogen to the argon shielding gas is a means of counteracting a loss of nitrogen and the formation of bubbles, although the corrosion of the tungsten electrode may increase. Using a filler metal of higher N solubility is another way of keeping nitrogen in the weld pool. The evaporation of manganese is impeded by the pressure of the shielding gas and a good coveridge of

5.5 Welding

125

Fig. 5.16 Weldment of MnCr77 sheet using a 3 kW continuous Nd:YAG laser a normal section through the seam (LOM). b Center of weld pool (LOM), (courtesy of L. Mujica Roncery)

the pool. This is best achieved if the bead is kept narrow. Mixtures of argon and helium raise the arc temperature to more than 10000 °C and are not recommended. It is essential to retain the content of N and Mn not only in respect to the strength level but also to subdue d-ferrite during solidification which would enhance degassing of N2. Experience with TIG welding of austenitic HNS demonstrates that flawless weldments without pores or hot tearing are feasible but that precipitates in the heat affected zone cause intercrystalline corrosion and some loss of ductility. This was corroborated recently in an investigation of welding stainless TWIP steels of high interstitial content as e.g. MnCr77 in Table 2.1. Solution annealing after welding will restore a fully austenitic structure. Continuous laser welding of sheet 1.2 mm thick was carried out in [10] without intercrystalline corrosion in the HAZ. The weld pool was about as wide as the sheet thick, i.e. the pool volume per unit length was very small and therefore the self-quenching rate high enough to prevent precipitation (Fig. 5.16). The hardness of the bead is slightly raised. In contrast to fusion welding experience with pressure welding of HNS suggests that electric resistance heating or friction goes without a loss of elements and appears to be applicable to HIS if general rules for high alloy steels are obeyed. Depending on the local cooling rate, subsequent solution annealing may be necessary.

5.6 Machining A high yield strength, intensive cold work hardening, and much ductility entail a strong resistance of the new HIS to metal cutting. This resistance is related to the specific fracture energy Ws of HIS (Tables A1, A3, A4) which is about 50 % higher than that of standard austenitic steel CrNi and 2 to 7 times as high as that of ferritic-pearlitic or quenched and tempered steels [11]. The chip formation on HIS is therefore accompanied by high forces and by heat. The latter is not readily dissipated because of a low thermal conductivity. A large rake angle is a way to

126

5 Manufacture

reduce the cutting force acting on the tool. This implies a sharper cutting edge because rake angle, wedge angle of the tool tip and clearance angle add up to 90°. The sharper the cutting edge the more fragile it is and the less brittle the tool material should be. This is the reason why mostly WC–Co sintered hardmetals are used which are abbreviated to tungsten carbide or WC tools. Because of their limited temperature resistance the cutting speed v is kept at a low level. In rough turning using M15 to M30 tool inserts with a chip breaker, v may be down to 25 m/min and in fine turning with M10 to M25 between 40 and 50 m/min. The feed (depth of cut) should be high enough to undercut most of the cold worked surface zone formed during the previous revolution but not overload the tool edge. In some applications a feed of 0.4 mm/revolution was used. In general the machine-tools should be quite rigid to keep vibrations low which would add to workhardening. From Hadfield steel we know that milling is possible but that drilling of deep holes may pose a problem. There is no experience with the new HIS, but a similar behaviour is expected. Hydro-jet-cutting with abrasives was applied to rough cuts and electro-spark-erosion to fine machining operations and slender holes.

References 1. Berns H, Nabiran N, Mujica L (2012) High interstitial stainless austenitic steel castings, steel research int. doi:10.1002/srin.201100332 2. Chesters JH (1957) Steelplant refractories. The United Steel Comp. Ltd., Sheffield 3. Körber F (1937) Einfluss der Beimengung auf die Reaktionen zwischen Eisenschmelzen, Eisen Mangan-Silikaten und fester Kieselsäure. Stahl u. Eisen 57(48):1349–1355 4. Körber F (1936) Zur Metallurgie der Eisenbegleiter. Stahl Eisen 56(3):77–104 5. Schacht CA (ed) (2004) Refractories handbook. CRC Press, Boca Raton 6. Rudolph S (1994) Betrachtungen zum Aufbau von Form- und Kernschlichten unter besonderer Berücksichtigung ihrer feuerfesten Bestandteile. Gießerei-Praxis 8:165–178 7. Riedner S, Berns H (2008) Wärmebehandlung hochfester, nichtrostender Austenite. HTM 63(2):84–94 8. Berns H, Riedner S (2008) Zusammenhang zwischen Konstitution und Wärmebehandelbarkeit hochfester austenitischer Stähle. HTM 63:337–341 9. Gavriljuk VG, Berns H (1999) High Nitrogen Steel. Springer, Berlin 10. Mujica Roncery L (2010) Development of high-strength corrosion-resistant austenitic TWIP Steels with C ? N, doctoral thesis, Ruhr University Bochum, Bochum 11. Berns H, Gavriljuk VG (2007) Steel of highest fracture energy, Trans Tech Publications. Key Eng Mater 345–346:421–424

Chapter 6

Assessment

The aim of this concluding Chapter is to summerise and assess the contents of previous chapters. On the scientific side the structure/property relation is of key interest. The technical side is mainly concerned with manufacture and application based on the achieved properties. The new high interstitial steels (HIS) are assessed by comparing them with known austenitic steels. Their acronyms and chemical compositions are given in Table 2.1.

6.1 From Structure to Properties The multiscale approach of investigating the new HIS, depicted in Fig. 1.1, is an attempt to relate macroscopic properties to microstructural features. Starting from the scale of electron structure it was demonstrated that joint alloying with carbon and nitrogen increases the concentration of free electrons in the CrMn austenite. Respective measurements of conduction electron spin resonance (CESR) were corroborated by ab initio calculations (Sect. 3.1.1). Free electrons enhance the metallic character of interatomic bonds and thus the ductility. At the same time they improve the homogeneity of the atomic distribution, i.e. reduce clustering of alloying atoms and support their short-range ordering. This more even distribution of alloying atoms in steels with C ? N was recorded by Mössbauer spectroscopy (Sect. 3.1.2). Three methods were employed to measure the degree of chemical homogeneity: (i) By CESR the size of clusters in HIS was estimated at 3 nm and their concentration at 1.55  1019 cm-3 (Sect. 3.1.3.1). (ii) Magnetic measurements revealed a cluster size of about 5 nm for CN107 and 3.5 nm for CN85, while MnC alloyed only with carbon showed a cluster size of about 8 nm. (iii) Measurements of the stacking fault energy (SFE) led to two maxima of dislocation splitting which hints to short-range decomposition. The difference DSFE between the two maxima decreased from CN107 to CN85. This points to less clustering,

H. Berns et al., High Interstitial Stainless Austenitic Steels, Engineering Materials, DOI: 10.1007/978-3-642-33701-7_6, Ó Springer-Verlag Berlin Heidelberg 2013

127

128

6 Assessment

which was highest for MnC (Sects. 3.1.3.3 and 3.1.3.4). In all, the SFE of HIS is inversely proportional to the concentration of free electrons. Originally the joint addition of C ? N was meant to circumvent costly pressure metallurgy: A part of N in high nitrogen steels (HNS) was replaced by C to give new HIS molten at normal pressure of air. Now it is evident that C ? N not only ease production but exert a beneficial influence on the electron structure of HIS. Their ductility is enhanced and their chemical homogeneity is improved on a nanoscale, i.e. on a length scale more than four orders of magnitude below the scale of microsegregations (Sect. 5.1.3). As a result of the more even atomic distribution the stability of austenite to phase transformation is raised. The lower the alloy concentration in the clusters the less they are prone to initiate precipitates and the less the areas in between are liable to provoke bcc transformation. This is reflected by thermodynamic equilibrium simulations which see a wide phase field of homogenous austenite in the Fe-18Cr-18Mn-C–N system but not in the Fe18Cr-18Mn-C system (Fig. 2.1c, a). In fact an optimal C/N ratio exists at which the austenite is most stable, i.e. extends to the lowest temperature (Fig. 2.3).

6.1.1 Mechanical Properties A key property of the new HIS is the specific fracture energy Ws which corresponds to the area below an engineering tensile stress/elongation curve. It profits from a high yield point, from intense cold work hardening and from ductility. Ws of CN96 is 100 % higher than that of the Hadfield carbon grade MnC, 60 % higher than that of standard low interstitial steel CrNi and 20 % above that of the nitrogen grade CrMnN (Table A1), which shows a higher DBTT in notch impact testing, though (Table A11). CN96 is topped only by the high manganese TWIP steel MnCr77 which is, however, less corrosion resistant. The individual contributions of the proof strength Rp0.2, the work hardening exponent n1 (Eq. 4.1) and the elongation at fracture are compared in Fig. 4.2 for HIS (shaded area) and reference steels. The high interstitial content of HIS leads to a proof strength which is almost three times that of low interstitial CrNi. In spite of a higher interstitial carbon content MnC does not by far live up to the proof strength of HIS. This is explained by a significantly stronger affinity of nitrogen atoms to dislocations (Sect. 3.2.1.1). Starting from a high yield point plastic straining of HIS is accompanied by a succession of structural changes: planar slip followed by twinning, occasional emartensite and a high density of dislocations (Sect. 3.2.1.3). Their joint strengthening effect delays necking and raises the uniform elongation of CN96 to 61 % compared to 45 % for MnC and 44 % for CrMnN. This evidences the advantage of alloying with C ? N compared to C or N. Reference steels of higher uniform or fracture elongation are either of low interstitial CrNi type or contain less chromium which is bound to enhance the concentration of free electrons (Fig. 3.6) but reduces the corrosion resistance.

6.1 From Structure to Properties

129

Effect of temperature. The fracture energy Ws is raised as the tensile test temperature is lowered to -100 °C (Table A6). This depends mainly on an increase of the yield strength, because the work hardening exponent n1 is hardly changed (Table A2) and the fracture elongation is moderately reduced. The higher yield strength is related to a decrease of SFE at subzero temperatures, which is caused by an increase in the concentration of free electrons (Sect. 3.2.2). Between -100 and -196 °C a severe embrittlement is observed (Fig. 4.4) which is known from high nitrogen but not from high carbon austenite. In notch impact tests a DBTT in the range of -100 °C is found for HIS based on a still generous toughness level of KV = 100 J (Table A11). The respective DBTT of the high nitrogen steel CrMnN is shifted to -21 °C, while the high carbon grade MnC starts from a lower energy at RT and reaches 100 J at -59 °C but without a distinct ductile- to-brittle transition. Compared to alloying with C the mobility of dislocations is enhanced by N or C ? N because of a higher concentration of free electrons. This improves the toughness at RT. As the testing temperature is lowered the number of glissile dislocations in a pile-up is higher for N and C ? N steels which tends to open microcracks if slip is blocked, e.g. by Lomer-Cotrell barriers. This ‘‘ductile’’ crack initiation resembles the onset of embrittlement (Sect. 3.2.3.1). At elevated temperatures serrations are observed in engineering stress/straincurves of HIS at 300 °C and again at 550–650 °C which are caused by an interaction of dislocations with interstitial or substitutional solute atoms, respectively [1]. In the upper temperature range a sensitization to intercrystalline corrosion points to a beginning precipitation. Actually the new HIS are designed to develop optimal properties in the as-quenched state of homogeneous austenite but some precipitation assists the short-time creep resistance at temperatures below 700 °C. The creep elongation is reduced, though, especially by grain boundary decoration (Sect. 4.1.4). Effect of strain rate. Increasing the strain rate of HIS MnCr82 at RT raised the proof strength without affecting the reduction of area [2]. At -196 °C an HNS showed a mixture of brittle and dimple fracture at low strain rate while fully dimple fracture prevailed at high strain rate [3]. This unusual behaviour was explained by local heating caused by a localization of slip (Sect. 3.2.3.1). If it comes to high velocity ballistic impact at RT the phenomenon of localised slip leads to adiabatic shear bands and a perforation of CN96 targets. The high fracture energy Ws measured by a low strain rate is not available at ballistic rates (Sect. 3.2.3.2). Effect of cyclic loading. In push/pull-tests at RT the microstructural changes revealed by TEM are similar to those after uniform straining except for additional nitride precipitates (Sect. 3.2.4.2). In rotating bending short cracks start from slip lines at the surface and follow crystallographic planes until, after two or three grains, they extend normal to the bending stress up to fracture. The fatigue limit at N = 107 cycles is not a true one but bound to come down, if cycling proceeds [4]. It is raised by previous cold working and in either case is about 50 % of the proof strength. This is a key advantage of HIS in respect to low interstitial steel CrNi which offers a distinctly lower proof strength and fatigue limit [5].

130

6 Assessment

6.1.2 Wear Behaviour Scratch tests have shown [2] that the specific scratch energy es (J/mm3) increases from 5.76 for CrNi to 12.81 for MnCr82 and 15.10 for MnC. The so-called fabvalue of these three steels is about the same, i.e. 0.82–0.85 on a scale of 0 to 1. At fab = 1 a chip is cut without deformation on either side of the groove, at fab = 0 the groove volume is plastically displaced into ridges along the groove just by ploughing without any removal of material. These results show that the energy es to cut a chip is almost tripled by a high content of interstitials. The equally high fab-value of the three steels supports the view that plasticity is more or less confined to the chip and only a minor part is spent on the vicinity of the grove, i.e. on work hardening of the surface. This lack of substantial cold work hardening is most likely the reason why the abrasion resistance of HIS CN96 is by only 10 % higher than that of low interstitial steel CrNi (Table 4.1). In contrast impact wear by mineral particles is accompanied by severe deformation of the wear surface into which mineral debris is embedded. The wear resistance of CN96 is higher by about 100 %. The repeated impacts bring about a thin surface layer of amorphous structure followed by a nanocrystalline one below, but only on CN96 and MnC and not on low interstitial CrNi (Sect. 3.2.4.1). This seems to be related to a vacancy-interstitial interaction [6]. In cavitation the surface is impacted by water jets of imploding bubbles. Here the wear resistance of CN96 is higher by more than an order of magnitude compared to CrNi (Table 4.1) and the wear surface is not affected by embedded mineral debris. It is interesting to note that the high carbon grade MnC is by far less resistant to wear by cavitation than the high interstitial grade CN96 although their rate of cold work hardening is equally high (Table A2). It is obvious to assume that the higher concentration of free electrons in the steel with C ? N (Sect. 3.1.1) promotes ductility and delays the initiation of cracks in the process of delaminating flakes of material from the surface. This is seen parallel to the much higher specific fracture energy Ws of CN96 in tensile tests compared to MnC (Table A1). Just as the well known high carbon Hadfield steel MnC, the new high interstitial steels promise a high wear resistance as long as sufficient work hardening of the surface is involved in the wear process. In contrast to MnC the new HIS passivate because they contain about 18 mass % Cr. This is of advantage if wear occurs in a corrosive environment.

6.1.3 Corrosion Resistance Wet corrosion in aqueous solutions tends to attack the weakest spot on the surface of stainless steels which is generally characterised by a reduced concentration of passivating elements. In this respect the chemical homogeneity is of major importance. As shown in Sect. 3.1.2 a high concentration of free electrons in HIS

6.1 From Structure to Properties

131

enhances short range atomic ordering and thus lowers short range decomposition on a nanoscale. The accumulation of chromium atoms in clusters is reduced and the solute atoms are spread more evenly. As microsegregations are the outmost result of short-range decomposition, joint alloying with C ? N promises a beneficial effect not only on the nanoscale but also on the microscale distribution of solute atoms and thus on the corrosion resistance. This is the more important as manganese, added to promote the solubility of nitrogen, is generally seen as unfavourable in respect to corrosion resistance [7]. On the other hand nitrogen is known to improve the resistance to pitting corrosion [8] and several explanations have been forewarded [9, 10]. It was also shown that carbon raises the pitting resistance if alloyed together with nitrogen [11]. Calculated phase diagrams (Figs. 2.9 and 2.10) reveal that the constitution of HIS allows of adding molybdenum to further increase the resistance to pitting which was verified experimentally (Sect. 4.3). The addition of copper to reduce general corrosion is limited by the evolution of gas from the melt (Fig. 2.11) and by a steep rise of TP, the temperature of beginning precipitation from austenite (Fig. 2.12). In conclusion one can say that the beneficial effect of C ? N makes up for the detrimental one of manganese, so that the new HIS come close to standard steel CrNi in respect to general corrosion. In respect to pitting corrosion they are distinctly better, though. A drawback of HIS is their sensitisation to intercrystalline corrosion, if the critical cooling time is surpassed. This may limit the size of the asquenched cross section.

6.1.4 Nonmagnetic State The new HIS are paramagnetic down to a temperature of 4 K and their susceptibility is quite small (Sects. 3.1.3.2 and 4.4). Their relative permeability remains at a low level even after plastic deformation and deep freezing indicating the absence of ferromagnetic traces. The austenite is very stable which amounts to a nonmagnetic behaviour of workpieces.

6.2 From Manufacture to Application The feasibility of manufacturing at reasonable costs is a prerequisite of transferring the excellent properties of new HIS to products. At this early stage of pilot production no serious assessment of costs is possible. However, the exchange of nickel by manganese makes the melt cheaper and the high strength tends to save cross-section, respectively weight of workpieces.

132

6 Assessment

6.2.1 Constitution and Hot Manufacture The higher the temperature and the slower the cooling rate, the closer the metallurgical processes come to equilibrium represented by phase diagrams. Therefore the results of respective calculations, visualised in Figs. B1 to B10, contain valuable information on the temperature range of solidification, hot working and solution annealing, especially if slow-cooling heavy cross-sections are concerned. The phase diagrams including the new HIS clearly point out that below the high temperature range of homogenous austenite a forbidden zone of carbide and nitride precipitation exists, which has to be quickly transgressed by quenching. The solidus temperature TS decreases and the begin of precipitation at TP increases, if the interstitial content is raised (Fig. 5.1). This reduces the range of homogenous austenite in which hot working is performed. The situation is aggravated by enhancing the rate of solidification as shown by a Scheil simulation in Fig. 5.2. Therefore soaking at the initial temperature of hot working is important to reduce the effect of segregation or even eliminate minor amounts of liquid phase which is most detrimental in respect to hot workability. Some deformation below TP is possible, if continuous straining leads to a dispersion of precipitates instead of grain boundary decorations which promote tearing. Quenching at the end of hot working, before precipitation has occurred, saves the cost of solution annealing and promotes grain refinement which enhances the yield strength. Of the new HIS (No. 1–11 in Table 2.1) GCN65 contained a few percent of dferrite indicating the lower end of interstitial content. TL and TS of this alloy are about 1380 and 1340 °C, respectively (Fig. B7). In comparison TL of the low interstitial standard steels Cr18Ni10 and Cr17Ni12Mo2 amounts to about 1,460 and 1450 °C and TS to about 1430 and 1420 °C, respectively [12]. In these grades precipitation is not a problem as far as hot working is concerned, but d-ferrite exists in equilibrium down to 1300 and 1250 °C, respectively. If segregation is taken into account areas of weaker d-ferrite are to be expected at the beginning of hot working. Under certain conditions this may lead to tearing. This weakness is not encountered in HIS which, however, require higher deformation stresses because of solid solution hardening by the interstitial elements. The constitution of the new HIS allows of alloying Mo (Fig. 2.9) to enhance the resistance to pitting corrosion. Grades with 1 mass % Mo have been manufactured (Figs. B4 and B5). The addition of Cu is restricted by the evolution of N2 gas (Fig. 2.11) and by a higher TP (Figs. 5.1 and B6). The forbidden zone of unwanted precipitation starts at TP and ends at about 500 °C. The lower TP, the lower is the temperature range of precipitation which is thereby retarded. In Fig. 2.3 the phase field of homogeneous austenite extends to the lowest temperature TP = Top at (C/N)op. This is made use of in steel CN85 (Fig. B1b), which is quite close to the optimal conditions marked by (+). The higher carbon content of CN96 or CN107 raises TP and thus reduces the critical cooling time (Fig. 5.13). Below 500 °C, aging occurs in cold worked HIS (Fig. 5.14a). At temperatures of C550 °C a sensitization to intercrystalline

6.2 From Manufacture to Application

133

corrosion is observed (Sect. 4.3.3) caused by a faint precipitation along grainboundaries followed by a severe one between 700 and 1000 °C (Figs. 5.9 and B11). Any stay in this temperature range from &500 °C to TP (Fig. 5.1) is likely to impair the structure of homogeneous austenite. It should be avoided or repaired by solution annealing and quenching. This applies to stress relief annealing, welding, brazing, hot coating and other hot processes. Creep loading is a tentative exception and the only diversion from a homogeneously austenitic structure (Sect. 4.1.4). Quenching from solution anneal temperature TSA (Table A16) must stay below the critical cooling time tc10/7 to prevent precipitation. This time becomes the shorter the higher the interstitial content and, as the nitrogen content of all new HIS is close to 0.6 mass %, the higher the C/N ratio is (Fig. 5.13). Therefore the steel selection has to be adapted to the as-quenched cross-section of a workpiece. This intense sensitivity to the cooling rate is not observed in standard stainless austenitic steels but is known from ferritic steels or ferritic-austenitic duplex steels [13, 14]. As described in Sects. 5.1 and 5.2 sand castings, centrifugal castings and ingots were produced on an industrial scale. The latter were electro-slag-remelted and hot worked by rolling and forging to different semi-finished products, solution annealed and quenched. No unexpected incident or scrap occurred.

6.2.2 Workhardening and Cold Manufacture Drawing, cold forging and coining operations are chipless, i.e. the degree of deformation stays below the tearing limit. Blanking, punching and machining go beyond this limit to cut off material. In both cases the Rp0.2 proof strength is a guideline for the stress required to initiate deformation. The ultimate tensile strength Rm, although not a true stress, corresponds to uniform elongation Au and just avoids necking in chipless deformation. The true fracture strength R stands for a separation or cut, if we set aside the different stress states. If we compare e.g. the high interstitial steel CN96 with the low interstitial standard steel CrNi (Table A1), we obtain the following ratios: Rp0.2 ? 600/221 = 2.71, Rm ? 1020/592 = 1.72, R ? 2547/1930 = 1.32. These results indicate that, at a given tool geometry, the new HIS require distinctly higher stresses to initiate plastic flow and keep it up. This calls for stronger machines to provide sufficient force but also bears on the tools. Thus the sheet thickness for blanking may have to be reduced and the slenderness of punched holes as well. The degree of deformation that a material will survive is strongly depending on the stress state. The elongation ratios of the above high and low interstitial steels are AU ? 61/70 = 0.87 and A ? 74/83 = 0.89. They stay below one but still promise excellent formability of HIS, provided the tools can take the load. Moderate heating up to e.g. 200 °C will help to lower the required flow stress (see Table A8) and enhance the obtainable deformation.

134

6 Assessment

The ratio of the workhardening exponents (Table A2) amounts to n1 ? 0.92/ 0.66 = 1.39. It underlines that the resistance to plastic flow of CN96 considerably surpasses that of CrNi. This is a disadvantage of cold forming HIS but an advantage of cold formed workpieces which offer a higher strength. It may further increase by natural age hardening at room temperature (Fig. 5.14b). A heat treatment of artificial aging will anticipate and define the increase of hardness (Fig. 5.14a). In Sect. 3.2.1 it was shown, that workhardening is brought about by a sequence of structural changes at such a rate that necking is delayed and elongation is enhanced. In consequence workhardening promotes strength and ductility simultaneously which is expressed by the specific fracture energy Ws containing both properties. The ratio of Ws is 675/422 = 1.60 (Table A1). This demonstrates that CN96 requires 60 % more energy than CrNi to deform a unit volume, which has to be provided by the press. In service CN96 may consume 60 % more energy before fracture.

6.2.3 Application The new HIS offer some excellent in-service properties (Chap. 4). They are related to structural features which were studied in great depth (Chap. 3). This sound structural foundation, summarised in Sect. 6.1, adds reliability to the properties measured. It is the combination of two or more properties which offers new areas of application. All of these profit from lower alloy costs. The forbidden temperature range of precipitation (Sect. 6.2.1) calls for a sufficient quenching rate. Therefore applications of moderately thick cross-sections are to be preferred. Although thin, sheet material is not seen in the centre of application because of restrictions in respect to welding (Sect. 5.5). The stress seems to rest on castings, hot rolled rods and rings as well as forgings, all machined to final size. Strip may be cold rolled and wire cold formed by drawing and forging as far as permitted by workhardening (Sect. 6.2.2). In the following a few combinations of properties are discussed to demonstrate the potential of new HIS in application. Yield strength + corrosion resistance. In respect to standard CrNi steel the yield strength of HIS is higher by a factor more than 2.5. This offers the chance to reduce the cross-section and save weight, i.e. go for light-weight construction e.g. in transportation. Workhardened + nonmagnetic corrosion resistant. Cold forging and surface rolling led to a surface hardness of 60 HRC in rings made of CN107 for roller bearings [15]. As this level of hardness is reached without ferromagnetic a-martensite, such rings combined with ceramic balls may be used as bearings in the vicinity of strong magnetic fields as demonstrated on a lab scale by the Schaeffler KG in DE 97424, Schweinfurt (Fig. 6.1a). Intermittent heating hardly affects the hardness at room temperature, although the hardness tested at elevated temperature declines (Fig. 6.1b).

6.2 From Manufacture to Application

135

Fig. 6.1 Cold worked surface of the inner ring of a roller bearing made of CN107 with a bore diameter of 25 mm. a Scatter band of the hardness profile HV1 measured along three paths (see arrows). b(1) Hardness HV0.05 of the cold worked surface measured at room temperature after holding at elevated temperature for 2 h or b(2) hot hardness measured at the elevated temperature after holding for 2h

Another application is seen in sea water. Here the higher resistance to pitting corrosion of HIS compared to martensitic stainless steels could be an advantage for bearings. Energy consumption + corrosion resistance. In road construction the hillside flanks are sometimes covered by a net of thick wire to catch falling rock. Their energy has to be consumed by the protective structure even at subzero temperatures. The exceptionally high specific fracture energy Ws of the new HIS offers an effective energy consumption which even increases down to -100 °C (Table A6). As HIS are stainless no protective coating is required. In contrast to their application under tensile loading a vehicle crash impact occurs at a lower hydrostatic stress, i.e. at an even higher ductility. At ballistic velocity of an impact, adiabatic shear bands are formed [16] which localize the strain and reduce the consumed energy (Sect. 3.2.3.2). Discs of CN107 and CN96, each 5 or 10 mm thick, were perforated by standard soft-core ammunition 7.62 9 51 WK (Nato level 1, STANAG 4569) at a velocity VZ of 835–840 m/s [17]. In an austenitic high nitrogen steel a ballistic limit of 535 m/s was found [18]. This suggests that a hard layer on top of HIS is liable to reduce the velocity by flattening or fracturing the bullet. In such a hard ductile-sandwich structure HIS could be a most effective ductile partner. A comparison of tensile stress–strain curves in Fig. 6.2 points to a higher specific fracture energy of

136

6 Assessment

Fig. 6.2 Engineering stress– strain curves, full lines = measured, dashed line = hypothetic. Steel A = low alloy martensitic steel with about 0.3 mass C, steel B = high alloy maraging steel X2NiCoMoTi 18-8-5 with 0.7 mass Ti, steel CN96 = solution annealed, CN96-40 = cold streched by 40 %, CN96-op = after optimal cold working

solution annealed steel CN96 (perforated) but a much higher yield strength of steels A and B (not perforated). After 40 % of cold stretching grade CN96-40 comes closer to steel B and after optimal cold working CN96-op may approach steel A. This leads to the proposal to generate a hard layer by surface rolling or severe peening a HIS and create an in situ sandwich structure. Wear and corrosion resistance. As shown in Table 4.1, the resistance to wear increases, if the degree of work hardening involved is raised. Compared to low interstitial standard steel CrNi the resistance W1 ab of high interstitial steel CN96 to abrasion is improved by a factor of only 1.1 and stays below that of Hadfield steel MnC which is not stainless though. The resistance W1 imp of CN96 to impact wear is, however, raised by a factor of 1.99, which is above that of MnC. Finally the factor of improving the resistance to cavitation W1 cav is up to 11.6, i.e. much higher than that of MnC (1.69). In contrast to impact wear by mineral particles no mineral debris are imbedded in the wear surface. These results suggest that the new HIS are suited to replace stainless CrNi steels in case of wear by impacting particles, and also replace steel MnC in aggressive surroundings as e.g. in deep pit or marine mining. In these applications CrNi or MnC castings often suffer from deformations because of unpredictably high service loads. In this respect the higher strength of HIS is bound to be an advantage. The highest profit is expected of applying HIS to fluid flow machines as e.g. pumps and armatures. In this case the corrosion resistance may be enhanced by alloying molybdenum. Strong + nonmagnetic. Nonmagnetic steels are used in the vicinity of alternating currents to prevent unwanted heating of structures and tools or near direct currents to prevent diverting of the magnetic field. The coils of transformers and magnets exert forces which are e.g. counteracted by nonmagnetic frames. Low cost and high strength HIS may be used to about -100 °C because of embrittlement

6.2 From Manufacture to Application

137

below (Figs. 4.4 and 4.10). Cold expanded retaining rings of moderate cross-section may be suited to hold the wiring on electric generator shafts. Elevated temperature service. The hot strength of CN85 derived in tensile tests up to 700 °C (Fig. 4.6) is not above that of hotwork tool steel H11. However, aging of cold worked CN96 at temperatures up to 550 °C keeps the hardness, measured at RT, at or above the initial level (Fig. 5.14a). This means that intermittent heating by friction or by residual heat after shut down is not necessarily accompanied by a drop of strength. The resistance of CN85 to short-time creep exceeds that of hot work tool steel H11 the more, the higher the temperature and the longer the duration of the test (Table A9, Fig. 4.9). Prestraining seems to have little effect on the time to creep rupture of CN85 but raises the initial strength. Cold working is, however, confined to smaller crosssections. This may be partly overcome by semi-hot working, to raise the initial strength and induce a dispersion of precipitates. Mechanical loading is usually higher during the stroke of a hot work press but thermal fatigue on the tool surface may be more severe in a die-caster and the more so, because the thermal expansion of austenitic HIS is higher and the thermal conductivity lower than that of martensitic tool steels. Thus the application of HIS for tooling is open for shop floor trial. Body friendly steel. About 10 % of the human population are allergetic to nickel, which may cause local inflammations of the skin or the tissue inside the body which stays in prolonged contact with e.g. zippers, wrist watches, costume jewellery, dental braces, piercings or implants [19]. This does not only apply to Ni coated parts but also to those made of stainless CrNi steels. High nitrogen steels were produced by pressure metallurgy and Ni exchanged by Mn. This is in line with high interstitial steels which do not require costly pressure equipment. By selecting the scrap the Ni content could be reduced in respect to that of the melts listed in Table 2.1. The addition of Mo would improve the corrosion resistance.

6.3 Pros and Cons of HIS The arguments pro and contra the use of new high interstitial austenitic steels are briefly summarised.

6.3.1 Pros The unique combination of high proof strength (Rp0.2 & 600 MPa) and cold work hardening exponent (n1 & 0.9) combined with a fracture elongation of A &70 % amounts to a tensile fracture energy between 600 and 700 J/cm-3 which is way

138

6 Assessment

above that of conventional stainless and other steels. In short: Stainless austenitic HIS are strong and tough! The intensive work hardening enhances the resistance to wear by impacting mineral particles and to wear by cavitation. In respect to the latter, HIS are proved to be superior to Hadfield steel, which is not stainless. The combination of high strength and good corrosion resistance of HIS especially in respect to pitting is unique. Quenched and tempered stainless steels may attain a higher proof strength but are less corrosion resistant. This applies vice versa to stainless duplex steels. HIS are nonmagnetic even at cryogenic temperatures or after severe cold working, which can lead to a hardness of 60 HRC without a trace of a-martensite. They even retain good properties after moderate heating. The industrial manufacture of HIS at normal pressure of air did not pose any special problems. The costs are reduced by refraining from alloying nickel.

6.3.2 Cons The numerous positive properties of HIS are contrasted by two negative ones. The first one is concerned with a tendency for precipitation, if the cooling rate during quenching from solution anneal temperature is too slow. This implies that the cross section of work pieces is limited and has to be matched with the C ? N content and C/N ratio of the steel, if a sensitisation to intercrystalline corrosion or embrittling grain boundary decoration is to be avoided. This applies also to the heat affected zone during welding. The second one is represented by an embrittling influence of nitrogen at subzero temperatures which is actually of ductile character. In notch impact tests the ductile-to-brittle transition temperature at KV = 100 J is close to -100 °C. Therefore the new HIS may be applied in the total range of climatic temperatures but not in the deep cryogenic range. Compared to HNS of equal interstitial content the DBTT of HIS appears to be lower, though [20–22].

References 1. Nikulin I, Kaibyshev R (2011) Deformation behavior and the Portevin-Le Chatelier effect in a modified 18Cr-8Ni stainless steel. Mater Sci Engin A 528:1340–1347 2. Schmalt F (2004) Nutzung der Löslichkeit von C ? N in nichtrostenden Stählen, doctoral thesis, Ruhr University Bochum, see also Fortschr. Ber (2005) VDI 5-702, VDI Verlag, Düsseldorf 3. Tomota Y, Nakano J, Xia Y, Inoue K (1998) Unusual strain rate dependence of low temperature fracture behavior in high nitrogen bearing austenite steels. Acta Mater 46:3099–3108

References

139

4. Mughrabi H (2010) Fatigue, an everlasting materials problem—still en vogue. 10th International Fatigue Congress (Fatigue 2010). Procedia Engineering 2:3–26 5. Nebel T, Martin U, Eifler D (2001) Wechselverformungsverhalten metastabiler austenitischer Stähle. HTM 56:314–320 6. Petrov Y (2012) Surface structure of different interstitial austenitic steels after impact wear. Int J Mat Res 103:551–553 7. Jargelius-Pettersson RFA (1998) The influence of N, Mo and Mn on the microstructure and corrosion resistance of austenitic stainless steels, doctoral thesis, Royal Institute of Technology, Stockholm. ISBN 91-7170-337-3 8. Forchhammer P, Engell HJ (1969) Untersuchungen über den Lochfraß an passiven austenitischen Chrom-Nickel-Stählen in neutralen Chloridlösungen, Werkstoffe u. Korrosion 20:1–11 9. Grabke HJ (1996) The role of nitrogen in the corrosion of iron and steels. ISIJ Intern 36:777–786 10. Mudali UK, Ningshen S, Raj B (2009) Passive films and localised corrosion—role of nitrogen. In: Svyazhin AG, Prokoshkina VG, Kossyrev KL (ed) Proceedings of 10-th international conference on high nitrogen steels (HNS 09), MISIS, Moskau, 6–8 July 2009, pp 271–280 11. Bernauer J (2004) Einfluss von Kohlenstoff als Legierungselement in stickstofflegierten Chrom-Mangan Stählen, doctoral thesis, ETH Zürich, No. 15457 12. Berns H (2002) Stainless steels suited for solution nitriding. Mat-wiss u Werkstofftechn 33:5–11 13. Oppenheim R (1982) Güteeigenschaften des Superferrit X1CrNiMoNb 28-4-2 für den Chemie-Apparatebau. Thyssen Edelstahl Techn Ber 8:97–110 14. Gümpel P, Chlibec G (1985) Untersuchungen über das Werkstoffverhalten des ferritischaustenitischen Stahles X2CrNiMoN22-5. Thyssen Edelstahl Techn Ber 11:3–8 15. Riedner S, Berns H, Tyshchenko AI, Gavriljuk VG, Schulte-Noelle C, Trojahn W (2008) Nichtmagnetisierbarer warmbeständiger nichtrostender Stahl für Wälzlager. Mat-wiss und Werkstofftechn 39:448–454 16. Armstrong RW, Walley SM (2008) High strain rate properties of metals and alloys. Int Mater Rev 53(3):105–128 17. Berns H, Riedner S, Gavriljuk V, Petrov Y, Weihrauch A (2001) Microstructural changes in high interstitial stainless steels due to ballistic impact. Mater Sci Eng A 528:4669–4675 18. Shi J, Dong H, Liu YL, Gu YL, Rong F, Hui WJ, Speidel MO (2004) Ballistic behavior of nitrogen alloyed austenitic steel plates for anti-terrorist use: steel grips suppl. High Nitrogen Steels 2:239–245 19. Uggowitzer PJ, Magdowski R (1996) Nickelfree high nitrogen austenitis steels. ISIJ Intern 36:901–908 20. Harzenmoser MAE (1990) Massiv aufgestickte austenitisch-rostfrei Stähle und Duplexstähle, doctoral thesis, ETH Zürich 21. Riedner S (2010) Höchstfeste nichtrostende austenitische CrMn-Stähle mit (C ? N), doctoral thesis, Ruhr-University, Bochum 22. Bernauer J (2004) Einfluss von Kohlenstoff als Legierungselement in stickstofflegierten Chrom-Mangan Stählen, doctoral thesis, ETH Zürich, No 15457

Appendix A Tables

Table A.1 Tensile properties of HIS and reference steels at RT Steel CN85 CN85x CN96 CN107 MnCr82 CrMnN

MnC

CrNi

MnCr77

Rp0,2 [MPa] Rm [MPa] R [MPa] Au [%] A [%] Z [%] Ws [J/cm3]

370 829 1131 45 46 33 330

221 592 1930 70 83 86 422

443 881 1635 86 100 – 751

552 1000 2136 53 67 64 603

561 1002 2143 52 67 61 607

600 1020 2547 61 74 69 676

604 1075 2545 62 74 65 694

494 951 2635 68 78 68 651

626 1014 2679 44 63 77 569

Table A.2 Constants of the Ludwigson fit (Fig. 4.1b) at RT and of CN96 at subzero testing temperature TT [°C] n1 K2 eK2 -n2 Steel K1 CN85x CN96 CN107 GCN65 GCN88 GCN98 GCN115 GCN85 MnCr82 CrMnN

2260 2148 2326 1916 2254 2290 2600 2108 2343 2115

0.451 0.919 0.924 0.513 0.530 0.546 0.579 0.551 0.616 0.842

6.05 6.37 6.37 5.94 6.07 6.13 6.11 6.03 6.04 6.40

424 582 584 380 434 460 452 417 420 603

19.50 0.02 0.01 20.89 17.86 14.89 12.72 15.58 10.36 0.07 (continued)

H. Berns et al., High Interstitial Stainless Austenitic Steels, Engineering Materials, DOI: 10.1007/978-3-642-33701-7, Ó Springer-Verlag Berlin Heidelberg 2013

141

142

Appendix A

Table A.2 (continued) Steel K1 MnCr77 CrNi MnC CN96 TT 0 -30 -60 -80 -100

n1

K2

eK2

-n2

1919 1544 2754

1 0.661 0.888

6.10 5.39 5.98

446 219 395

-10-5 6.85 4.679

2222 2282 2499 2490 2617

0.927 0.936 0.963 0.949 0.940

6.44 6.58 6.66 6.75 6.85

629 718 784 851 944

0.015 0.013 0.007 0.008 0.011

Table A.3 Effect of Mo and Cu on the tensile properties of HIS at RT Steel CN96 CN94Mo1 CN96Cu2 CN107

CN103Mo1

Rp0,2 [MPa] Rm [MPa] R [MPa] Au [%] A [%] Z [%] Ws [J/cm3]

575 1060 2441 61 74 42 673

600 1020 2547 61 74 69 676

571 999 2428 58 72 43 646

544 997 2243 56 69 42 608

604 1075 2545 62 74 65 694

Table A.4 Effect of interstitial content on the tensile properties of HIS castings at RT, N & 0.6 mass % Casting Centrifugal Sand Steel Rp0,2 [MPa] Rm [MPa] R [MPa] Au [%] A [%] Z [%] Ws [J/cm3]

GCN65 422 779 1586 46 61 63 423

GCN88 501 910 1661 49 59 54 472

GCN98 491 919 1597 47 54 51 433

GCN1,15 554 1040 1823 58 65 49 585

GCN85 477 831 1278 44 46 41 329

Table A.5 Grain size of as-quenched steels, ASTM number according to ISO 643 C1 Steel ASTM Steel ASTM Number Number CN85x CN96 CN107 CN94Mo1 CN103Mo1 CN96Cu2

6 5.2 3.2 7 5 6

GCN85 GCN82 CrMnN CrNi MnC

-0.2 5.4 9 5.7 4.8

Appendix A

143

Table A.6 Tensile properties of HIS and reference steel CrMnN at subzero testing temperatures TT TT [°C] 22 0 -30 -60 -80 -100 -196 CN107

Rp0,2 [MPa] Rm [MPa] R [MPa] Au [%] A [%] Z [%] Ws [J/cm3]

604 1075 2545 62 74 65 694

644 1097 2648 58 69 63 687

693 1161 2754 58 70 64 736

775 1236 2797 55 69 61 749

878 1325 2704 52 65 61 776

926 1386 2674 52 66 58 811

1393 1613 1702 6 6 2 75

CN96

Rp0,2 [MPa] Rm [MPa] R [MPa] Au [%] A [%] Z [%] Ws [J/cm3]

600 1020 2547 61 74 69 676

652 1057 2267 57 71 70 670

733 1122 2418 54 69 66 706

783 1204 2782 56 70 61 781

857 1258 2831 54 67 62 745

945 1350 2902 50 63 63 821

1419 1679 1854 6.5 6.5 6.5 96

CN85

Rp0,2 [MPa] Rm [MPa] R [MPa] Au [%] A [%] Z [%] Ws [J/cm3]

552 1000 2136 53 67 64 603

592 1030 2313 59 75 66 698

675 1100 2412 57 72 64 723

719 1160 2915 54 69 65 736

824 1250 2704 48 63 62 730

891 1310 3026 46 61 62 739

– – – – – – –

CN85x

Rp0,2 [MPa] Rm [MPa] R [MPa] Au [%] A [%] Z [%] Ws [J/cm3]

561 1002 2143 52 67 61 607

596 1086 1923 52 65 51 667

648 1117 2259 54 69 61 706

733 1197 2519 51 66 60 733

829 1279 2546 54 72 58 856

855 1339 2393 54 67 47 836

– – – – – – –

GCN85

Rp0,2 [MPa] Rm [MPa] R [MPa] Au [%] A [%] Z [%] Ws [J/cm3]

477 831 1278 44 46 41 329

516 878 1315 48 54 42 411

553 925 1554 43 48 48 394

614 977 1687 38 42 47 400

669 1052 1623 45 47 40 431

718 1122 2311 43 43 52 421

– – – – – – –

CrMnN

Rp0,2 [MPa] Rm [MPa] R [MPa] Au [%] A [%] Z [%] Ws [J/cm3]

626 1014 2679 44 63 77 570

694 1065 – 44 64 73 609

767 1149 – 47 64 73 672

801 1223 – 47 64 63 726

917 1295 – 48 65 64 778

1012 1409 – 45 60 62 765

– – – – – – –

144

Appendix A

Table A.7 Tensile properties of CN85 at elevated temperatures, after heating the specimens were soaked for 1 h at the testing temperature TT TT (°C) Rp0,2 [MPa] Rm [MPa] Au (%) A (%) 20 300 400 500 550 600 625 650 700 1)

561 312 319 259 235 213 216/2071) 203 212

1002 760 679 637 607 549 482 479 413

52 54 39 38 33 30 29 24 19

67 70 57 47 42 37 33 27 26

Upper/ lower yield point

Table A.8 Loss of strength at slightly elevated testing temperature TT of specimens cold stretched by 20 % and aged at TA Steel CN85 CN96 TA [C°] TT [°C] Rp0,2 [MPa] Rp1 [MPa] Rm [MPa] R [MPa] Ag [%] A [%] Z [%] Ws [J/cm3]

– 20 1017 1101 1235 2532 26 37 62 441

360, 2 h 100 130 956 933 969 946 1110 1084 2148 2198 23 23 35 33 60 58 371 344

– 22 1092 1092 1247 2676 29 46 66 541

150 866 884 1053 1918 27 38 55 383

360, 2 h 100 130 964 920 950 912 1138 1109 2687 2735 33 31 46 44 67 68 498 461

150 891 886 1096 2712 31 44 67 449

Table A.9 Results of creep tests under constant temperature and stress of steel CN85, SA = solution annealed, CW = additionally cold worked e(_emin ) t(_emin ) ef tf Temperature Stress e_ min [°C] [MPa] [% / h] [%] [h] [%] [h] SA

600

650

700

350 370 400 300 350 400 270 300

3.3510-03 5.2510-03 4.3310-02 1.1410-06 6.1410-02 1.5110-01 8.1210-02 1.4710-01

1.78 1.64 3.42 2.54 5.39 5.58 1.76 1.95

47.99 24.21 4.69 91.78 5.84 1.80 5.40 4.84

6.48 6.06 9.37 6.32 8.42 9.12 19.95 12.2

239.39 161.4 27.61 420.15 35.23 14.49 74.15 29.63 (continued)

Appendix A

145

Table A.9 (continued) Temperature [°C] CW

600 650

700

Stress [MPa]

e_ min [% / h]

e(_emin ) [%]

t(_emin ) [h]

370 400 320 350 370 250 270

2.7410-03 3.4410-03 6.8710-03 8.4110-03 7.0610-03 2.7710-02 2.2010-02

0.18 0.32 0.27 0.13 0.31 0.41 0.12

18.48 14.79 16.67 5.45 3.60 15.13 8.64

tf [h]

ef [%] 0.80 1.04 1.75 0.52 1.54 9.57 6.94

173.42 86.88 106.58 44.20 18.10 105.84 64.19

Table A.10 Vickers hardness measured at room temperature: macrohardness HV30, microhardness HV0.1 on fracture face after tensile test at RT Steel HV30 HV0.1 HV30 CN85 head1) gauge2) CN96

CN107 CN94Mo1 CN103Mo1 CN96Cu2 GCN65 GCN88 GCN98 GCN115 GCN85 CrNi MnC CrMnN MnCr82 MnCr77 1)

237

270

270 277 276 254 264 265 271 278 244 141 211 262 248 212

619

after tensile test at °C 500 550 600 257 263 255 476 441 415 after % elongation by cold drawing 16 20 28 360 363 400

590

650 250 387

700 256 371

37 423

630

460 741 526 566 547

undeformed specimen head,

2)

within deformed gauge length

Table A.11 ISO-V notch impact toughness KV at RT and ductile-to-brittle transition temperature DBTT at KV = 100 J, in brackets = specimens of transverse taking Steel KV [J] DBTT [°C] CN85x CN85 CN96

272 325 (233) 364 (296)

-88 -94 (-112) -91 (-49) (continued)

146

Appendix A

Table A.11 (continued) Steel

KV [J]

DBTT [°C]

CN107 CN94Mo1 CN103Mo1 CN96Cu2 GCN65 GCN88 GCN98 GCN115 GCN85 CrMnN MnCr82 CrNi MnC

333 (297) 271 271 207 250 265 317 249 290 313 187 2881) 280

-49 (-53) -93 -87 -96 -47 -94 -21 -77 -1961) -59

1)

After G. Kalla and G. Uhlig, Bänder, Bleche, Rohre (1992) 136-142

Table A.12 Results of impact wear tests: wear rate between 3000 and 12000 impacts, surface hardness and hardness penetration measured in a section normal to the wear surface Steel Total mass loss Wear rate Surface hardness Hardness penetration [mg] [lg/impact] [HV0.05] [mm] CN85 CN96 CN107 CrMnN MnC MnCr77 CrNi

48.5 61.5 60 53 58 63 116

4.4 4.7 5.3 4.6 5.2 5.7 9.6

– 613 726 538 761 547 524

– 0.4 0.4 0.4 0.9 0.3 1.1

Table A.13 Effect of molybdenum and copper on mass loss Dm, corrosion rate v and removal w during uninterrupted submersion tests in aqueous solutions lasting for 120 h CN96 CN94Mo1 CN96Cu2 CN107 CN103Mo1 10 %H2SO4

1 %HCl

3 %HCl

Dm [g] v [g/m2h] w [mm/a] Dm [g] v [g/m2h] w [mm/a] Dm [g] v [g/m2h] w [mm/a] density [g/cm3]

16.7670 131.2 150.3 1.9125 19.4 22.3 16.7284 164.4 188.3 7.65

8.7784 36.1 41.2 4.0429 42.6 48.7 13.7921 115.5 132 7.67

0.6995 5.7 6.5 3.7070 38 43.4 13.5929 115.4 131.9 7.66

2.2984 17.9 20.7 0.9131 8.9 10.3 16.1793 157.7 181.7 7.60

0.0922 0.76 0.88 2.0997 21.5 24.7 13.2903 113.9 130.9 7.62

Appendix A

147

Table A.14 Effect of carbon on mass loss Dm, corrosion rate v and removal w during uninterrupted submersion tests in aqueous solutions lasting for 120 and 600 h, respectively 120 h GCN65 GCN88 GCN98 GCN115 10 %H2SO4 1 %HCl

3 %HCl

600 h 10 %H2SO4

Dm [g] v [g/m2h] w [mm/a] Dm [g] v [g/m2h] w [mm/a] Dm[g] v [g/m2h] w [mm/a]

0.1285 1.22 1.4 0.9632 9.2 10.5 12.7642 121.5 139.2

0.3901 3.25 3.7 0.8582 7.1 8.2 14.2557 118.9 136.2

0.1131 1 1.1 0.6578 5.5 6.3 13.5281 112.8 129.2

0.1138 1 1.1 0.4318 3.6 4.1 12.7844 106.8 122.2

Dm [g] v [g/m2h] w [mm/a] density [g/cm3]

0.1248 0.24 0.27 7.65

0.2421 0.40 0.46 7.65

0.1489 0.25 0.29 7.65

0.1177 0.2 0.22 7.65

Table A.15 Characteristic data of current density/SHE potential tests at room temperature in (a) 0.5 m H2SO4, b 3 % NaCl (see Fig. 4.21) (a)

UR [V]

UPP UP [lA/cm2]

CN85x CN96 CN107 CN94Mo1 CN103Mo1 CN96Cu2 GCN65 GCN88 GCN98 GCN15 GCN85 CrMnN CrNi

-0.30 -0.24 -0.24 0.11 -0.025 -0.11 -0.07 -0.07 -0.07 -0.07 -0.31 -0.04 -0.13

-0.25 -0.22 -0.22 -0.10 -0.25 -0.08

(b) CN85x CN96 CN107 CN94Mo1 CN103Mo1 CN96Cu2

UR [V] -0.23 -0.39 -0.31 -0.37 -0.27 -0.29

UP [V] -0.19 -0.16 -0.10 -0.16 -0.16

0.21 0.21 0.22 0.31 0.29 0.22 0.22 0.22 0.22 0.22 0.21 0.19 0.28

UA

UB

URP

UB’

iPP

iP

i0

iRP

iB’

0.29 0.40 0.44 0.45 0.40 0.45 0.29 0.36 0.36 0.36 0.28 0.47 0.4

1.08 1.05 1.10 1.10 1.09 1.15 1.11 1.11 1.11 1.11 1.12 1.04 1.09

1.40 1.32 1.31 1.33 1.31 1.33 1.32 1.32 1.32 1.32 1.32 1.40 -

1.63 1.64 1.63 1.63 1.61 1.65 1.63 1.61 1.61 1.61 1.62 1.61 -

30 54 69 16 10 45 17 2

11000 2000 460 54 49 100 71 63 8900 70

2.1 6.2 3.2 6.8 2.8 2.8 1.5 1.5 1.5 1.5 2.2 2 1.8

2650 3100 4200 3400 4000 2800 930 4300 4940 4980 2950 1950 -

40 50 70 110 77 29 20 65 40 100 92 20 -

UA [V] -0.05 0.05 -0.05 -

UB [V] -0.19 (0.69) 0.20 0.91 0.44 0.88 -0.04

2

ip [lA/cm2]

i0 [lA/cm2]

3 3.2 0.55 0.38 3.1

80 2.1 2 0.43 1 4 (continued)

148

Appendix A

Table A.15 (continued) (b) UR [V] UP [V]

UA [V]

UB [V]

ip [lA/cm2]

i0 [lA/cm2]

GCN65 GCN88 GCN98 GCN115 GCN85 MnCr821) CrMnN CrNi

-0.05 -0.07 -0.02 -0.13 0.10 0.08

0.18 1.24 1.25 1.28 0.02 (0.65)2) -0.03 0.74 0.26

7.6 4.4 5.6 1.3 3.1 0.9

4 1.4 1.4 1 22 3 0.8

1)

-0.34 -0.34 -0.43 -0.51 -0.20 -0.53 -0.18 -0.10

Taken from [1.5]

-0.11 -0.09 -0.14 -0.28 0.03 0.06 2)

In brackets: derived from a linear plot

Table A.16 Temperature TSA and time tSA of solution annealing applied to the new HIS, pN2 is the nitrogen pressure calculated by Thermo-Calc to retain the level of nitrogen at the surface tSA [min] pN2 [mbar] Steel TSA [°C] CN96 CN107 CN85 CN94Mo1 CN103Mo1 CN96Cu2 GCN65 GCN88 GCN98 GCN115 MnCr82 CrMnN CrNi MnC

1100 1175 1100 1125 1175 1200 1100 1150 1175 1200 1100 1050 1100 1050

30 30 30 30 30 30 45 45 45 45 30 45 30 30

62.2 106.5 59.8 72.8 114.6 342.7 42.5 76.8 76.6 150.3 -

Table A.17 Steel CN96 cold drawn up to 37 % reduction of cross-section. a Results of tensile tests at RT, b Costants of the Ludwigson fit (a) Reduction [%] Rp0.2 [MPa] Rp1 [MPa] Rm [MPa] R [MPa] Au [%] A [%] Z [%] Ws [J/cm3] Rp0.2/Rm [-]

0 600 606 1020 2547 61 74 69 676 0.59

16 810 914 1164 2699 32 45 63 501 0.70

20 988 1103 1243 2624 26 39 59 469 0.79

29 1103 1241 1320 2825 30 35 58 453 0.84

37 1190 1379 1410 2766 12 27 58 368 0.84 (continued)

Appendix A

149

Table A.17 (continued) (a) Reduction [%] (b) Reduction [%] 0 16 20 29 37

Constants K1 2411 2224 2216 2166 2038

n1 0.518 0.283 0.233 0.183 0.117

K2 6.10 6.06 6.03 5.93 5.65

eK2 445 428 415 377 284

-n2 14.94 21.90 24.23 27.97 39.80

Table A.18 Liquidus temperature TL, solidus temperature TS and the begin of precipitation at TP of HIS listed in Table 2.1 (see Figs. B1 to B10) Steel TL [°C] TS [°C] TP [°C] CN85x CN96 CN107 CN94Mo1 CN103Mo1 CN96Cu2 GCN65 GCN88 GCN98 GCN115

1364 1362 1351 1362 1351 1352 1385 1364 1355 1349

1307 1292 1256 1289 1258 1269 1344 1305 1270 1251

1006 1047 1146 1031 1119 1139 981 1006 1118 1167

Appendix B Figures

H. Berns et al., High Interstitial Stainless Austenitic Steels, Engineering Materials, DOI: 10.1007/978-3-642-33701-7, Ó Springer-Verlag Berlin Heidelberg 2013

151

152

Appendix B

CN85x

(a) 1500

temperature [°C]

0.8

L

1400

L+F

1300

1.2 p N2 [bar] L+G

L+F+A

F

1200

A

F+A

1100

A+M2 N

1000 900

A+M2N+M23 C6

800 A+M2N+M 23 C6 + σ

700 0

0.3

0.6 0.9 (C+N) [mass%]

1.2

1.5

N [mass%]

(b) 1500

temperature [°C]

1400

0.8

0.6

0.4

0.2

L+F+G L+G

L

1300

F+A

L+A

1200

A

CN107

1100 1000

A+M2N

0

1p N2 [bar]

A+M 23 C 6

CN96 CN85

900 A+M2N+M23 C6

800 700

A+M 2N+M 23 C6 + σ 0

0.2

0.4 C [mass%]

0 0.11 0.250.430.67 1

0.6

1.5 2.3

0.8 4

9

∞

C/N

(c)

1.0 A

A

L

mole fraction [ - ]

0.8 0.6 0.4 0.2

M2N

M23 C6

G

σ

0 700

900 1100 1300 temperature [°C]

1500

Fig. B.1 Figs. B.1–B.10 Calculated phase diagrams for the actual chemical composition of HIS listed in Table 2.1, a Plot over C+N in which the dotted line represents the steel content, b Plot over C/N in which the dotted line represents the C/N ratio of the steel and (+) in Fig. B.1b the TP of three steels for comparison, c Mole fraction in dependence of temperature. The target phase field of homogeneous austenite extends from the solidus temperature TS to the begin of precipitation at TP, which are listed together with the liquidus temperature TL in Table A.18

Appendix B

153 CN96

(a) 1500

0.8 L

temperature [°C]

1400

1.2 [bar]

L+G

L+F L+F+A

F

1300

L+A

1200 A

F+A

1100 1000

A+M2N

900

A+M2N+M23C6

800 A+M2N+M23C6 + σ

700 0

0.3

0.6 0.9 (C+N) [mass%]

1.2

1.5

0.2

0

N [mass%] 0.8

(b) 1500

0.4 1 pN2 [bar]

L+G

1400

temperature [°C]

0.6

L+F

L

1300

L+A

1200

A

1100 1000

A+M23C 6 A+M2N

900

A+M2N+M23 C6

800

A+M2N+M23 C6 + σ

700 0

0.2

0.4 C [mass%]

0 0.11 0.250.430.67 1

0.6

1.5 2.3

0.8 4

9

∞

C/N

(c)

1.0 A

A

L

mole fraction [ - ]

0.8 0.6 0.4 0.2 M2N

0 700

Fig. B.2

M23C 6

G

σ

900 1100 1300 temperature [°C]

1500

154

Appendix B

CN107

(a) 1500

0.8 1.2 pN2 [bar]

L

temperature [°C]

1400

L+G

L+F L+F+A

1300 F

L+A A

1200

F+A

1100 A+M23 C6

1000 F+A +M23 C 6

900

A+M2N+M23 C6

800 A+M2N+M23 C6 + σ

700 0

0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 (C+N) [mass%] N [mass%] 1

(b) 1500 1400

temperature [°C]

0.8

0.6

0.4

0

L

1300

L+A

1200

A

1100

A+M23 C6

1000

A+M2N

900

A+M2N+M23 C6 A+M2N+M23 C6 + σ

800 700

0.2

1 pN2 [bar]

L+G

0

0.2

0.4

0.6

0.8

1

C [mass%] 0 0.11 0.250.430.67 1

1.5 2.3

4

9

∞

C/N

(c)

1.0 A

A

L

mole fraction [ - ]

0.8 0.6 0.4 0.2 0

M23 C6

700

Fig. B.3

M2N

G

σ

900 1100 1300 temperature [°C]

1500

Appendix B

155 CN94Mo1

(a) 1500 L

temperature [°C]

1400

L+G 1.2 p N2 [bar]

0.8

L+F L+F+A

1300

L+A

1200

A

F+A

1100 1000

C6 A+ M 23

900

A+M2N+M23C6

800 A+M2N+M 23C6+ σ

700 0

0.3

0.6 0.9 (C+N) [mass%]

1.2

1.5

0.2

0

N [mass%] 0.8

(b) 1500

0.4 1 pN2 [bar]

L+G

1400

temperature [°C]

0.6

L

1300

F+L

L+A

1200

A

1100 A+M23C6

1000 A+M2N

900 A+M2N+M23C6

800 A+M2N+M23 C6 + σ

700 0

0.2

0.4 0.6 C [mass%]

0 0.11 0.250.430.67 1

1.5 2.3

0.8 4

9

∞

C/N

(c)

1.0

mole fraction [ - ]

A

A

0.8

L

0.6

0.4 0.2 M2N

700

Fig. B.4

M 23C 6

G

σ

0

900 1100 1300 temperature [°C]

1500

156

Appendix B CN103Mo1

(a) 1500 0.8

L+F

1300 temperature [°C]

L+G

L

1400

L+F+A

F

1.2 pN2 [bar]

L+A

1200

A

F+A

1100 A+M23C 6

1000 900

A+M23C6 +σ

800

A+M2N+M23 C 6

A+M2N+M23C6+ σ

700 0

0.3

0.6

0.9

1.2

1.5

0.2

0

(C+N) [mass%]

(b) 1500

1

1 p [bar] N2

L+G

1400 temperature [°C]

N [mass%] 0.6 0.4

0.8

L

1300

L+A A

1200 1100

A+M23 C 6

1000

A+M2N

900

A+M2N+M23C 6

800

A+M2N+M23 C6 +σ

700 0

0.2

0.4

0.6

0.8

1

C [mass%] 0 0.11 0.25 0.43 0.67 1

1.5 2.3

4

9

∞

C/N

(c)

1.0 A A

L

mole fraction [ - ]

0.8 0.6 0.4 0.2 M2N

0 700

Fig. B.5

M23C 6

G

σ

900 1100 1300 temperature [°C]

1500

Appendix B

157 CN96Cu2

(a) 1500

0.8 1.2 pN2[bar] L+G

L

1400

L+F L+F+A

temperature [°C]

1300

L+A

F

1200

A

F+A

1100

A+M23C 6

1000 A+M2N+M23C6

900 800

A+M2N+M23C6+σ

700 0

0.3

0.6

0.9

1.2

1.5

0.2

0

(C+N) [mass%]

(b) 1500

1p

L+G

1400 temperature [°C]

N [mass%] 0.6 0.4

0.8

N2 [bar]

L

1300

L+A A

1200

A+M23C 6

1100

A+M2N

1000 900

A+M2N+M23C6

800 700 0

0.2

0.4

0.6

0.8

C [mass%] 0 0.11 0.25 0.43 0.67 1

(c)

1.5 2.3

4

9

∞

C/N

1.0 A

A

L

mole fraction [ - ]

0.8 0.6 0.4 0.2

M23C 6 σ

M2N

G

0 700

Fig. B.6

900 1100 1300 temperature [°C]

1500

158

Appendix B

(a)

GCN65 1500

L

1400

1.2 p N2 [bar] 0.8

L+F

L+G

temperature [°C]

1300 1200

A F+A

1100 1000 A+M2N

900 A+M2N+ σ

800

A+M2N+M 23C 6+σ

700 0

(b)

0.3

0. 6

1500

0.6 0.9 (C+N) [mass%]

0. 5

0. 4

N [mass%] 0. 3 0. 2

1.5

0. 1

0

1 [bar] L

1400

L+F

1300 temperature [°C]

1.2

L+F+A F+A

1200

F+A+M23 C 6

1100 A

1000 900

A+M23 C 6 A+M2N A+M2N+M23 C6

800

A+σ

A+M2N+M23C6+σ

700 0

0.1

0

(c)

0.2 0.3 0.4 C [mass%]

0.11 0.25 0.43 0.67

1 1.5 C/N

0.5 2.3

0.6

4

9

1.0 L

A A

0.8 mole fraction [ - ]

∞

0.6 0.4 F

0.2

F

σ

0

M 23C 6 700

M2N 900

A 1100

1300

temperature [°C]

Fig. B.7

G 1500

Appendix B

159 GCN88

(a)

1500

0.8

L

temperature [°C]

1400

L+F

L+F+A L+A

F

1300

1.2 p N2 [bar] L+G

1200

A F+A

1100

A+M2N

1000 900

A+M2N+M23C6

800

A+M2N+M23 C6 +σ

700

0

0.3

0.6

0.9

1.2

1.5

(C+N) [mass%] N [mass%]

(b)

0.8 0.7 0. 6 0. 5 0.4 0.3 0. 2 0. 1 1500 L+G

temperature [°C]

1400

0

1 p N2 [bar] L L+F

1300

L+F+A

1200

A

1100 A+M23C6

1000

A+M2N

900

A+M2N+M23 C6

800 A+M2N+M23 C6 +σ

700

0

0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 C [mass%]

0

(c)

0.11 0.25 0.43 0.67

1 1.5 C/N

2.3

4

9

∞

1.0 L

0.8 mole fraction [ - ]

A

A

0.6 0.4

0.2 F

σ M2N M C 23 6

G

0

700

900

1100

1300

temperature [°C]

Fig. B.8

1500

160

Appendix B GCN98

(a)

1500

0.8 1.2 p [bar] N2

L

temperature [°C]

1400

L+G

L+F F

1300

L+F+A

1200

L+A A

F+A

1100 A+M23 C6

1000

F+A +M23 C6

900 800

A+M2 N+M23 C6

A+M2N+M23 C6 +σ

700 0

0.3

0.6

0.9

1.2

1.5

0.2

0

(C+N) [mass%] N [mass%]

(b)

0.8

0.6

1500

0.4 1 pN2 [bar]

L+G

L

temperature [°C]

1400 1300

L+F L+F+A

L+A

1200

A

1100 A+M23C6

1000

A+M2N

900

A+M2N+M23 C6

800 A+M2N+M23C6 + σ

700 0

0.2

0.4

0.6

0.8

C [mass%] 0

(c)

0.11 0.25 0.43 0.67

1 1.5 C/N

2.3

4

9

∞

1.0 L A

A

mole fraction [ - ]

0.8

0.6 0.4

0.2 M2N M23 C6

G

σ

0 700

900

1100

1300

temperature [°C]

Fig. B.9

1500

Appendix B

161 GCN115

(a)

1500

temperature [°C]

0.8 1.2 p [bar] N2 L+G

L

1400

L+F

1300

L+F+A

F

L+A A

1200

F+A

1100 A+M23 C6

1000

F+A +M 23C6

900

A+M2N+M23 C6

800 A+M2N+M23 C6 +σ

700 0

(b)

0.3

1500

1.5

0.2

0

1 p N2 [bar]

L+G

L

1300

L+F

L+A

1200

A

1100

A+M23 C6 A+M2N

1000 900

A+M2N+M23 C6

800

A+M 2N+M 23 C6 +σ

700

1.2

N [mass%] 0.6 0.4

0.8

1 1400

temperature [°C]

0.6 0.9 (C+N) [mass%]

0

0.2

0.4

0.6

0.8

1

C [mass%] 0

(c)

0.11 0.25 0.43 0.67

1 1.5 C/N

2.3

4

9

∞

1.0 L A

mole fraction [ - ]

0.8

0.6 0.4

0.2

M2N M23 C6

σ

0 700

Fig. B.10

A

900

1100 1300 temperature [°C]

G

1500

162

Appendix B

Fig. B.11 Content of pseudo-pearlite after solution annealing and isothermal holding at temperature Tis for 1 h and subsequent quenching in water

About the Authors

Hans Berns born 1935, studied ferrous metallurgy, worked in the speciality steel industry from 1959–1979, doctorate TH Aachen l964, habilitation TU Berlin l975, held the Chair of Materials Technology at the Ruhr University, Germany, from 1979–2000, since then working as emeritus. Valentin Gavriljuk born 1938, studied metal physics, worked at the Institute for Metal Physics, Kiev, Ukraine as post-graduate (l962), senior scientific researcher (1973), leading scientific researcher (1986), head of department of physical principles for design of steel and alloys since l989, doctorate l965, habilitation l986. Sascha Riedner born l975, studied mechanical engineering at the Ruhr University Bochum, Germany, majoring in engineering materials, diploma 2004, doctorate 2009, since then working in R+D and customer services at the German Specialty Steel Company (DEW) in Hagen.

H. Berns et al., High Interstitial Stainless Austenitic Steels, Engineering Materials, DOI: 10.1007/978-3-642-33701-7, Ó Springer-Verlag Berlin Heidelberg 2013

163

Index

A Ab initio calculation, 22 Adiabatic shear bands, 66, 135 Aging, 123, 134 Amorphous structure, 71 Application, 134 Atomic bonds, 21 configuration, 22, 30 distribution, 28, 46, 127 interaction, 30 Austenite homogeneous, 7, 22, 116 retained, 34 stability, 128

B Ballistic impact, 65, 135 Bearing ring, 134 Body friendly, 137 Boiling, 9 Brittleness. See Embrittlement

C C/N ratio, 9, 45 C+N influence constitution, 8, 118 electron structure, 24, 29, 128 stacking fault energy, 45 strength, 88 Carbides, 8 Casting, 113 Chemical composition, 18

Chemical homogeneity. See Homogeneity Cleanness, 116 Cleavage-like fracture, 61 Clustering, 28, 36, 127 Cluster size, 41, 43, 127 Cold drawing, 123 Cold manufacture, 133 Cold worked structure, 51 Cold work hardening, 56 Constitution, 7 C/N ratio, 9, 111, 152 interstitial content, 8 steels investigated, 149 substitutional content, 11 tramp elements, 16 Contra HIS, 137 Cooling time, 120, 132 Coordination spheres, 30 Corrosion, 99, 130, 134 current density/potential, 102, 147 intercrystalline, 106, 118, 120 submersion, 100, 146 Creep, 91 Cross-section, 111, 120, 132 Cyclic loading, 68

D DBTT, 62, 93, 129, 138, 145 Degassing, 9 Dilatometry, 35 Dislocation barrier, 64 density, 74 pile up, 48

H. Berns et al., High Interstitial Stainless Austenitic Steels, Engineering Materials, DOI: 10.1007/978-3-642-33701-7, Ó Springer-Verlag Berlin Heidelberg 2013

165

166

D (cont.) pinning, 48 splitting, 44, 51

E Electron spin resonance, 26, 38 Electron structure, 4, 22, 127 Elevated temperature, 88, 137 Embrittlement, 2, 129, 136 cleavage-like, 61 covalent, 21 precipitation, 120 ESR, 112, 116

F Fatigue push/pull, 73 repeated impacts, 68 rotating bending, 94, 129 Fermi level, 23 Ferrite-d, 9 Foaming, 9 Forging, 115 Fractography, 90, 95 Fracture energy, 87, 128, 134, 136 Free electron, 21, 38, 127 concentration, 27 distribution, 24

G Grain size, 49, 116, 142

H Hadfield steel, 17, 36, 47, 52 Hall-Petch equation, 49 Hardness, 58, 93, 135, 145 Heat treatment, 116 High intersticial steels (HIS), 2, 17 High nitrogen steels (HNS), 1 Homogeneity nano scale, 36, 128 segregation, 28, 114 Hot manufacture, 132 Hot working, 115

I Impact toughness, 93, 123, 145 Ingot, 112 Interaction energy, 33

Index Interatomic bonds, 21, 24, 127 Interstitial clouds, 49 content, 8 Iron structure, 24

L Liquidus temperature, 112, 149 Ludwigson equation, 86, 141 Ludwik equation, 51

M Machining, 125 Magnetic measurement, 36, 41, 127 Magnetic properties, 106 Manufacture, 111, 131 Martensite (a), 32, 34, 70 Martensite (e), 50, 54, 57, 66, 74 Melting, 112 Microstructure. See Structure Mössbauer spectroscopy, 29 Multiscale approach, 4

N Nanocrystals, 71 Nitrides, 8, 66, 75, 118 Nitrogen dislocation, 128 DOS, 24 grainboundary, 49, 75 pressure, 8 solubility, 9 yielding, 48 Nonmagnetic, 131, 134

O Ordering, 29, 63, 127

P Paramagnetic, 41 Phase diagrams, 7, 152 Phases encountered, 7 Planar slip, 54 Precipitation dissolution, 116 during fatigue, 75 during impact, 66 grainboundary, 91, 120 range, 132 temperature, 112, 149

Index Pro HIS, 137 Properties, 85 corrosion, 99 magnatic, 106 mechanical, 85, 128 wear, 95 Pseudo pearlite, 115, 118

Q Quenching, 132, 133 continuous, 120 interrupted, 118

R Reference steels, 17 Refractories, 114 Remelting, 112, 116 Retained austenite, 34 Retaining rings, 137 Ring rolling, 115 Rolling, 115

S Scheil simulation, 113 Segregation, 28, 111, 114 Selection of steel, 17 Sensitization, 120 Serration, 88, 129 Short range decomposition, 28, 36 Short range ordering, 28, 36, 63, 127 Sigma phase, 7, 12, 15 Solidification, 7, 113 Solidus temperature, 112, 149 Solution annealing, 116, 148 Stacking fault energy, 44, 50, 127

167 Steels investigated, 18 Strain rate, 61 Stress-strain curves, 86, 124, 136 Structure, 21, 127 after loading, 47 as-quenched, 22 Subzero temperature, 59, 88, 89, 129 Superparamagnetic clusters, 41, 43 Surface, 114, 117

T Tensile properties, 47, 85, 128, 141 comparison, 87 creep, 91, 144 elevated temperature, 88, 129, 144 strain rate, 61, 129 subzero temperature, 59, 88, 129, 143 Thermo-Calc, 7 Thermodynamic stability, 23, 29, 34, 46 Tramp elements, 16 TRIP, 4 Twinning, 47, 53, 58 TWIP, 3, 4, 58

W Wear, 95, 136 abrasion, 96 cavitation, 99 impact, 68, 96, 146 Welding, 124 Work hardening, 47, 52, 86, 133 Work hardening mechanism, 56

Y Yielding, 48