Perovskite Ceramics: Recent Advances and Emerging Applications (Elsevier Series on Advanced Ceramic Materials) [1 ed.] 0323905862, 9780323905862

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Perovskite Ceramics

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Elsevier Series in Advanced Ceramic Materials

Perovskite Ceramics Recent Advances and Emerging Applications

Edited by

 Luis Clabel Huama´n Jose Victor Anthony Garcia Rivera

Elsevier Radarweg 29, PO Box 211, 1000 AE Amsterdam, Netherlands The Boulevard, Langford Lane, Kidlington, Oxford OX5 1GB, United Kingdom 50 Hampshire Street, 5th Floor, Cambridge, MA 02139, United States Copyright © 2023 Elsevier Ltd. All rights reserved. No part of this publication may be reproduced or transmitted in any form or by any means, electronic or mechanical, including photocopying, recording, or any information storage and retrieval system, without permission in writing from the publisher. Details on how to seek permission, further information about the Publisher’s permissions policies and our arrangements with organizations such as the Copyright Clearance Center and the Copyright Licensing Agency, can be found at our website: www.elsevier.com/permissions. This book and the individual contributions contained in it are protected under copyright by the Publisher (other than as may be noted herein). Notices Knowledge and best practice in this field are constantly changing. As new research and experience broaden our understanding, changes in research methods, professional practices, or medical treatment may become necessary. Practitioners and researchers must always rely on their own experience and knowledge in evaluating and using any information, methods, compounds, or experiments described herein. In using such information or methods they should be mindful of their own safety and the safety of others, including parties for whom they have a professional responsibility. To the fullest extent of the law, neither the Publisher nor the authors, contributors, or editors, assume any liability for any injury and/or damage to persons or property as a matter of products liability, negligence or otherwise, or from any use or operation of any methods, products, instructions, or ideas contained in the material herein. ISBN: 978-0-323-90586-2 For information on all Elsevier publications visit our website at https://www.elsevier.com/books-and-journals

Publisher: Matthew Deans Acquisitions Editor: Gwen Jones Editorial Project Manager: Rafael Guilherme Trombaco Production Project Manager: Anitha Sivaraj Cover Designer: Matthew Limbert Typeset by STRAIVE, India

Dedication

To my parents, Marcelina Huama´n and Juan Clabel, for their endless love, support, and encouragement. —Jose Luis Clabel Huama´n To Max, Nina, Sami, and Maritza —Victor Anthony Garcia Rivera In memory of Prof. Josue Olortegui Obregon —Jose Luis Clabel Huama´n and Victor Anthony Garcia Rivera

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Contents

Contributors Preface

Part One 1

2

3

xiii xvii

Advanced research

Introduction to perovskites Luis Ortega-San-Martin 1 What is a perovskite? 2 Symmetry matters: From the bumped journey that started with the pursuit of the appropriate symmetry of the mineral perovskite to the search for modern lasers 3 An incredible structural versatility that can lead us to the next generation of solid-state refrigerators: Hybrid perovskites 4 Chemical flexibility: Toward the next generation of electrocatalyzers proposed by modern machine learning tools 5 Never abandon your perovskite; others might find it useful 6 Conclusions References Methods for the synthesis of ceramic materials with perovskite structure Jos e Luis Clabel Huama´n, J.C. Sczancoski, Euclydes Marega, Jr. and Alexandre H. Pinto 1 Introduction 2 Synthesis and processing of perovskite ceramics 3 Conclusions References Antiferrodistortive phase transition in doped strontium titanate ceramics: The role of the perovskite lattice vacancies Alexander Tkach 1 Introduction 2 Phase transition of undoped ST ceramics 3 Effect of isovalent doping on the phase transition 4 Acceptor doping/oxygen vacancy effect 5 Donor doping/strontium vacancy effect 6 Conclusions References

1 3 6

8 12 17 21 24 24 31

31 32 62 62 77 77 78 82 85 90 94 94

viii

4

5

6

Contents

Polaronic effects in perovskite oxides Marius Adrian Husanu and Dana Georgeta Popescu 1 Introduction 2 Response of Frohlich and Holstein polarons to external fields 3 Electron-phonon interaction (EPI) in transition metal oxides 4 Outlook and perspectives Acknowledgments References Further reading Compensation and screening of ferroelectricity in perovskite oxides Dana Georgeta Popescu and Marius Adrian Husanu 1 Introduction 2 Proper ferroelectrics 3 Improper ferroelectrics 4 Rashba ferroelectrics 5 Conclusions and outlook Acknowledgments References Crystal structures of copper oxide-based perovskite compounds Takeo Oku 1 Introduction 2 Atomic observation by HREM 3 Crystal structures of Tl-based copper oxides 4 Modulated superstructures of Tl-based copper oxides 5 Crystal structures of Pb-based copper oxides 6 Crystal structures of Hg-based copper oxide 7 Crystal structures of Bi-based copper oxide 8 Structures of lanthanoid-based copper oxides 9 Oxygen ordering in YBa2Cu3O7 x 10 Y-based copper oxides with high JC 11 CO3- and BO3-based copper oxides 12 Defects, interfaces, and surface structures 13 Summary Acknowledgment References

99 99 102 108 116 117 117 124 125 125 127 136 142 145 146 146 155 155 156 158 163 168 175 180 182 188 191 196 197 206 207 207

Contents

Part Two Oxide and halide perovskites: Functional and advanced applications 7

8

9

10

Perovskite-structured ceramics in solid oxide fuel cell application Nurul Akidah Baharuddin, Hamimah Abd Rahman, Abdullah Abdul Samat, Nafisah Osman, Nur Syafkeena Mohd Affandi, and Suhaida Dila Safian 1 Introduction 2 Perovskite-structured electrolyte for SOFCs 3 Perovskite-structured electrodes for SOFCs 4 Perovskite-structured coating for the SOFC interconnect 5 Summary References Perovskite membranes for oxygen separation Daniel Dornellas Athayde, Julius Motuzas, and Wander Vasconcelos 1 Introduction 2 Perovskite structure and separation phenomena 3 Membrane synthesis and preparation methods 4 Performance and applications 5 Future perspectives 6 Final remarks References Perovskite lead-free dielectric ceramics: Highly promising materials for energy storage applications Mrinal Kanti Adak and Debasis Dhak 1 Introduction 2 Energy storage measurement principle 3 Energy storage performance of lead-free dielectrics 4 Breakdown mechanism in dielectric materials 5 Conclusion and future prospectus Acknowledgment References Perovskite-like structure ceramic materials and their design for electrical applications Armando Reyes-Montero, Rosalba Castan˜eda-Guzma´n, Marı´a Elena Villafuerte-Castrejo´n, Jos e A´lvaro Cha´vez-Carvayar, and Lorena Pardo 1 Introduction 2 The structural link to electrical properties 3 Properties of electroceramics 4 Lead-free piezoceramic materials 5 Processing techniques of polar electroceramics

ix

219 221

221 224 230 239 249 249 263 263 265 272 283 288 289 289 295 295 297 300 307 308 309 309 317

317 318 319 322 326

x

11

12

13

14

Contents

6 Electrical properties evaluation: Piezoelectric applications 7 Summary Acknowledgments References

328 331 331 331

Multiferroic perovskite ceramics: Properties and applications Jos e Luis Clabel Huama´n, Victor Anthony Garcia Rivera, Alexandre H. Pinto, and Euclydes Marega, Jr. 1 Introduction to multiferroic perovskite ceramics 2 Theory of magnetoelectric properties in ceramics 3 Applications of magnetoelectric composite ceramics References

339

Functional double perovskites from bulk to thin film heterostructures: The example of La2NiMnO6 Jasnamol P. Palakkal and Lambert Alff 1 Introduction 2 Functional properties of double perovskites 3 Summary Acknowledgment References

339 340 368 373 383 383 384 392 393 393

Perovskite-based emerging memories Firman Mangasa Simanjuntak, Tahta Amrillah, A. Syed Jalaluddeen, V. Bipin, and Suresh Kumar Garlapati 1 Introduction 2 FeRAM technologies 3 MRAM technologies 4 ReRAM technologies 5 Summary, challenges, and outlook Acknowledgments Authors’ contributions References

401

Perovskite-based light-emitting diodes Joni Welman Simatupang, Firman Mangasa Simanjuntak, and David James Tyler 1 Introduction to LED devices 2 State-of-the-art LED structures and technologies 3 Dielectric-optical properties and applications of perovskite-based LEDs 4 Challenges and performances 5 Roadmap of perovskite-based LEDs 6 Conclusions Acknowledgment References

485

401 401 426 446 456 462 462 462

485 487 497 507 510 512 513 513

Contents

15

16

Light storage perovskites: Synthesis, mechanisms, and applications Victor Vendruscolo, Douglas Lourenc¸o Fritzen, Elaine Andrade de Mattos, and Lucas Carvalho Veloso Rodrigues 1 Light storage materials 2 Synthesis and designing of persistent luminescent perovskites 3 Persistent luminescent mechanism 4 Applications of persistent luminescent perovskites 5 Conclusions References Halide-based perovskites in photonics: From photocatalysts to highly efficient optoelectronic devices Luan Passini, Jeferson Almeida Dias, Giovanna Ferreira Bigotto Gonc¸alves, Sajjad Ullah, Elias Paiva Ferreira Neto, and Danilo Manzani 1 Introduction 2 Synthesis methods of MHPs 3 Luminescent perovskites: Photoexcitation and emission processes of halide-based nano/QD perovskites 4 Perovskite solar cells (PSCs) 5 Photocatalysis by MHP 6 Conclusions Acknowledgments References

Index

xi

517

517 523 526 534 540 540 547

547 549 555 559 575 585 586 587 601

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Contributors

Mrinal Kanti Adak Department of Chemistry, Sidho-Kanho-Birsha University, Purulia; Department of Chemistry, Indian Institute of Technology Delhi, New Delhi, India Lambert Alff Advanced Thin Film Technology, Institute of Materials Science, Technical University of Darmstadt, Darmstadt, Germany Tahta Amrillah Department of Nanotechnology, Faculty of Advanced Technology and Multidiscipline, Universitas Airlangga, Surabaya, Indonesia Daniel Dornellas Athayde School of Engineering and Architecture, FUMEC University; Department of Chemical Engineering, Federal University of Minas Gerais, Belo Horizonte, Brazil Nurul Akidah Baharuddin Fuel Cell Institute, Universiti Kebangsaan Malaysia, Bangi, Selangor, Malaysia V. Bipin Department of Metallurgical and Materials Engineering, Indian Institute of Technology Madras, Chennai, India Rosalba Castan˜eda-Guzma´n Institute for Applied Sciences and Technology, National Autonomous University of Mexico, Ciudad Universitaria, Mexico City, Mexico ´ lvaro Cha´vez-Carvayar Materials Research Institute, National Autonomous Jose A University of Mexico, Ciudad Universitaria, Mexico City, Mexico Jose Luis Clabel Huama´n Department of Physics and Material Science, Sa˜o Carlos Institute of Physics, University of Sa˜o Paulo, Sa˜o Carlos, Sa˜o Paulo, Brazil Elaine Andrade de Mattos Department of Fundamental Chemistry, Institute of Chemistry, University of Sa˜o Paulo, Sao Paulo, Brazil Debasis Dhak Department of Chemistry, Sidho-Kanho-Birsha University, Purulia, India Jeferson Almeida Dias Institute of Science, Technology, and Innovation, Federal University of Lavras, Sa˜o Sebastia˜o do Paraı´so, MG, Brazil

xiv

Contributors

Douglas Lourenc¸o Fritzen Department of Fundamental Chemistry, Institute of Chemistry, University of Sa˜o Paulo, Sao Paulo, Brazil Suresh Kumar Garlapati Department of Materials Science and Metallurgical Engineering, Indian Institute of Technology Hyderabad, Hyderabad, India Giovanna Ferreira Bigotto Gonc¸alves Sa˜o Carlos Institute of Chemistry (IQSC), University of Sa˜o Paulo (USP), Sa˜o Carlos, SP, Brazil Marius Adrian Husanu Surface Science, National Institute of Materials Physics, Magurele, Romania Danilo Manzani Sa˜o Carlos Institute of Chemistry (IQSC), University of Sa˜o Paulo (USP), Sa˜o Carlos, SP, Brazil Euclydes Marega, Jr. Department of Physics and Material Science, Sa˜o Carlos Institute of Physics, University of Sa˜o Paulo, Sa˜o Carlos, Sa˜o Paulo, Brazil Nur Syafkeena Mohd Affandi Physics Department, Faculty of Applied Sciences, Universiti Teknologi MARA, Arau, Perlis, Malaysia Julius Motuzas FIM2 Lab—Functional Interfacial, Materials and Membrane Laboratory, School of Chemical Engineering, The University of Queensland, Brisbane, QLD, Australia Elias Paiva Ferreira Neto Department of Chemistry, Federal University of Santa Catarina (UFSC), Floriano´polis, SC, Brazil Takeo Oku Department of Materials Science, The University of Shiga Prefecture, Hikone, Japan Luis Ortega-San-Martin Departamento de Ciencias, Seccio´n Quı´mica, Pontificia Universidad Cato´lica del Peru´ (PUCP), Lima, Peru Nafisah Osman Physics Department, Faculty of Applied Sciences, Universiti Teknologi MARA, Arau, Perlis, Malaysia Jasnamol P. Palakkal Advanced Thin Film Technology, Institute of Materials Science, Technical University of Darmstadt, Darmstadt, Germany Lorena Pardo Materials Science Institute of Madrid, CSIC, Madrid, Spain Luan Passini Sa˜o Carlos Institute of Chemistry (IQSC), University of Sa˜o Paulo (USP), Sa˜o Carlos, SP, Brazil

Contributors

xv

Alexandre H. Pinto Department of Chemistry and Biochemistry, Manhattan College, Riverdale, NY, United States Dana Georgeta Popescu Surface Science, National Institute of Materials Physics, Magurele, Romania Hamimah Abd Rahman Faculty of Mechanical and Manufacturing Engineering, Universiti Tun Hussein Onn Malaysia, Batu Pahat, Johor, Malaysia Armando Reyes-Montero Materials Research Institute, National Autonomous University of Mexico, Ciudad Universitaria, Mexico City, Mexico Victor Anthony Garcia Rivera Facultad de Ciencias Fı´sicas, Universidad Nacional Mayor de San Marcos - UNMSM, Lima, Peru Lucas Carvalho Veloso Rodrigues Department of Fundamental Chemistry, Institute of Chemistry, University of Sa˜o Paulo, Sao Paulo, Brazil Suhaida Dila Safian Physics Department, Faculty of Applied Sciences, Universiti Teknologi MARA, Arau, Perlis, Malaysia Abdullah Abdul Samat Faculty of Mechanical Engineering Technology, Universiti Malaysia Perlis (UniMAP), Arau, Perlis, Malaysia J.C. Sczancoski CMDF-UFSCar, Federal University of Sa˜o Carlos, Sa˜o Carlos, SP, Brazil Firman Mangasa Simanjuntak School of Electronics and Computer Science, University of Southampton, Southampton, United Kingdom Joni Welman Simatupang Electrical Engineering Department, President University, West-Java, Indonesia A. Syed Jalaluddeen Department of Materials Science and Metallurgical Engineering, Indian Institute of Technology Hyderabad, Hyderabad, India Alexander Tkach Department of Materials and Ceramic Engineering, University of Aveiro, CICECO—Aveiro Institute of Materials, Aveiro, Portugal David James Tyler Manchester Fashion Institute, Manchester Metropolitan University, Manchester, United Kingdom Sajjad Ullah Institute of Chemical Sciences, University of Peshawar, Peshawar, Pakistan

xvi

Contributors

Wander Vasconcelos Department of Metallurgical and Materials Engineering, Federal University of Minas Gerais, Belo Horizonte, Brazil Victor Vendruscolo Department of Fundamental Chemistry, Institute of Chemistry, University of Sa˜o Paulo, Sao Paulo, Brazil Marı´a Elena Villafuerte-Castrejo´n Materials Research Institute, National Autonomous University of Mexico, Ciudad Universitaria, Mexico City, Mexico

Preface

Perovskite ceramics have emerged as promising materials from both fundamental and technological viewpoints, for instance, sensors, photovoltaic solar cells, light sources, and memories, among others. Their functionalities and intriguing physical/chemical properties are an ideal playground for fascinating areas of applications. Nonetheless, there are still many fundamental issues that remain to be understood and investigated. Hopefully, the success in the understanding of both the fundamentals for advanced research and applications on oxide and halide perovskites could result in relevant fundamental momentum and significant practical contributions to the field of advanced perovskite research. This book organized into two parts: the first on the fundamentals for advanced research and the second an extensive review of up-to-date functional applications on oxide and halide perovskites. Each part incorporates some critical topics. In this order of ideas, Part I (Chapters 1–6) offers a comprehensive overview of the development of perovskite ceramics, synthesis methods of ceramics perovskites, polaronic effects in perovskite oxides, compensation, and screening of ferroelectricity in perovskite oxides and crystal structures of copper oxide-based perovskite compounds. A comprehensive discussion of the diverse applications of perovskite ceramics is included in Part II (Chapters 7–16). The topics include multiferroic perovskite ceramics, perovskite lead-free dielectric ceramics, design for electrical applications of perovskite ceramics, functional double perovskites, perovskite-based memories, light storage perovskites, perovskite-structured ceramics in solid oxide fuel cell applications, perovskite membranes for oxygen separation, perovskite-based light-emitting diodes, and applications of halide-based perovskites in photonics. Therefore, this book provides a comprehensive study of the perovskite ceramic from our authors’ advanced-level research and their advanced applications. In this context, we hope that this book contributes to undergraduate and postgraduate students, researchers, and industrial technology. Finally, we gratefully acknowledge the contributions of many experts and colleagues around the world in this field. Thank you for accepting this challenge, and for your great effort and dedication to this project. Jose Luis Clabel Huama´n Victor Anthony Garcia Rivera

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Part One Advanced research

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Introduction to perovskites Luis Ortega-San-Martin Departamento de Ciencias, Seccio´n Quı´mica, Pontificia Universidad Cato´lica del Peru´ (PUCP), Lima, Peru

1

More than 180 years have passed since Gustav Rose received from Russia an unknown mineral sample whose name now gives this and many other books their title (Rose, 1839). With the formula CaTiO3, and known as perovskite, this mineral has a crystalline arrangement that describes one of the most versatile crystalline structures that exist: the perovskite structure. Until the 1920s, however, this mineral and other inorganic compounds with similar formulas and structures were not more than mineralogical and chemical curiosities. At the end of that decade, VM Goldschmidt’s group found that the perovskite structure could be adopted by numerous substances of the general formula AMX3 (where A and M were different metal cations and X oxide or fluoride ions) as long as specific size requirements between the ions were satisfied. That discovery stimulated research into perovskites, some of which were even used as pigments for exterior paints in the 1940s. However, the technological boom based on perovskites, which also came in the 1940s, was based on their ferroelectric properties and their immediate use as capacitors, first demonstrated using BaTiO3 (Cross & Newnham, 1987). Since Rose’s first publication, tens of thousands of works (papers, books, theses, etc.) related to materials with a perovskite structure have been published (Fig. 1.1). A similar number of patents involving perovskite-related materials exists, which brings the number of available documents to more than a quarter of a million. Hundreds, if not thousands, of researchers have somehow been involved with these types of substances. Even numerous Nobel Prize winners have been involved in studies with compounds from this vast family. Some of them worked with perovskites during their early stages of research, others for long periods, and only a few received the award for perovskite-related works. Table 1.1 provides this information. This table is probably not complete but gives us an idea of how relevant this field of research has been in the last century. Is it possible to summarize all the available information on perovskites in a short introduction? Furthermore, how can we summarize all this without repeating what other generic works (Tilley, 2016) have already done? In recent years, this has become even more complicated as the available information has grown almost exponentially. According to Fig. 1.1 from the Scopus database, in 2021 alone, more than 10,000 works related to perovskites were published primarily, but not only, on the photovoltaic properties of perovskites for their application in solar cells. Staying up to date in this field is already a utopia. However, perovskites are more than photovoltaic promises. There are huge expectations in the field of photodetectors, light-emitting diodes Perovskite Ceramics. https://doi.org/10.1016/B978-0-323-90586-2.00015-2 Copyright © 2023 Elsevier Ltd. All rights reserved.

4

Perovskite Ceramics

Fig. 1.1 Academic documents involving perovskite-structured materials published from 1950 to 2021. The data shown in the figure were extracted from the Scopus© database and include regular papers, conference papers, reviews, and books, but not editorials, erratum, notes, or retracted documents. The period before 1981 is inset due to the tremendous scale difference with the current publishing pace. The document search included publications that mentioned the word perovskite in the title, abstract, or keywords and also all documents that did not contain the word “perovskite” but were devoted to well-known perovskites (barium titanate, lead titanate, Ruddlesden-Popper phases, etc.).

Table 1.1 Nobel Prize laureates who worked on perovskite-related issues. Nobel Laureate

Date and field

Linus Pauling

1954, Chemistry

Willard Libby

1960, Chemistry

Giulio Natta

1963, Chemistry

Time period and research involved on perovskites

Perovskite related?

1920s: Worked on complex hexahalides of the type A2AMX6, now considered as double perovskites 1970s: Worked on lanthanum cobaltite perovskites as possible exhaust catalysts 1920s: Worked on the crystal structure of simple perovskite halides

No

No

No

Introduction to perovskites

5

Table 1.1 Continued Nobel Laureate

Date and field

Philip Warren Anderson

1977, Physics

Nevill Francis Mott

1977, Physics

Johannes Georg Bednorz

1987, Physics

K. Alexander M€uller

1987, Physics

Pierre-Gilles de Gennes

1991, Physics

John B. Goodenough

2019, Chemistry

M. Stanley Whittingham

2019, Chemistry

Time period and research involved on perovskites

Perovskite related?

1980s–2010s: Worked on the explanation of the superconducting mechanism of layered perovskite oxides 1970s–1990s: Worked on the electronic and magnetic properties of several perovskites, including layered superconductors 1980s–2010s: Worked on (and also discovered) the superconducting properties of several layered perovskite oxides 1980s–2010s: Worked on (and also discovered) the superconducting properties of several layered perovskite oxides 1960s: Worked on the magnetic properties and double exchange mechanisms on manganite perovskites 1960s–2010s: He has contributed to a number of areas of perovskite research, from magnetic properties to electronic properties 1970s–1990s: Has contributed to the understanding of mixed conductors, especially those related with tungsten bronzes (some of which have the perovskite structure with vacancies in the A-site)

No

No

Yes

Yes

No

No

No

The table shows the laureates, the date and field of their Nobel Prize, the period in which they did research involved in perovskites, and if the prize was related to their research on perovskites.

6

Perovskite Ceramics

and also for field-effect transistors. Even today, properties with potential applied interest are still being discovered in compounds with this structure. Is there a way to be aware of all the research trends? Within this vast ocean of information (which now includes the present book on ceramic perovskites), the objective of this short introduction is modest. I will briefly describe what we consider a material with a Perovskite or Perovskite-type structure; then I will focus on some areas of interest outside of the leading research trends, that is, away from solar cells or ferroelectric substances. Occasionally, I will mention some historical precedents related to the properties or compounds covered in this chapter. As far as possible, these properties will be associated with the most outstanding characteristics of this “simple” and, at the same time, complex structure.

1

What is a perovskite?

Despite the many years of perovskite research, the concept of what a perovskite is is still sometimes complex to establish. A few years ago, I tried (Ortega-San-Martin, 2020) to summarize the history of perovskite research, emphasizing the evolution of the generic concept of a “Perovskite-type structure.” After so many years, the only clear thing is that the idea of a “perovskite-type” or “perovskite family” is extremely broad. In some extreme cases, it becomes difficult to see the relationship between the structure of a family member and that of the ideal perovskite structure. Fig. 1.2 shows, for example, four members of the family drawn using Vesta software (Momma & Izumi, 2011) from data deposited in the Crystallographic Open Database (Grazˇulis et al., 2012). At first glance, these are four different crystal structures, but the key to classifying them as perovskites is how we interpret the basic building units of each structure. Fig. 1.2A is the prototypical cubic perovskite showing the three-dimensional arrangement of MX6 octahedra forming cubo-octahedral voids filled with element A. Fig. 1.2B shows the Bi5FeTi3O15 structure, which is the n ¼ 4 member of the Aurivillius layered group of the general formula (Bi2O2) (An1BnO3n+1). As Aurivillius originally suggested (Aurivillius, 1949), the easiest way to understand this layered structure is considering it as a stacking of “Bi (Fe0.25Ti0.75)O3” perovskite layers between bismuth oxide (Bi2O2) layers. Fig. 1.2C is the structure of CsLaNb2O7, a member with n ¼ 2 of the family of layered structures known as Dion-Jacobson. Finally, the last structure (Fig. 1.2D) is the post-perovskite that results from the transformation of the perovskite MgSiO3 at the core-mantle boundary due to extreme pressure and temperature conditions. If we consider as family members all the variants with A, B, or X-site defects or excesses, the number of possibilities multiplies. Additionally, if we add the hexagonal polytypes, in which the octahedra share faces, the variety of substances that adopt the Perovskite structure becomes almost unmanageable and even challenging to classify (although there are reasonable attempts in the literature) (Mitchell, 2002; Tilley, 2016). Yet this is without considering the case of the now ubiquitous hybrid inorganic-organic perovskites, which will be mentioned later on. No wonder that,

Introduction to perovskites

7

Fig. 1.2 Different examples of perovskite structures. (A) Ball and stick model of the ideal cubic perovskite structure (top) and the polyhedral view of the same structure (bottom); (B) n ¼ 4 member of the Aurivillius family of the general formula (Bi2O2)(Bin1BnO3n+1); (C) n ¼ 2 member of the Dion-Jacobson laminar structure of the general formula An1BnO3n+1; (D) postperovskite structure, a high-pressure and -temperature variety. Data for the structures are from open data deposited in the Crystallographic Open Database (COD) (data numbers are (B) 1521427, (C) 2004717, (D) 9009217).

in recent years, voices have emerged asking to limit the term "perovskite" to something more manageable and for structures closer to that of the original mineral (Akkerman & Manna, 2020). What types of perovskites will be considered in this chapter? We will focus mainly on the “classic ones,” those with the generic formula AMX3 and a three-dimensional arrangement of vertex-sharing BO6 octahedra (Fig. 1.2A). These will include all possible distortions due to the change of A, M, and X, regardless of the nature of these elements or ions. In fact, we will include compounds with either A, M, and X being traditional monatomic ions (metal ions, halides, oxide, etc.) or polyatomic ions, molecules, and/or any other moiety that fits into the structure (the hybrid inorganic-organic perovskites).

8

2

Perovskite Ceramics

Symmetry matters: From the bumped journey that started with the pursuit of the appropriate symmetry of the mineral perovskite to the search for modern lasers

In Ortega-San-Martin (2020), I focused on describing the problem that initially existed to correctly determine the appropriate symmetry of the mineral perovskite. It took more than 100 years to find the correct information. Before the discovery of X-rays, research focused on elucidating the crystalline class. This slow process generated conflicting literature and interesting debates among crystallographers, mineralogists, and chemists. The debate fluctuated between the cubic, rhombic, and monoclinic systems and was closed in 1912 when two studies rightly concluded that the crystals of the mineral perovskite were rhombic (B€oggild, 1912). After the discovery of X-ray diffraction, the original debate was reopened because the first description of the atomic arrangement of the mineral Perovskite assigned it a cubic centrosymmetric space group. It became difficult to explain the inconsistency, and it got more complicated after the work of Victor M. Goldschmidt’s research group in the 1920s (Goldschmidt, Barth, Lunde, & Zachariasen, 1926). This group synthesized numerous oxides of the formula AMO3, all with the same basic arrangement as the perovskite mineral. As expected, some of the newly obtained compounds were not cubic, but the group did not pay much attention to their symmetry. They focused on something more interesting: they were the first to note the incredible versatility of this structure toward ion substitution as long as they met certain size relations. Those relations were parametrized by the well-known tolerance factor. ðr + r X Þ t ¼ pffiffiffiffi A 2 ðr B + r X Þ

(1.1)

where rA, rB, and rX are the tabulated ionic radius of the A-, B-, and X-site ions. Usually, t values are between 0.8 and 1.1 for perovskite-structured phases. Thanks to these studies, we now know that if we carefully change the A, M, and X ions, the basic building block of the perovskite structure (the vertex shared octahedra shown in Fig. 1.2A) is retained. Still, the symmetry of the compound usually changes. That was not easy to explain with the diffraction equipment of the early 20th century, so the end of the debate on perovskite symmetry had to wait until the 1950s. The verdict: CaTiO3 shows a centrosymmetric rhombic distortion (Kay & Bailey, 1957). All this debate was not in vain. The studies carried out in the first half of the last century revealed the two most essential characteristics of the perovskite structure. First, the incredible versatility of the perovskite structure to accommodate a considerable number of different ions (even at the same time) in its different crystallographic positions. Today, it is common to find perovskites with up to four different ions in the same position [see, for example (La0.80Nd0.05Sr0.10Ca0.05)(Cr0.90Ni0.10)O3 (OrtegaSan-Martı´n et al., 2018)]. Second, it was possible to see that all this compositional

Introduction to perovskites

9

variety and change of ions also caused structural changes not only of crystallographic interest but also of critical practical importance. We know that when ions of different sizes are introduced, they reorganize according to their size: either through cooperative displacements of the X ions (which results in BX6 octahedra tilting) or by small displacements of the A and M ions from their ideal positions (Glazer, 1972). These changes are dependent on the ion concentration as observed for the PbTi1-xZrxO3 system (Fig. 1.3), a perovskite solid solution of outstanding practical importance. These small structural changes (which are also dependent on temperature and pressure) have been systematized in recent years (Howard & Stokes, 2005; Stokes, Kisi, Hatch, & Howard, 2002). There is even a specific software [SPuDs, updated in 2021 (Lufaso, 2021)] that predicts the distortion of a compound once its composition is known. Understanding these distortions has also been key to understanding the fundamental properties of perovskites, such as ferroelectricity, the main exploited property of perovskites in practical applications. The most popular ferroelectric compounds are titanates with the formula PbTi1-xZrxO3 (PZT) and Ba1-XSrxTiO3 (BTO). The importance of these substances lies mainly in the fact that their structure is not centrosymmetric in a specific temperature

Fig. 1.3 Structural transitions in perovskites. (A) Compositional dependence of the structural transitions observed in the PbZr1xTixO3 perovskite ceramics. (B) High-temperature centrosymmetric ideal structure for 0  x  1. (C) Tetragonal noncentrosymmetric modification observed when x ¼ 0.6. The practical utility of these perovskites depends on the symmetry adopted by the solid solution. A: From Li, M. J., Xu, L. P., Shi, K., Zhang, J. Z., Chen, X. F., Hu, Z. G., Dong, X. L., & Chu, J. H. (2016). Interband electronic transitions and phase diagram of PbZr1  xTixO3 (0.05  x  0.70) ceramics: Ellipsometric experiment and first-principles theory. Journal of Physics D: Applied Physics, 49(27), 275305. https://doi.org/10.1088/0022-3727/49/27/275305. Licensed under CC by 3.0 license, http://creativecommons.org/licenses/by/3.0/; B and C drawn using Vesta.

10

Perovskite Ceramics

range. This is why the original debate on the exact symmetry of the mineral Perovskite makes sense: symmetry matters. Centrosymmetric perovskites cannot show the same properties as their noncentrosymmetric counterparts. For example, the off-center displacement of the Pb and Zr/Ti ions observed in Fig. 2.1C with respect to Fig. 2.1B is the key to their well-known dielectric properties. This displacement results in a permanent polarization of these ceramics that can be modified under an applied load (piezoelectricity), under an applied electric field (inverse piezoelectric effect), or even by the change in temperature (pyroelectric effect). These properties can only be observed in noncentrosymmetric crystals such as BTO, PZT, and other similar perovskites. These perovskites and their properties have been vital in the current technological advance since they have allowed the manufacture of capacitors, actuators, transducers, etc. These properties and their basic explanation in connection with perovskites have been extensively dealt with in the literature and even in basic materials science and solid-state chemistry textbooks. Nevertheless, piezoelectricity and pyroelectricity are not the only properties that require a noncentrosymmetric crystal structure. Some optic effects such as the linear electro-optic effect (Pockels effect) and nonlinear optical properties, which find several applications in our everyday life, also depend on the absence of a center of symmetry. Perovskites dominate this field, and the most important one with electro-optic effects is by far lithium niobate, LiNbO3 (Weis & Gaylord, 1985). This highly distorted perovskite has been widely used since the 1970s in applications such as acoustic wave transducers, acoustic delay lines, acoustic filters, optical amplitude modulators, optical phase modulators, second-harmonic generators, Q-switches, beam deflectors, phase conjugators, dielectric waveguides, etc. Simple devices such as the Green lasers used in classrooms, for recreational purposes, or even for ophthalmologic applications rely on one of such noncentrosymmetric properties: nonlinear optical properties (NLO). In general, NLO materials can act by adding the frequency of several incident photons to generate a photon of higher energy. The most common are secondary harmonic generators (SHG), which double the frequency of incident radiation. The most important consequence of this property is that NLO materials can expand the wavelengths of existing laser sources to much wider spectral regions, increasing their applicability in many technological devices. For UV-lasers, for example, those applications include semiconductor fabrication, laser micromachining, photolithography, and attosecond pulse generation. Depending on the required wavelength, they can even be used for ultrahigh-resolution photoemission spectrometry and photoelectron emission microscopy (Tran, Yu, Rondinelli, Poeppelmeier, & Halasyamani, 2016). Unfortunately, the NLO crystals used for laser applications are not yet perovskites. In recent years, however, more and more compounds with this structure show interesting NLO properties, most of them being conventional halide perovskites (Dong et al., 2019; Xu, Li, et al., 2020). Still, they have also been observed in the so-called antiperovskites such as K3(B6O10)X, with X ¼ Cl or Br (Al-Ama, Belokoneva, Stefanovich, Dimitrova, & Mochenova, 2006; Wu et al., 2011). Antiperovskites? Yes. Antiperovskites are simply those perovskites in which the positions of cations and anions are reversed (in AMX3 antiperovskites, X represents metal cations

Introduction to perovskites

11

while A and M are anions). The K3(B6O10)X ceramics are an example and have a three-dimensional arrangement of corner-shared XK6 octahedra with a tridimensional borate moiety filling the position of traditionally large A cations. Fortunately, the abovementioned perovskites crystallize in noncentrosymmetric space groups and show promising SHG NLO properties. “Promising” might be considered an overused word in the perovskite field, but past discoveries have shown that many promises are fulfilled. And how do promises come true in the perovskite world? The answer is simple: by changing some ions for others. We have already said that one of the most outstanding characteristics of this type of structure is that it can accommodate almost any combination of ions. That is the first lesson for those who work with perovskites: if you cannot find the right property in your composition, change it until you find what you are looking for. The basic structure will prevail, but you may modify its distortion and symmetry, which could result in new or even unexpected properties. To date, these kinds of attempts have not been very successful with antiperovskites. The reason? Because when potassium is replaced by other metals such as sodium, rubidium, or their mixtures, the result is a centrosymmetric perovskite (Mutailipu, Poeppelmeier, & Pan, 2021). This again highlights the importance of knowing the symmetry of the synthesized compounds. Nevertheless, this should not be considered a defeat for antiperovskites, as many have shown properties of great interest. Among them are very fast ion conductivities of interest for Li-ion Batteries, giant magnetoresistance, and even promising superconducting properties (Wang et al., 2019). If the strategy of changing ions from one site of the perovskite does not work, we change those from another site, and, if necessary, we introduce complete molecules in certain positions. This is not the Olympics, so any type of doping is allowed. In fact, the introduction of complex molecular cations such as N-methyl-N0 -diazabicyclo [2.2.2] octonium in metal-free perovskites such as A(NH4)X3 (where X is Cl, Br, or I) has shown quite promising results (Kasel et al., 2019). In addition to this, the latter type of compounds also has ferroelectric properties comparable to conventional ceramics such as barium titanate (Li & Ji, 2018). Will they replace BTO in its numerous applications? Expectations are high, but practical applications are still awaited. As can be seen, the search for materials with properties of interest is linked to the design of new compounds that not only have a particular structure but also crystallize in noncentrosymmetric space groups. Predicting the final space group of a perovskitestructured compound is not an easy task. However, the path is reasonably known: certain combinations of A and M elements may lead to structural features (MX6 octahedra rotations and/or A site ordering) that prevent the presence of a symmetry center (Young, Lalkiya, & Rondinelli, 2016). The 19th-century debate over the perovskite’s mineral symmetry was not consequently useless debate. In any case, if this symmetry-related path to perovskite-based lasers is not satisfactory, there is a considerable research effort based on exploiting another property of some perovskites: the semiconducting properties of halide perovskites (Liao, Jin, & Fu, 2019). Halide perovskites are bandgap semiconductors with high photoluminescence quantum yield, narrow emission bandwidths, and tunable color emissions, making them very promising light-emitting materials and also excellent gain

12

Perovskite Ceramics

materials for laser applications. As a result of their large degree of bandgap tunability, which is usually a result of the superb flexibility of the perovskite structure to ion substitution, lasing wavelengths of perovskite materials range from near-infrared to ultraviolet (Lei, Dong, Gundogdu, & So, 2021; Liao et al., 2019). Considering that these perovskites are reasonably easy to synthesize, enormous expectations have been placed in this area.

3

An incredible structural versatility that can lead us to the next generation of solid-state refrigerators: Hybrid perovskites

As already indicated, a common strategy to find new properties, or improve existing ones, involves a well-planned ion change on the perovskite structure. Using the tolerance factor mentioned above (Eq. 1.1), this planning becomes an “easy” task if we work with well-known spherical ions. But… How do we proceed when big nonspherical molecular ions (such as N-methyl-N0 -diazabicyclo [2.2.2] octonium mentioned before) are introduced into the structure? We will return to this soon, but first, another question: when did we realize that this structure could accommodate molecular ions in its crystallographic positions? almost a century ago … but it was not known until later. In fact, the first researchers to structurally characterize perovskites containing polyatomic ions did so without knowing that they were describing the first examples of this structure. These early chemists include the young Linus Pauling, a future Nobel prize winner in Chemistry. He did so during his PhD while elucidating crystalline structures of complex salts discovered in the 19th century (Pauling, 1924). Among the compounds studied in those times are those of the formula (NH4)3MX6, with X ¼ Fe, Al, Si, Pt, and X ¼ F, Cl. These halides were later classified as elpasolites or cryolites, which in reality are full members of the perovskite family. Their apparent complex formula can be simplified as A(A0.5B0.5)X3, where the ammonia ion occupies all the A positions and half of the B positions. As shown in Fig. 1.4A, the main difference from a conventional perovskite is that the AX6 and BX6 octahedra [e.g., “(NH4) F6”] are arranged alternately. These compounds are commonly formulated as A2B0 BX6 (sometimes B0 can also be occupied by A ions) and are generically known as double perovskites. The crystalline structure of many such compounds was elucidated at the beginning of the last century, but they always remained a crystallographic curiosity (Babel, 1967). The disorder that exists in the halide positions of these perovskites probably hindered early characterizations, thus relegating these compounds to the exotic part of the structural family of perovskites (Udovenko, Laptash, & Maslennikova, 2003). The systematic structural characterization of perovskites with more complex polyatomic ions began in the late 1960s. From this moment on, the possibility of inserting polyatomic ions in both A-sites and the anion sites was realized. Compounds such as A2BB0 (CN)6 (with the cyanide ion in the X site) were characterized at this time

Fig. 1.4 The crystal structures of two hybrid perovskites. (A) is (NH4)3FeF6, probably the first crystal structure of a perovskite containing polyatomic ions and (B) is [CH3NH3][Mn(N3)3], an example of a metal organic framework-like hybrid inorganic-organic perovskite. Drawn using VESTA from COD datafile numbers 1010085 and 4121027. (References for VESTA and COD are given in the main text.)

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Perovskite Ceramics

(Swanson & Ryan, 1973). It was also discovered that old compounds with big polyatomic ions in the A-site also had the perovskite structure. This was the case of the methylammonium perovskites such as (CH3NH3)BX3 with B ¼ Sn and/or Pb (Weber, 1978a, 1978b) and related compounds such as [N(CH3)4]2CsFe (CN)6 (Babel, 1982). These unconventional compounds proved the incredible versatility of the perovskite structure, but their properties were never the subject of significant attention until recently. The phenomenal photovoltaic (PV) properties of such halide perovskites resulted in an unexpected research boom due to their promise to substitute siliconbased PVs. Perovskites consisting of organic and inorganic ions are considered a new family known as the hybrid inorganic-organic perovskites (HIOPs). The possibility of using dozens of different polyatomic ions in the A-, B-, and Xsites (Fig. 1.5) has resulted in the synthesis of hundreds of HIOPs. Consequently, it has been necessary to rewrite the tolerance factor so that it can accommodate the different geometries that we can find in molecular ions (Eq. 1.2): ðr + r Xeff Þ   t ¼ pffiffiffiffiAeff 2 rB + hXeff 2

(1.2)

In this new form, rB is the tabulated ionic radius of the B-site metal ion, rAeff is the effective radius of the A-site cation, and rXeff and hXeff are the effective radius and length of the X-site molecular ions. This new factor depends on the geometry of the polyatomic ions, but calculations to date have shown that most hybrids with the perovskite structure fall within tolerance values between 0.8 and 1.0 (Kieslich, Sun, & Cheetham, 2014, 2015). The chemical and structural variability that these organic and inorganic fragments allow helps to expand the perovskite family and the concept itself of “the perovskite family” to almost unsuspected limits. The new possibilities are enormous, although on some occasions this relationship is more fictitious than real (as in the case of “0D” hybrid perovskites) (Saparov & Mitzi, 2016). As expected, the greater the number of possible compounds, the more new properties appear, which include (Li et al., 2017): excellent photovoltaic properties as a result of being direct bandgap semiconductors with bandgaps ranging from 1.2 to 2.8 eV, promising light-emitting capabilities, competitive ferroelectric properties (including relaxor ferroelectric properties), ferroelastic properties, and multiferroic properties. Many of these properties appear (or intensify) due to the structural transitions that almost all hybrid perovskites present, multiplying their potential applications. Unlike the structural transitions of traditional perovskites (oxides or halides such as BaTiO3), the phase transitions in these types of compounds are more complex. These transitions not only involve the usual displacements of the A- and B-sites or the tilting of the BX6 octahedra. They also involve order-disorder transformations of the A-site organic/inorganic moieties (Xu, Du, Zhang, & Chen, 2016). An example of this type of transformation (order-disorder coupled with octahedra tilting) is shown in Fig. 1.6 for (3-ammoniopyrrolidinium) RbBr3 (drawn with VESTA using crystallographic

Introduction to perovskites

15

Fig. 1.5 A selection of different A-site and X-site ions present in hybrid organic-inorganic perovskites. Color schemes: N, blue; O, red; C, black; H, gray or light pink; Cl, purple or green; Br/I, purple; S, light yellow; P, yellow or purple; M ¼ Ag or Au, yellow; B, turquoise; and F, pink. From Li, W., Stroppa, A., Wang, Z.-M., & Gao, S. (2020). Hybrid organic-inorganic perovskites (1st ed.). Wiley-VCH Verlag GmbH & Co. KGaA. https://doi.org/10.1002/9783527344338.

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Perovskite Ceramics

Fig. 1.6 Order-disorder transition in (3-ammoniopyrrolidinium)RbBr3 hybrid inorganicorganic perovskite. Order-disorder transitions such as the one pictured in this figure result from the disorder created in the organic ion site (A-site) above a specific temperature. At low temperatures (left), hydrogen bonding with the surrounding perovskite ions fixes amine atomic positions. At higher temperatures (right), amine atoms are disordered, so we see a blurred amine in the center. Such ferroelectric-paraelectric, order-disorder transformations might result in practical applications in photovoltaics and other electronic devices. The high entropy change has also been identified as a useful mechanism for barocaloric effects.

data from Pan et al., 2017). These complicated transitions, many of which involve polar amines such as the one pictured in Fig. 1.6, may cause ferroelectric ordering. This in turn increases their potential uses either in photovoltaic applications or in other polar-based electronic devices (Xu, Kopyl, Kholkin, & Rocha, 2019). Some of these transformations depend on the hydrogen bonds established at low temperatures between the organic ion and the anions in the structure, which are broken when the temperature is raised. These order-disorder transformations result in large entropy changes not observed in traditional perovskites. Moreover, this order-disorder behavior is easily affected by external fields, resulting in large caloric effects (Li et al., 2017). This opens a new window of application for perovskite materials that could be very promising in the framework of future global warming. Let us elaborate on this. Caloric effects can be divided into different groups depending on the stimulus that causes the change in temperature of the material. Barocaloric and elastocaloric materials are those materials that change their temperature under pressure (depending on the stress direction, we have either one or the other). On the other hand, electrocaloric materials are those that respond to an electric field, and finally, magnetocaloric materials are affected by magnetic field changes (Moya, Kar-Narayan, & Mathur, 2014). This ability of changing their temperature under different stimuli can be used in the refrigeration area, especially for solid-state refrigeration (Szafranski, Wei, Wang, Li, & Katrusiak, 2018). Let us remember how refrigeration works. The essence of conventional refrigeration lies in the ability of gases to be expanded and compressed under changing pressure conditions. The continuous release and application of pressure on these gases allow temperature changes in conventional

Introduction to perovskites

17

refrigerators roughly from  18 to 8°C (ΔT ¼ 20 K). Hydrofluorocarbons are still the most commonly used gases in general appliances (freezers or air-conditioning systems), but they either have important global warming effects or are hazardous chemicals. Consequently, there is an important research effort to find suitable substitutes that can change their temperature under an external stimulus. This is where caloric effects on perovskites may find promising applications (Barman, Kar-Narayan, & Mukherjee, 2019) (I know, too many uses of the “promising” word). It is important to note that caloric effects are not new to perovskites. They have been known for years in traditional ceramics such as the doped phases of BaTiO3 and PbTiO3, and even Pr1-xLaxNiO3, but the induced temperature changes have been minimal (ΔT 14 K), on the other hand, have also been found in some oxofluorides, such as those studied by Pauling 100 years ago, albeit at very high pressures (close to or greater than 1 GPa) (Flerov, Gorev, Molokeev, & Laptash, 2016). Conventional appliances cannot generate such high pressures, so the future relies on materials that respond to easily reached pressures. This seems to have been achieved in hybrid perovskites such as [(CH3)4N]Mn [N3]3, [TPrA]M[dca]3 (M ¼ Mn, Cd, TPra is the tetrapropylammonium cation and dca is the dicyanamide anion) (Bermu´dez-Garcı´a et al., 2017; Salgado-Beceiro et al., 2020). The order-disorder effects on these perovskites occur near room temperature and, in some cases, can generate a ΔT  5 K with pressures as low as 70 bar. This finding has opened a new window of application for the family of hybrid perovskites. In fact, similar phenomena have been observed in numerous perovskites, so a promising future is expected. The variables to optimize are known: a large and reversible entropy change from a phase transition, a significant pressure dependence of the transition temperature (δTt/δP, known as the barocaloric coefficient), and a transition temperature close to that of the desired application (Bermu´dez-Garcı´a, Sa´nchezAndu´jar, & Sen˜arı´s-Rodrı´guez, 2017). The next step is finding the perovskites that satisfy them, although there are many candidates, as shown in Fig. 1.7. Would these new perovskites satisfy all the stringent characteristics needed for solid-state refrigeration materials (F€ahler et al., 2012) to become the real future of refrigeration?

4

Chemical flexibility: Toward the next generation of electrocatalyzers proposed by modern machine learning tools

We have already seen that the chemical flexibility of the perovskite structure for the accommodation of cations and anions of different sizes and forms is remarkable. The inorganic-organic hybrid perovskites and classical ferroelectric oxides such as the PbZr1xTixO3 are clear examples: in all these cases, the ion substitutions produce structural changes that in turn impact their final properties and potential applications.

18

Perovskite Ceramics

Fig. 1.7 Barocaloric data for several hybrid inorganic-organic perovskites. The image shows the barocaloric coefficient (δTt/δP) dependence on (A) their solid-solid phase transition entropy change (ΔStransition) and (B) their transition temperature (Tt). In these perovskites, the entropy change of the phase transition (ΔStransition) results from configurational, rotational, and/or vibrational contributions and determines the maximum barocaloric effect attainable. The barocaloric coefficient (δTt/δP) arises from the material’s pressure dependence and determines the operational temperature window of the material. Adapted with permission from Bermu´dez-Garcı´a, J. M., Sa´nchez-Andu´jar, M., & Sen˜arı´sRodrı´guez, M. A. (2017). A new playground for organic–inorganic hybrids: Barocaloric materials for pressure-induced solid-state cooling. Journal of Physical Chemistry Letters, 8(18), 4419–4423. https://doi.org/10.1021/acs.jpclett.7b01845. Copyright 2017 American Chemical Society.

In the previous cases, the substitutions do not necessarily involve changes in the oxidation states of the present ions. However, when this occurs, profound property changes may result. For example, the electrical and magnetic properties of LaMnO3 dramatically change when some of the lanthanum is replaced by an alkaline earth cation, such as calcium and barium. The substitution of a + 3 cation by a + 2 one reduces the positive charge of the A site, which must be compensated in two possible

Introduction to perovskites

19

ways: either by increasing the oxidation state of manganese (a part of Mn+3 becomes Mn+4), or by forming oxygen vacancies. Depending on the preparation atmosphere, a combination of both may result. In LaMnO3, for example, the partial substitution of La by Ca or Ba implies a change from an insulating-antiferromagnetic oxide to a ferromagnetic conducting compound, which also shows colossal magnetoresistance near room temperature ( Jin et al., 1994). This dramatic change in properties is not only of interest for magnetic-related applications (in data reading and recording, for example) but has been of great interest since the excellent catalytic properties were discovered in the 1970s on doped perovskites such as La1 xSrxCo1 yRuyO3 (Voorhoeve, Johnson, Remeika, & Gallagher, 1977). To a large extent, these properties depend on the interplay between the different oxidation states of the transition metals and the oxygen vacancies. Both variables can be modified with relative freedom simply by introducing the desired quantities of the elements of interest in A and B positions. The amount of different ions that can be introduced in each perovskite site is not yet limited, allowing us to synthesize complex combinations such as the abovementioned (La0.80Nd0.05Sr0.10Ca0.05)(Cr0.90Ni0.10)O3 (Ortega-SanMartı´n et al., 2018) to obtain the desired property. Just as oxygen vacancies, A-site vacancies are also possible. The judicious combination of all these phenomena generates a myriad of different properties that can be exploited in numerous areas related to energy generation (batteries, fuel cells, photocatalysts, catalysis, thermoelectrics, and solar thermal conversion), or for the catalytic transformation of various substances (CO2 reduction to methanol, the reduction of nitrogen to ammonia, methane reforming, etc.) (Bian et al., 2020; Hwang et al., 2017). The applications mentioned above arise from adequately exploiting the ionic conduction resulting from the oxygen or A-cation site vacancies and the electrical conduction resulting from the B-site mixed valence. If we combine an adequate selection of ions with an optimized preparatory process, we can obtain new materials with properties of great interest (Irvine et al., 2021), such as excellent photocatalytic activity for the generation of hydrogen by water splitting, proton-conducting oxides for the next generation of energy conversion, potential thermoelectric phases capable of directly converting heat into electricity, electrolytic materials for fuel cells and reversible electrode materials for either hydrogen generation (power to fuel) or electricity generation from hydrogen (fuel to power), as shown in Fig. 1.8. There are many issues to overcome, but the time for promises is over: the challenges are now bringing the results obtained with laboratory prototypes to real applications. Among the most interesting challenges in recent years is the low-cost manufacture of clean fuels such as H2 through the electrolysis of water (H2O ! H2 + ½O2) without the use of scarce and valuable metals such as iridium or ruthenium. Clean and easy production of H2 is also beneficial for a large number of chemical industries, including ammonia production. One of the main obstacles of this process is the so-called oxygen evolution reaction (OER). Consequently, the development of low-cost and robust OER catalysts is critical to solving this efficiency problem in water splitting (Song et al., 2020). Ba0.5Sr0.5Co0.8Fe0.2O3–δ (BSCF) (Suntivich, May, Gasteiger, Goodenough, & Shao-Horn, 2011) is a remarkable perovskite for this purpose not only for its excellent

20

Perovskite Ceramics

Fig. 1.8 Power to fuel (left) and fuel to power (right) modes proposed for reversible consumption and generation of hydrogen using proton-conducting perovskites. This form of power to fuel and fuel to power mode is one of the possibilities envisioned for solid oxide fuel cells which can also function using oxide ion conductors; in such cases, O2 are the mobile ions (usually at the expense of increasing operation temperatures). From Irvine, J., Rupp, J. L. M., Liu, G., Xu, X., Haile, S., Qian, X., et al. (2021). Roadmap on inorganic perovskites for energy applications. Journal of Physics: Energy, 3(3), 031502. https:// doi.org/10.1088/2515-7655/abff18. Licensed under a Creative Commons Attribution 4.0 license. https://creativecommons.org/licenses/by/4.0/.

electrocatalytic properties but also for how its synthesis was conceived. In this case, unlike what many of us have done in the past (prepare a lot of compounds varying the composition until the perovskite with the best properties among those synthesized is obtained), there was a previous study of which variable was the one that should be optimized before carrying any laboratory trials. A previously known idea was followed: perovskites with good catalytic properties contained transition metals with certain degrees of d orbital occupancy. So after reviewing the various compounds in the literature and potential ion combinations, BSCF was successfully proposed. For the latter example, it was considered that the OER occurring on perovskite surfaces follows the adsorbate evolution mechanism (AEM), which involves electronproton transfers on the transition metal active centers, so their electronic population had to be optimized. Another successful approach to the problem has been the recognition that oxygen vacancies also play an essential role [through the lattice oxygenmediated mechanism (LOM)] in the OER, so the search for perovskites with some oxygen vacancies is also an important path to find the best perovskite electrocatalyst (Pan et al., 2020). As expected, the future relies on identifying the best combination of A and B elements to have an adequate number of d electrons and generate sufficient oxygen vacancies. This also involves finding the appropriate synthesis processes and optimal manufacturing procedures. And how do we do that? Again by trial and error? Not quite. The current trend is to replace laboratory work with “machine” work (at least a part of it). Recent searches for potential catalysts, which include other families of perovskites such as the Ruddlesden-Popper phases (Xu, Pan, Zhong, Ran, & Shao, 2020), are carried out by combining robust computational studies based on DFT coupled (or not) with machine learning and artificial intelligence ( Jacobs, Hwang, Shao-Horn, & Morgan, 2019; Li, Achenie, & Xin, 2020; Weng et al., 2020). These computational tools, which are now used in all fields of perovskite research (Tao, Xu, Li, & Lu, 2021), may save us a lot of experimentation effort, although they give

Introduction to perovskites

21

good results only if calculations are preceded by some experimentation, as reported by Weng et al. (2020) (Fig. 1.9). This new tendency has also impacted the most critical parameter of perovskite stability: the tolerance factor. The need for an accurate variable that helps in finding new perovskites for different applications has led to the rewriting of the tolerance factor (Bartel et al., 2019) as shown in the following equation: τ¼

  ðr A =r B Þ rX  nA  rB ln ðr A =r B Þ

(1.3)

where nA is the oxidation state of A and ri is the ionic radius of ion i. According to this new proposal, a perovskite structure is found when τ is below 4.18. Will these “machine-based” analyses end the work of laboratory enthusiasts? On the contrary, it will guide chemists and materials scientists toward the most promising materials that should be pursued in the laboratory.

5

Never abandon your perovskite; others might find it useful

The almost limitless versatility of the perovskite structure to accommodate ions of different types, sizes, and charges (limited only by the different versions of the tolerance factor) has placed it at the center of the research of thousands of scientists worldwide. It has already been indicated that if the desired property is not found for a particular combination of ions, they can be exchanged for others until we achieve the expected result (or at least an improved one). Sometimes this can be achieved by systematic or even random ion substitutions, but in other situations it is convenient to know beforehand the lattice effects that some ions can provoke (Attfield, 2001), isolate their influence, and get an optimized result. In this frantic process of making substitutions, some interesting compounds are “abandoned” in forgotten laboratory corners or drawers. Years later, someone studying another property finds that “your perovskite” is the most suitable for a new, unexpected cutting-edge application. This revival of already known phases usually ends up with the discovery of very impressive properties widely opening an unexplored perovskite research field. This is a recurring event with perovskites. The first example is probably that of barium titanate. A patent for its use as a white pigment was filed in 1917 but never used as such (Nielsen & Goldschmidt, 1920). However, its ferroelectric properties, discovered during World War II, caused a complete scientific and technological revolution in the electronics industry. Another classic example is that of the 1990s boom on colossal magnetoresistive manganites. In this latter case, the basic magnetic properties of manganites had already been studied in the 1950s (Rao, Cheetham, & Mahesh, 1996). This has occurred again with an old hexagonal acquaintance that has recently drawn the attention of those who work in photonics: BaTiS3. This thioperovskite adopts a classic hexagonal perovskite structure shown in Fig. 1.10, which can be

Fig. 1.9 An example of a workflow diagram including machine learning capabilities to accelerate the discovery of new oxide perovskites for their use as electrocatalyzers. In this example, the authors synthesize some perovskites (blue section); after testing their properties, they perform symbolic regression (a machine learning-related analysis, light brown boxes) to find a suitable descriptor for their compound search. Once the descriptor is chosen, a group of possibilities are selected (green boxes), and then the most promising ones are experimentally tested (dark brown boxes). From Weng, B., Song, Z., Zhu, R., Yan, Q., Sun, Q., Grice, C. G., et al. (2020). Simple descriptor derived from symbolic regression accelerating the discovery of new perovskite catalysts. Nature Communications, 11(1), 1–8. https://doi.org/10.1038/s41467-020-17263-9. Licensed under Creative Commons CC BY 4.0 license, http://creativecommons.org/licenses/by/4.0/.

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23

Fig. 1.10 (A) A comparison of the wavelength-dependence of the absolute birefringence values for various materials in the infrared. (B) As observed, BaTiS3 shows the highest reported birefringence among anisotropic crystals, and it is of an order of magnitude larger than quartz. Octahedra symbols indicate the fully transparent region of BaTiS3. Mid-wave infrared (MWIR) and long-wave infrared (LWIR) transparency regions are highlighted. B: Adapted by permission from Niu, S., Joe, G., Zhao, H., Zhou, Y., Orvis, T., Huyan, H., et al. (2018). Giant optical anisotropy in a quasi-one-dimensional crystal. Nature Photonics, 12, 392– 396. https://doi.org/10.1038/s41566-018-0189-1. Copyright 2018, Springer Nature.

considered a quasi-one-dimensional arrangement, a delight for optical studies. This thioperovskite was known for more than 60 years (Hahn & Mutschke, 1957), but no one had yet bet on it as a promising optical material. Nevertheless, it has been recently found to have giant optical anisotropy (Niu et al., 2018). Such materials are of technological importance in photonics due to their broadband birefringence and large dichroism, which might find use as polarizers, wave plates, or phasematching elements. Not bad for a perovskite forgotten for many years. The halide-perovskite case is also worthy of mention: many of the compounds that are now being studied for photovoltaic applications were already known more than six decades ago, especially the perovskite CH3NH3PbBr3 (Weber, 1978a). At that time, however, no one thought of using them for that purpose. The interest of the last decade in luminescent perovskites also fits into this wave of rediscovery (Kim, Cho, & Lee, 2016). Even if at the end of the 1990s some layered perovskites were proposed as potential organic-inorganic light-emitting diodes (Chondroudis & Mitzi, 1999), the explosion of research has occurred more recently, after the discovery of highefficiency quantum dot light-emitting diodes based on the previously mentioned halide perovskites (either with or without organic substituents) (Song et al., 2015; Zhang et al., 2015). Will the future of display technology rely on these perovskites? The rediscovery of known materials for different applications may occur fast, as has been the case of [TPrA]M[dca]3 (M ¼ dipositive metal ion) hybrid perovskites. These hybrids have been found to form stable glasses before they decompose (Shaw et al., 2021). So not only are they promising barocaloric materials (Bermu´dez-Garcı´a,

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Perovskite Ceramics

Sa´nchez-Andu´jar, Castro-Garcı´a, et al., 2017), but they might be used as thermoelectric materials (materials that can convert heat into electricity, for example). This promising application results from the increased electrical conductivity and lower thermal conductivity in the amorphous state versus the crystalline form. This latter finding opens the door to other perovskite-related glasses, which, if optimized, may be the core for the future recovery of waste heat from electric and electronic devices.

6

Conclusions

In this brief introduction, which has a significant bias due to my own preferences and limitations, an attempt has been made to show the enormous variability of properties found in materials with a perovskite structure. We have seen that the concept “material with a perovskite structure” goes beyond the first classic ceramic materials (calcium or barium titanate) and can be extended to include complex formulations with various types of polyatomic ions that occupy the positions of calcium, titanium, and oxygen. As the concept is expanded, the possibilities of preparing new compounds are notably increased. With more perovskites on “the market,” new properties appear (or those already known are improved), which multiplies the potential applications of this type of compounds. Today’s growing and unmanageable literature is a challenge for anyone looking to enter into the world of perovskites. In this summary, areas as broad as those of the ferroelectric perovskites or photovoltaic applications have not been considered (they will indeed be dealt with in other chapters of this book). We have also not addressed some groups of perovskites, such as hexagonal ones or even the various layered families. This does not mean that no exciting discoveries have been made in those areas; it means that there is no room here for everything. The world of perovskites is huge, not just because almost every day a new perovskite is synthesized but because the revival of already known phases accelerates the discovery of new ones and stimulates the research in related areas creating a mixed sense of frustration (too much to read) and excitement (so many interesting things to read). What hidden properties remain on the thousands of compounds already made? After more than 180 years of research, the field of perovskites research is far from being fully harvested.

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Methods for the synthesis of ceramic materials with perovskite structure

2

Jos
e Luis Clabel Huama´na, J.C. Sczancoskib, Euclydes Marega, Jr.a, and Alexandre H. Pintoc a Department of Physics and Material Science, Sa˜o Carlos Institute of Physics, University of Sa˜o Paulo, Sa˜o Carlos, Sa˜o Paulo, Brazil, bCMDF-UFSCar, Federal University of Sa˜o Carlos, Sa˜o Carlos, SP, Brazil, cDepartment of Chemistry and Biochemistry, Manhattan College, Riverdale, NY, United States

1

Introduction

The rapid and continuous modernization of society has been the driving force to stimulate technological innovations in different fields of science and technology. For uninterrupted progress in the coming years, search and production of novel multifunctional materials for a wide range of potential technological applications are crucial (Hench et al., 2011; Passarelli et al., 2020; Pomerantseva et al., 2019). One of the fundamental concepts in materials science establishes that the final physicochemical properties of any material are a direct manifestation of its structural and morphological features (Shyamaldas et al., 2020). Examples of such features include the degree of crystallinity, chemical composition, chemical size and particle size distribution, particle shapes, agglomeration/aggregation tendency, chemical surface, specific surface area, porosity/densification, and so on (Ansari et al., 2019; Clabel, Awan, Pinto, et al., 2020; Suttiponparnit et al., 2011; Ternero et al., 2021). The manipulation and effective control of these correlated parameters for a specific purpose are seen as a complex puzzle for researchers working in this area. Hence, a critical step is selecting the appropriate synthetic approach to design the desirable material. In a consensus, researchers from academia and industry prefer choosing or developing synthetic routes of easy handling, cost-effective, environmentally friendly, scalable production, and replicable for a wide variety of materials (Kharissova et al., 2019; Kumar et al., 2021; Mahin et al., 2021; Yang et al., 2022). Depending on the particular synthesis method, different experimental variables play a key role in the performance of inorganic materials, including the produced form (thin film, powder, colloidal suspension, etc.), time and temperature processing, medium (solid or liquid), pH condition, precursor concentration, type of solvent, gas atmosphere, nature of the substrate employed in the deposition process, etc. (Arau´jo et al., 2016; Clabel et al., 2014, 2016; Clabel, Nicolodelli, et al., 2021; Ribut et al., 2019; Shahi et al., 2021). Moreover, the optimized combination of multivariate experimental variables is another critical issue in tuning the material’s physicochemical properties synthesized by suitable production technology. Perovskite Ceramics. https://doi.org/10.1016/B978-0-323-90586-2.00004-8 Copyright © 2023 Elsevier Ltd. All rights reserved.

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According to the literature (Ali et al., 2021; Clabel, Awan, Lozano, et al., 2020; De Oliveira et al., 2020; Rivera et al., 2014), all fabrication techniques employed for the formation and growth of inorganic materials are classified into two categories: bottom-up and top-down. In the top-down category, precursors in the bulk state are brought down to the nanoscale through physical, chemical, or mechanical processes. Despite the versatility in forming a wide variety of materials on a large scale, there is the probability of some unwanted effects in the final products as high concentration of defects, poor morphological control, and leaving remaining traces of impurities in the lattice (Clabel, Awan, Pinto, et al., 2020; Vinod & Jelinek, 2019). In contrast, the bottom-up methods are based on the spontaneous self-assembly phenomenon of individual atoms and molecules to form stable crystals or nanoparticles (Sneharani & Byrappa, 2020). Such methods are successfully employed for crystal growth (single particles or hierarchical architectures) at multiscale levels. Sometimes, the limitation of the low yield is compensated by the improved ability to provide a controlled growth of monodisperse nanoparticles (Clabel, Awan, Pinto, et al., 2020; Vinod & Jelinek, 2019). Thus, the present chapter introduces the most popular top-down and bottom-up approaches employed to form and grow perovskites, a promising group of ceramic oxides well known in the scientific community because of their advanced technological applications (Karvounis et al., 2020; Neophytou et al., 2019). The chapter describes the traditional top-down method known as a solid-state reaction. In the sequence, the main bottom-up methods, including coprecipitation, sol-gel processes, conventional and microwave-assisted hydrothermal/solvothermal routes, chemical vapor deposition, electron beam physical vapor deposition, molecular beam epitaxy, radio frequency magnetron sputtering, and pulsed laser deposition, have been briefly addressed.

2

Synthesis and processing of perovskite ceramics

This chapter describes the main methods widely used to make perovskite structures. They have been classified according to the state of the starting reactants: solids, nanoparticles, and thin films. Within each section, we proceed from the most energy-demanding processes, usually at high temperatures, to those that can take place at lower temperatures.

2.1 Solid-state reaction method The solid-state reaction method (SSRM) is based on preparing a mixture of the oxides or carbonates of the metallic ions. This procedure usually involves mixing the oxide or carbonate reactants by grinding them. Then, the mixture undergoes a milling procedure in a low boiling-point solvent, for example, ethanol. This milling procedure can be done either in a ball or in a planetary mill. After milling, the mixture is subjected to a heat treatment at temperatures usually higher than 1000°C to promote ionic diffusion through the reaction mixture necessary to form the product (Rasaki et al., 2021). The SSRM presents many advantages, such as a straightforward synthetic procedure, the adaptation of the process for different oxides, and the possibility of obtaining a gram-scale quantity of product (Seo et al., 2017; Yang et al., 2019). However, the

Synthesis of ceramic materials with perovskite structure

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SSRM presents some disadvantages, such as producing large particles and broad particle size distribution. The inhomogeneous composition of the products is also a concern, generally with undesired crystal phases or unreacted initial materials along with the desired material (Kakihana, 1996). Another disadvantage is the possibility of oxide volatilization during the extended high-temperature heat treatment (Bhalla et al., 2000). Besides the harm it can cause to the final product composition, it also imposes a safety risk when the volatile oxide is toxic. For instance, lead (II) oxide (PbO) can damage the lungs, liver, kidney, and brain cells, as indicated in laboratory studies with mice (Dumkova´ et al., 2017). Over the years, the SSRM has been a widespread method used to prepare several types of perovskite materials, such as titanates (ATiO3) (Buscaglia et al., 2005), ferrites (AFeO3) (Maleki, 2019), manganites (AMnO3) (Im et al., 2007), zirconates (AZrO3) (Tripathi et al., 2018), niobates (ANbO3) (Vittayakorn et al., 2008), and layered perovskites (Schaak & Mallouk, 2002). Recently, the SSRM has been used to prepare cathode materials for solid oxide fuel cells (SOFCs) (Gunanto et al., 2021). For instance, Yu et al. prepared strontium titanate doped with varying amounts of iron (SrFe1
xTixO3
δ), with x varying between 0 and 0.15 (Yu et al., 2014). They used SrCO3, Fe2O3, and TiO2 as sources of Sr+2, Fe+3, and Ti+4, respectively. The heat treatment was carried out in three cycles, the first one at 950°C for 10 h, the second one at 1200°C for 10 h, and the third one at 1300°C for 10 h. The highest conductivity (72 S cm
1) was obtained for the sample with x ¼ 0.05, heated at 650°C. The conductivity behavior vs. temperature presented a semiconductor to a metallic phase transition (Yu et al., 2014). These samples showed a coefficient of thermal expansion (CTE) decreasing with increasing Ti content. The possibility of tuning the CTE with varying Fe content is suitable for these SOFCs since the cathode CTE should be compatible with other SOFCs to decrease the chances of cracking and delamination in the fabrication and operation of SOFCs (Kaur & Singh, 2020). In another example of perovskites applied to SOFCs, Xue and coworkers used the SSRM to prepare the double perovskite YBaCo2
xFexO5+ δ (x ¼ 0.0, 0.2, 0.4, 0.6). The authors used Y2O3, BaCO3, Co3O4, and Fe2O3 as sources of metallic ions. The samples were ground, pressed into pellets, and calcined at 1000°C for 12 h. These pellets were turned into powder, ground again, pressed to turn into pellets, and finally sintered at 1100°C for 20 h (Xue et al., 2011). These samples presented increasing CTE with increasing Fe content. Another example of fuel cells is the protonic ceramic fuel cell (PCFC). In the PCFCs, the proton-conductive perovskite is used as an electrolyte. The PCFCs have lower operating temperatures and long-term operational stability than other fuel cell technologies. These advantages can be attributed to PCFCs transferring protons across their solid electrolytes, whereas SOFCs transfer O
2 (Rasaki et al., 2021). Barium zirconate (BaZrO3) is a base material for PCFC electrolytes, mainly because structural defects can be induced when BaZrO3 is doped by lanthanide ions, leading to the creation of oxygen vacancies. A comparison between the transferring reactions occurring in SOFCs and PCFCs is shown in Fig. 2.1. Thermochemical energy storage (TCES) devices are materials with the potential for storing energy due to thermochemical reactions. Metal oxides are used in TCES because of the redox reactions occurring at high temperature operating conditions.

34

Perovskite Ceramics

Fig. 2.1 Schematic representation of the ionic transport mechanism of the traditional ceramic fuel cells and PCFCs. Adapted from Hao, L., Yang, Y., Huan, Y., Cheng, H., Zhao, Y. Y., Wang, Y., Yan, J., Ren, W., & Ouyang, J. (2021). Achieving a high dielectric tunability in strain-engineered tetragonal K0.5Na0.5NbO3 films. npj Computational Materials, 7(1). https://doi.org/10.1038/s41524-02100528-2, with the permission of Nature.

Especially, metal oxides with perovskite structures are promising materials for thermochemical energy storage (TCES) because of their high oxygen mobility (Lucio et al., 2019). Aiming to explore this potential from perovskites, Lucio et al. prepared a chromium-doped calcium manganite (CaCr0.1Mn0.9O3
δ). The CaCr0.1Mn0.9O3
δ was prepared using CaCO3, MnO2, and Cr2O3. The starting materials were homogenized in a planetary ball mill for 48 h. Then, the mixture was dried at 85°C for 5 h and calcined at 1100°C for 14 h (Lucio et al., 2019). All the samples presented CaMn2O4 as a secondary phase. TGA/DSC isothermal and nonisothermal experiments revealed an oxidation temperature of about 90°C and energy storage density higher than 150 J/g. These results indicate that CaCr0.1Mn0.9O3
δ is a promising material for TCES devices. Thermochemical fuel production (TCFP) uses solar energy to convert CO2 or H2O into CO or H2. The TCFP is a means of CO2 valorization since the H2/CO mixture produced can be converted to liquid hydrocarbons through Fisher-Tropsch reactions (Haeussler et al., 2018). Metal oxide perovskites can dissociate CO2 or H2 according to the following sequence of reactions: ABO3

> ABO3
δ + δ=2 O2

(2.1)

ABO3
δ + CO2 ðH2 OÞ

> ABO3 + δ CO ðH2 Þ

(2.2)

In Eq. (2.1), the stoichiometric perovskite ABO3 is partially reduced to the nonstoichiometric compound ABO3
δ. This step is endothermic and uses solar heat. In

Synthesis of ceramic materials with perovskite structure

35

Table 2.1 List of materials prepared by the SSRM used in the TCFP applications, presenting the temperature of reduction, and the quantity of O2 and CO evolved. O2 evolved (μmol g2 1)

CO evolved (μmol g2 1)

Material

T reduction (°C)

1st Cycle

2nd Cycle

1st Cycle

2nd Cycle

LaCoO3 LaMn0.5Co0.5O3 LaFe0.75Co0.25O3 LaMn0.5Ni0.5O3 LaFe0.75Ni0.5O3 SrFeO3 SrFeO3 Ba0.5Sr0.5FeO3 Ba0.5Sr0.5Fe0.2Co0.8O3 La0.6Sr0.4FeO3 La0.5Sr0.5MnO3 La0.5Sr0.5CaO3 Y0.5Sr0.5MnO3 La0.6Sr0.4Co0.2Fe0.8O3 La0.5Sr0.5Mn0.5Fe0.5O3 La0.5Sr0.5Mn0.5Co0.5O3

1300 1400 1300 1400 1300 1200 1100 1000 1000 1200 1400 1400 1400 1200 1300 1300

369 83 59 54 114 795 782 582 550 337 248 311 551 465 214 538

86 84 74 59 84 47 52 31 – 27 141 132 90 46 78 91

123 145 117 97 150 100 96 136 – 53 269 210 112 62 135 152

22 152 65 112 62 72 69 78 – 18 215 168 105 24 79 125

Adapted from Nair, M. M., & Abanades, S. (2018). Experimental screening of perovskite oxides as efficient redox materials for solar thermochemical CO2 conversion. Sustainable Energy & Fuels, 2(4), 843–854. https://doi.org/10. 1039/c7se00516d.

the step represented by Eq. (2.2), the partially reduced compound ABO3
δ is oxidized by CO2 or H2O, producing CO or H2; this second step is exothermic (Nair & Abanades, 2018). For studying the TCFP, Nair and Abanades prepared different types of doped and undoped perovskites using the SSRM and measured their capacity for producing CO and H2. The results obtained by them are summarized in Table 2.1. The variety of perovskite metal oxides studied in that article confirms how the SSRM is versatile regarding obtaining several compositions and allowing the doping of these oxides with various metal ions. Additionally, these results review the ability of the SSRM to generate functional materials, despite all the compositional and particle size distribution shortcomings related to SSRM.

2.2 Coprecipitation method The coprecipitation method is based on the solubilization of the metal cation precursors in the solvent, followed by the addition of a precipitating agent. The precipitating agent causes the simultaneous precipitation of all metallic cations present in the final product. The precipitation can be caused either by forming an insoluble species or by the attainment of supersaturation conditions in the solution. Before adding the

36

Perovskite Ceramics

precipitating agent, it might be necessary to control parameters such as solution pH, temperature, or addition of surfactants. These parameters can influence the final material’s structure, morphology, and property. Additionally, the precipitation step may not be enough to render the product with the desired crystal structure. In this case, it might be necessary to perform some additional heat treatments to make the precipitate attain the desired crystal structure. The usual steps of the coprecipitation method are summarized in Fig. 2.2, adapted from Athayde et al. (2016). The development of perovskite-based catalysts is an area that has benefited from the coprecipitation method. For instance, LaAlO3 prepared by the coprecipitation method was used as a catalyst for the methane oxidative coupling reaction (Sim et al., 2019). In the preparation of the LaAlO3, the authors controlled the solution pH in the range between 6 and 10 units. The sources of La+3 and Al+3 were, respectively, La(NO3)36H2O and Al(NO3)39H2O, whereas the precipitating agent was Na2CO3. After the precipitation, the powder obtained was heated at 950°C for 5 h. The XRD results revealed that phase pure LaAlO3 was obtained only at pH 6, whereas at pH 7 and 8 it led to the formation of La(OH)3 and La2O3 as undesired phases. At pH 9, La2O3 was obtained almost exclusively, and at pH 10, La(OH)3 was the only phase obtained. The precipitating pH also influenced the morphology of the final

Precipitating agent Homogeneously dispersed particles

Precursor A Precursor B

Washing

Suspension of A and B

Decantating of particles by adding a precipitating agent (supersaturation condition)

Filtration

AB Heat treatment

Fig. 2.2 Schematic representation of the usual steps of the coprecipitation method. From Athayde, D. D., Souza, D. F., Silva, A. M. A., Vasconcelos, D., Nunes, E. H. M., Da Costa, J. C. D., & Vasconcelos, W. L. (2016). Review of perovskite ceramic synthesis and membrane preparation methods. Ceramics International, 42(6), 6555–6571. https://doi.org/10.1016/j. ceramint.2016.01.130, with the permission of Elsevier.

Synthesis of ceramic materials with perovskite structure

37

product. At pH 6, semi-spherical particles with diameters around 100 nm were obtained. At pH values 7 and 8, rod-shaped larger particles were obtained. At pH values 9 and 10, round-shaped particles with diameters around 1.5 μm were obtained. The LaAlO3 morphologies obtained at different precipitating pH values can be seen in Fig. 2.3. Regarding the catalytic performance, the best CH4 conversion was approximately 33%, and the samples obtained precipitated at pH 7 and 9. The highest yield, about 16%, was obtained by the sample precipitated at pH 8. The highest yield by the pH 8 sample was explained based on well-developed oxygen vacancies and electrophilic lattice oxygen. The oxygen evolution reaction (OER) is essential in the water-splitting process and metal-oxygen and metal-air batteries (Tahir et al., 2017). In this sense, Matienzo et al. prepared by coprecipitation the following six metal oxides: LaFeO3, LaCoO3, LaNiO3, PrCoO3, Pr0.8Sr0.2CoO3, and Pr0.8Ba0.2CoO3 to test them as electrocatalysts for the OER in alkaline water electrolysis (Matienzo et al., 2020). All the metal ions were introduced in the solution as their nitrate salts, and the KOH was used as the precipitating agent. The precipitates obtained on each synthesis were subjected to a thermal treatment at 700°C for 3 h to crystallize as the perovskite metal oxide of the intended composition. After the thermal treatment, all the oxides were obtained with the intended composition, except for LaFeO3, which presented La2O3 as a secondary phase. In relation to the OER, the best performance was obtained by Pr0.8Sr0.2CoO3, which presented OER activity comparable to the one obtained by the commercial Ru nanoparticles supported on carbon (Matienzo et al., 2020).

Fig. 2.3 SEM images of the LaAlO3 products after heat treatment for precipitating pH values between 6 and 10. From Sim, Y., Yoo, J., Ha, J. M., & Jung, J. C. (2019). Oxidative coupling of methane over LaAlO3 perovskite catalysts prepared by a co-precipitation method: Effect of co-precipitation pH value. Journal of Energy Chemistry, 35, 1–8. https://doi.org/10.1016/j.jechem.2018.10.002, with the permission of Elsevier.

38

Perovskite Ceramics

Still, about the use of coprecipitation to prepare catalytic materials, Li et al. prepared lanthanide-doped BaCeO3
δ-doped with lanthanide (LN) ions, such as La+3, Y+3, or Pr+3 (Ba(Ce0.9LN0.1)O3
δ) (Li et al., 2021). These Ba(Ce0.9LN0.1)O3
δ materials were used as catalytic supports impregnated by Ru, the catalyst used in the ammonia synthesis reaction. BaCeO3
δ was synthesized by mixing Ba(NO3)2 and Ce(NO3)3 solutions, in the mole ratio of 8:10 for Ba+2:Ce+3, followed by the addition of an ammonium oxalate solution as the precipitating agent. The precipitate formed was dried and heat treated at 800°C for 6 h. For preparing the Ba(Ce0.9LN0.1)O3
δ materials, the only modification was the addition of the nitrate salt of the lanthanide ion along with the mixture of Ba(NO3)2 and Ce(NO3)3. In this case, the mole ratio was Ba+2:Ce+3:LN+3 ¼ 8:9:1. The XRD results showed that all the BaCeO3
δ and Ba(Ce0.9LN0.1)O3
δ samples presented small quantities of BaCO3 and CeO2. However, BaCO3 disappeared from all samples after catalyst activation with an H2/N2 gas mixture, probably due to BaCO3 amorphization. Regarding the ammonia synthesis reaction, the best performance was obtained for the catalysts made up of 2.5% Ru/BaCe0.9La0.1O3
δ, followed, in order, by 2.5% Ru/BaCe0.9Y0.1O3
δ, 2.5% Ru/BaCeO3
δ, and 2.5% Ru/BaCe0.9Pr0.1O3
δ. The best performance of 2.5% Ru/BaCe0.9La0.1O3
δ could be due to an enhancement of the strong metal-support interaction (SMSI) effect between the La+3 and Ru. This higher SMSI effect could lead the support to form strong chemical bonds with Ru and inhibit the Ru particle aggregation (W. Li, et al., 2021). Lanthanide ions are useful not only as doping ions but also as A-type constituents of the perovskites. For instance, Nguyen et al. prepared HoFeO3 powders to be used as anodes in lithium-ion batteries. Fe+3 and Ho+3 solutions in ethanol form precipitate upon the addition of ammonia (Nguyen et al., 2022). The resulting precipitate was treated at 850°C. The XRD indicated the formation of a crystalline single-phase HoFeO3 material. TEM analysis revealed interconnected particles with diameters of approximately 50–100 nm. Energy-dispersive X-ray spectroscopy (EDX) coupled to the TEM revealed homogeneous distribution of Ho, Fe, and O throughout the sample, as can be seen in Fig. 2.4. The electrochemical characterization indicated high-capacity retention, good cyclability, and high Li+ diffusion coefficient, which are reasonable parameters for HoFeO3 application in lithium-ion batteries. The coprecipitation method is also useful when preparing complex perovskite metal oxides. For example, perovskites are doped with more than one type of metallic ion. For instance, Varandiili et al. prepared LaFeO3 codoped with Co+2 and Pd+2 (LaFe1
x
yCoxPdyO3 [(x, y) ¼ (0,0), (0.40, 0), (0.38, 0.05)]) (Varandili et al., 2018). The sources of all metallic ions were their nitrate salts, and three different precipitating agents were used in this study: NH4OH, NaOH, and (NH4)2CO3. Heat treatment was performed either at 800°C, 900°C, or 1000°C for 1 h. In general, the LaFeO3, LaFe0.6Co0.4O3, and LaFe0.57Co0.38 Pd0.05O3 powders presented single-crystalline materials when calcined at 800°C or 900°C, regardless of the precipitating agent. Also, in general, using NH4OH yielded smaller particles than using NaOH.

Synthesis of ceramic materials with perovskite structure

39

Fig. 2.4 Electronic microscopy characterization of the HoFeO3 nanocrystals: (a) Bright-field transmission electron microscopy (BF-TEM) image, (b) High-resolution transmission electron microscopy (HR-TEM) image, (c) Selected area electron diffraction (SAED) pattern, (d) Energy dispersive X-ray spectroscopy (EDX) mapping of different elements present in the HoFeO3 samples, from left to right: Dark-field image, analysis of Ho (shown in yellow), analysis of Fe (shown in light blue), and analysis of O (shown in light green). From Nguyen, A. T., Phung, V. D., Mittova, V. O., Ngo, H. D., Vo, T. N., Thi, M. L. L., Nguyen, V. H., Mittova, I. Y., Le, M. L. P., Ahn, Y. N., Kim, I. T., & Nguyen, T. L. (2022). Fabricating nanostructured HoFeO3 perovskite for lithium-ion battery anodes via co-precipitation. Scripta Materialia, 207. https://doi.org/10.1016/j.scriptamat.2021.114259, with the permission of Elsevier.

Additionally, as more doping ions are incorporated, the smaller the particle size becomes for the same precipitating agent, according to Scherrer’s equation calculations.

2.3 Sol-gel process and its variants Metal or nonmetal alkoxides are compounds where there are metal or nonmetaloxygen-carbon bonds. A classic example of a nonmetal alkoxide is tetraethyl orthosilicate (Si(OC2H5)4), popularly known as TEOS. In comparison, a classic example of metal alkoxide is titanium isopropoxide (Ti(OCH(CH3)2)4). The sol-gel method became initially recognized by the hydrolysis and condensation reaction of alkoxides, such as TEOS (Cushing et al., 2004). In general, the hydrolysis of alkoxides in water can be represented by the following equation (Burda et al., 2005): M ðORÞx + n H2 O ! M ðORÞx
n ðOHÞn + nROH

(2.3)

40

Perovskite Ceramics

In Eq. (2.3), M(OR)x represents the alkoxide, where M can be a metal or a nonmetal. The species M(OR)x
n(OH)n is the product of the alkoxide hydrolysis, which can further undergo a polycondensation step as shown in Eq. (2.4) (Burda et al., 2005): 

2 M ðORÞx
n ðOHÞn ! ðOHÞn
1 ðORÞx
n
M
O
M ðORÞx
n ðOHÞn
1



+ H2 O (2.4) The polycondensation product, [(OH)n
1 (OR)x
n
M
O
M (OR)x
n (OH)n
1], is an inorganic polymer forming a three-dimensional network due to metal or nonmetal oxyanions (Leite, 2004). This inorganic polymer can undergo an aging procedure, where the polycondensation reactions continue until the polymer is converted to a gel. Then, this gel should be dried to remove water and other volatile liquids from the gel network. If the gel is dried to a thermal procedure, it will yield a xerogel, whereas if it is dried under supercritical conditions, it will yield an aerogel. Finally, after the drying process, the xerogel or aerogel undergoes heat treatment, such as calcining or sintering, to fill the voids left by the evaporation of the volatile liquid (Cushing et al., 2004). This step tends to cause particle agglomeration, which can be avoided by tuning variables such as pH, temperature, and the use of surfactants. A variant of the sol-gel method commonly used to prepare metal oxides with perovskite structures is the Pechini method. Metallic salts are dissolved in a glycol solution, i.e., ethylene glycol, containing some chelating agent, such as citric acid, glycolic acid, lactic acid, or EDTA, in excess. This excess is necessary for the chelating agent to form metal complexes with the metallic salt. When this glycol, chelating agent, and metallic salt solution are heated up at temperatures of approximately 150°C, some esterification reactions start to happen between the carboxyl group of the chelating agent and the hydroxyl group of the glycol. The condensation of these two functional groups in a new bond is accompanied by water elimination (Kakihana, 2009). Due to the relatively high concentrations of the glycol and the citric acid, this esterification reaction happens in many different points of the reaction medium; an intricate polymeric network is formed, which in the macroscopic scale can be observed as the formation of a highly viscous gel. This gel is heated up at temperatures of approximately 300°C to eliminate most of the organic material and convert the material from a gel-like to a powder-like aspect. Then, this powder is subject to calcination procedures in temperatures higher than 500–900°C to produce the desired binary or multinary transition metal oxide. The main steps of the Pechini method are presented in Fig. 2.5 (Danks et al., 2016). One advantage of the Pechini method is that the metallic cations are well dispersed in the polymeric network during the gel formation. This fact tends to generate a more compositionally homogeneous final product. Consequently, it helps to reduce the chances of phase segregation, which is a prevalent issue, mainly in the case of the synthesis of multinary materials (Leite, 2004).

Fig. 2.5 Schematic representation of the Pechini method to prepare polyesterified gels containing metallic cations. From Danks, A. E., Hall, S. R., & Schnepp, Z. (2016). The evolution of “sol-gel” chemistry as a technique for materials synthesis. Materials Horizons, 3(2), 91–112. https://doi.org/10.1039/c5mh00260e, with the permission of the Royal Society of Chemistry.

42

Perovskite Ceramics

However, the method presents some shortcomings as well. For instance, the high content of organic compounds is one of them. This high content of organic compounds can be detrimental for many final applications as a high content of residual carbon can be left in the product, even after multiple heat treatments (Mackenzie, 1988). Another shortcoming is the difficulty separating the nucleation and growth steps of the nanoparticle formation, which tends to generate agglomerated powders (Hodgson & Weng, 2006). Hammouda et al. used a variation of the Pechini method to prepare different cobaltates, such as BaCoO3, CeCoO3, LaCoO3, and SrCoO3 (Hammouda et al., 2017). These materials were prepared by reacting their respective metal nitrate aqueous solution with citric acid in excess. The resulting solution was heated at 100°C to eliminate the water excess, leading to a spongy material. This spongy material was dried at 100°C and then calcined at 700°C to obtain the desired perovskite material. All materials were phase pure according to the XRD results. Then, these materials were applied as catalysts for phenol degradation and mineralization by adding potassium monopersulfate (PMS). The phenol degradation-mineralization activity presented the following trend from the highest to the lowest active catalyst: SrCoO3, LaCoO3, BaCoO3, and CeCoO3. The authors explained this activity trend with the crystallinity of the materials. The more crystalline the materials were, the better was their phenol degradation-mineralization activity. Still, in the area of phenol degradation, Wang et al. prepared CaMnO3 using a modified sol-gel method by using aqueous solutions of Ca(NO3)24H2O and Mn(NO3)2 with ethylene glycol and citric acid (T. Wang, et al., 2020). The liquid part of the solution was evaporated at 85°C, producing a sol, which was dried at 80°C overnight to produce a gel. This gel underwent two additional heat treatments, the first at 400°C for 2 h, and the second at 900°C for 3 h. For the phenol degradation experiments, the peroxydisulfate was used to generate radical species. Besides CaMnO3, other Mn-based catalysts, such as β-MnO2, Mn2O3, CaMn2O4, and Ca2MnO4 were tested. Among all these catalysts, the best activity was obtained for CaMnO3. The explanation for this result is based on the fact that it is the tolerance of Mn multivalence (Mn+3 and Mn+4) in the CaMnO3 perovskite structure. GdCoxFe1
xO3 (x ¼ 0; 0.2; 0.5; 0.8; 1) was tested as a catalyst in the dry reform of the methane reaction (Yafarova et al., 2019). In the sol-gel version used to prepare these materials, metallic nitrates were used as the source of metallic ions, dissolved in deionized water and citric acid. This solution was dehydrated at 90°C, then heated at 240°C, yielding a powder. This powder was further calcined at 500°C for 2 h. The XRD results indicated the formation of single-phase materials. In general, the catalytic tests presented higher activity as the cobalt content in the catalyst increased. One possible explanation for this trend is that the cobalt presence in the B-site of the perovskite catalyst would help suppress the occurrence of parallel reactions, such as the water-gas shift reaction, which would decrease the rate of the dry reforming methane reaction. One demonstration of the ability of the sol-gel method to generate metal oxides with complex compositions is provided in the paper published by Lee et al. (2019).

Synthesis of ceramic materials with perovskite structure

43

In that paper, the author prepared the PrBa0.5Sr0.5Co2
xNixO5+ δ (x ¼ 0, 0.1, 0.2, and 0.3) with the double-perovskite structure. These powders were prepared by dissolving metallic nitrates, ethylene glycol, and citric acid in distilled water. The solvent was evaporated, leading to a gel formation. The gel was heat-treated at 280°C, further treated at 600°C for 4 h to remove organic impurities. Finally, the resulting powder was heat-treated at 1000°C for 4 h. In general, the PrBa0.5Sr0.5Co2
xNixO5+ δ powders presented were obtained as phase-pure materials, except for when x ¼ 0.3, which presented some Ba4Co3O10 as the secondary phase. The PrBa0.5Sr0.5Co2
xNixO5+ δ double-perovskite powders were tested as catalysts both for the oxygen reduction reaction (ORR) and oxygen evolution reaction (OER) for rechargeable zinc-air batteries. The best electrochemical performance was obtained for the sample with x ¼ 0.1, which was attributed to the different chemical states of the cobalt cations due to the presence of nickel cations in the lattice. The sol-gel method is also appropriate for preparing doped ferrite materials for gassensing applications. For instance, Cao et al. prepared lanthanum ferrite doped by chloride (LaFeO3
xClx) for x ¼ 0, 0.15, 0.3, and 0.6 (Cao et al., 2017). These powders were prepared from La(NO3)36H2O and Fe(NO3)39H2O. The Cl doping was accomplished by partially replacing the Fe(NO3)39H2O with FeCl36H2O. The metallic cations were dissolved in deionized water. Then, polyethyleneglycol was added to the mixture at 80°C, forming a sol, which was dried to form a gel. The gel was dried and ground to be calcined at 600°C for 2 h. The Cl doping was assessed by EDX elemental mapping in the SEM. In comparison, XPS results indicated that the Cl might be present in the lattice and the surface as FedCl. These powders were used for sensing the following gases: ethanol, DMF, acetone, NH3, and CH2Cl2. Ethanol vapor was the gas detected with the highest response for all the LaFeO3
xClx compositions, the levels of ethanol increased with increasing Cl content.

2.4 Hydrothermal and Solvothermal methods Historically, in the middle of the 19th century, Sir Roderick Murchison (1792– 1871) was the first to use the word “hydrothermal” (“hydro” ¼ water; “thermal” ¼ heat) to describe the influence of water at elevated temperature and pressure on the geological formation processes of different rocks and minerals (Somiya, 1990). In 1845, the pioneering study performed by Schafthaul et al. (Schafthaul, 1845) on the growth of tiny quartz crystals in a Papin’s digester was considered the first publication on hydrothermal research (Fig. 2.6). In the past two decades, hydrothermal routes have revolutionized nanoscience and nanotechnology due to the flexibility of producing several inorganic materials with improved physicochemical properties (Wojnarowicz et al., 2020). According to the literature (Byrappa & Yoshimura, 2012), hydrothermal synthesis is defined as any heterogeneous reaction performed in an aqueous medium under high temperature (>100°C) and pressure (>1 atm) in a closed system. In particular, the conventional hydrothermal (CH) reactions involve heating the chemical solution in closed vessels, inducing an autogenous pressure inside the system as the temperature

44

Perovskite Ceramics

Fig. 2.6 Typical hydrothermal autoclave manufactured in stainless steel, illustrating the respective main parts. From TOPTION, Teflon Lined Hydrothermal Synthesis Autoclave Reactor, http//Pt. Toptiontech.Com/Laboratory-Reactor/Hydrothermal-React. (2021).

exceeds the typical boiling point of the solvent (Einarsrud & Grande, 2014). For industrial and laboratory operations, such reactions for the growth of metal oxides can be conducted at temperatures and pressures at subcritical and supercritical water states (Aux
em
ery et al., 2020; Darr et al., 2017). Among various solution-based synthetic approaches available currently for advanced material processing/synthesis (precipitation, sol-gel, sonochemical, and oxidant peroxo methods) (Camargo & Kakihana, 2001; Kamali et al., 2021; Styskalik et al., 2017), the CH route is technologically attractive and is often preferred for many types of research due to its simplicity, easy handling, and low cost (Yoshimura & Byrappa, 2008). In a hydrothermal reaction, the formation of the desired crystalline phase as well as the control/manipulation of morphological features (size, shape, and size distribution) is dictated by adjusting different experimental parameters, such as the type of chemical precursor (chloride, nitrate, acetate, etc.), reaction temperature, time processing, solute concentration, pH condition, the addition of growth modifier agents (mineralizers, surfactants, polar organic molecules, hydrophilic polymers, etc.) (Liu & Shaw, 2016; Mamakhel et al., 2020; Bhosale et al., 2015; Rao et al., 2012). When the chemical reactions inside a closed vessel are carried out in nonaqueous media (alcohol, organic solvents, etc.) or organic solvent-water mixtures, the processing is more appropriately called solvothermal synthesis (Burda et al., 2005). Hydrothermal/solvothermal syntheses are carried out in a specific apparatus, popularly called an autoclave, which must resist corrosive solvents and high temperatures and pressures along with the reactions (2012). For safety issues, such autoclaves are commercially machined from high-strength metal alloys, especially stainless steel, without any welding. To avoid direct contact with the inner walls of the autoclave, the solution can be placed inside a Teflon chamber, which acts as an additional protective barrier to prevent corrosion and contamination from the steel. Generally, the total volume of the Teflon chamber is filled below 80% with the desired precursor

Synthesis of ceramic materials with perovskite structure

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solution to guarantee a suitable vapor pressure (Einarsrud & Grande, 2014; Rabenau, 1985). Another essential feature is that autoclaves are designed to be powered by indirect heating systems as laboratory ovens, furnaces, and collar-type electric resistance heaters. Fig. 2.6 shows two typical standard autoclaves employed in CH synthesis. In hydrothermal reactions, the physicochemical properties of water, such as ionic product, surface tension, density, dielectric constant, and viscosity, are significantly affected by the increase of temperature and pressure (Peterson et al., 2008; Rabenau, 1985). Hence, the formation and growth mechanisms of metal oxides via hydrothermal processing is a complex puzzle, so that one of the most acceptable is based on dissolution-precipitation processes (Ahn et al., 2013; Canu & Buscaglia, 2017). In a brief explanation, temperature gradients are maintained in the solvent due to an uneven heat transfer through the autoclave walls. Such a phenomenon promotes a density difference in the solution, leading to buoyancy-driven convective flows (Ma et al., 2020) (Fig. 2.7A). The dissolution of precursors in the form of ions or molecules MICROWAVE HEATING GLASS REACTION VESSEL

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Fig. 2.7 Schematic representation of (A) conventional and (B) microwave heating mechanisms in hydrothermal reactions carried out in aqueous medium. In conventional heating (A), both thermal conduction and convection govern the heating process. In contrast, microwave heating (B) is able to provoke a fast and uniform heating of the solution due to the water molecular motion in the external oscillating electric field of microwaves. From Sweygers, N., Alewaters, N., Dewil, R., & Appels, L. (2018). Microwave effects in the dilute acid hydrolysis of cellulose to 5-hydroxymethylfurfural. Scientific Reports, 8(1). https://doi.org/10.1038/s41598-018-26107-y with the permission of Nature Research.

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occurs in the hottest zone of the autoclave. On the other hand, the solvent supersaturation condition is reached by transporting ions or molecules to the coldest zone (Yang & Park, 2019). The supersaturation plays a significant role in the nucleation and crystal growth processes to precipitate the desired products (Bahrig et al., 2014). In general, the addition of stabilizing agents in the reaction medium is a strategy often employed in CH to suppress the aggregation state of clusters and particles, control the size, and manipulate the growth direction and the exposed crystallographic facets of the final crystals (Basnet & Chatterjee, 2020; Einarsrud & Grande, 2014) (Fig. 2.7). Enormous progress has been made in producing perovskites with improved multifunctional properties through morphological control in hydrothermal/ solvothermal systems (Feng et al., 2005; Kalyani et al., 2015; Niederberger et al., 2004). This control has been normally achieved by adjusting different reaction parameters such as temperature, pressure, power density, gas flow rate, etc. In this context, Chen, Zheng, et al. (2019) investigated the effect of different reaction times on the growth of butterfly-like calcium titanate (CaTiO3) dendrites under hydrothermal treatment. Tetramethylammonium hydroxide was employed as a mineralizer and surface modifier. For times less than 30 min, the samples exhibited irregular and compact agglomerates, indicating a high concentration of amorphous precursor in the composition of the final product. Crystallization of dendrites was only observed with times longer than 60 min. Such hierarchical microarchitectures demonstrated a promising photocatalytic response for the degradation of organic dyes under ultraviolet-visible irradiation. In a further study, Chen, Bao, et al. (2019) hydrothermally synthesized CaTiO3 tetragonal microrods by using different potassium hydroxide (KOH) concentrations (from 0.25 to 6.0 mol L
1). The authors concluded that using potassium titanate (K2Ti6O13) nanofibers as a titanium source played a crucial role in inducing the growth of these microrods. The experimental results also proved that KOH concentration directly impacted particle size and surface features. At low KOH concentrations (from 0.25 to 6.0 mol L
1), unique V-type holes along the [010] crystallographic direction were noted on two end crystal facets of the tetragonal microrods. Such peculiarity was ascribed to a fraction of K2Ti6O13 nanofibers still embedded in CaTiO3 microrods, due to a slow formation and crystallization kinetics of this perovskite. The hydrothermal reactions have also been regarded as a versatile approach to building perovskite-based semiconductor heterostructures to produce potential photocatalysts for water splitting (Bin Adnan et al., 2018). For example, Ha et al. (2016) synthesized strontium titanate/titanium oxide (SrTiO3/TiO2) heterostructures via hydrothermal processing (160°C for 12 h) by using morphology-controlled TiO2 (nanocubes, nanoparticles, nanospheres, and nanofibers) as both template and precursor in a strontium hydroxide (Sr(OH)2) solution. An essential finding of this study was that, by adjusting the Sr/Ti molar ratios, the structure, interface, and phase component of SrTiO3/TiO2 heterostructures were tailored. According to these authors, SrTiO3 crystals nucleated and grew on the surface of TiO2 nanocrystals via a dissolution-precipitation mechanism involving a reaction between [Ti(OH)6]2
and Sr2+ species.

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2.5 Microwave-assisted methods The conventional hydrothermal/solvothermal technologies are based on indirect heating processes, i.e., heat conduction proceeds from autoclave walls to the central volume of the solution. Hence, heat is transferred from outside to inside in solid particles during the hydrothermal growth stage (Schanche, 2003; Sweygers et al., 2018). Depending on wall thickness, autoclave material, and input processing parameters (heating rate, reaction time, and temperature), thermal energy losses and nonuniform heating of the solution are practically inevitable in these reaction systems (Gude et al., 2013; Roy & Prasad, 2018). Such setbacks can lead to long processing times to produce and crystallize the desired inorganic materials, sometimes yielding a morphological system composed of a broad particle size distribution (Zito et al., 2018). These drawbacks could only be overcome when Komarneni et al. (1992) innovated the traditional hydrothermal systems by implementing microwave radiation as an energy source to synthesize crystalline inorganic materials. This new technology, termed by them as “microwave-hydrothermal process” (Leonelli & Komarneni, 2015), revolutionized the field of solution-based chemical routes, bringing many benefits such as an increase in reaction kinetics by one to two orders of magnitude, formation and selective crystallization of inorganic phases, excellent control of the reaction parameters, and improved reproducibility of synthesis (Gao et al., 2021; Komarneni et al., 1999; Roy & Prasad, 2018). Due to the successful use of microwaves, electromagnetic waves in the frequency range 0.3–300 GHz, the temperature can be raised uniformly throughout the entire volume of the solution (volumetric heating), eliminating the thermal gradients and favoring homogeneous nucleation with posterior controlled growth (Rosa et al., 2014; Schanche, 2003) (Fig. 2.7B). Fundamental studies have explained microwave heating by two main mechanisms, namely dipolar polarization and ionic conduction (Bilecka & Niederberger, 2010; Kappe et al., 2009; Sweygers et al., 2018; Xiong et al., 2021) (Fig. 2.8). In the first

Fig. 2.8 Schematic representation of (A) dipolar polarization and (B) an ionic conduction mechanism responsible for the microwave heating. From Martins, C. P. C., Cavalcanti, R. N., Couto, S. M., Moraes, J., Esmerino, E.A., Silva, M.C., … Cruz, A. G. (2019). Microwave processing: Current background and effects on the physicochemical and microbiological aspects of dairy products. Comprehensive Reviews in Food Science and Food Safety, 18(1), 67–83, with the permission of Wiley.

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one, the molecules periodically rotate back and forth, trying to align their dipoles with the oscillating electric field of microwaves at a specific frequency. This cyclic process results in molecular friction and collision events with neighboring molecules, whose lost energy is transformed into heat energy for the reaction medium. When the molecular dipoles do not have enough time to realign (when the frequency of the electric field is higher than the relaxation time of the dipoles) or reorient (when the time to change the direction of the electric field is longer than the response time of the dipoles) themselves with the applied electric field, there is no heating. Hence, the time of orientation and disorientation of the polar molecules is dependent on the frequency of microwave radiation, reflecting on the amount of heat generated in the system. For example, at a typical microwave frequency of 2.45 GHz, the molecular dipoles have time to align with the oscillating electric field but not to follow it precisely (Bilecka & Niederberger, 2010; Gude et al., 2013; Kappe, 2004; Kappe et al., 2009) (Fig. 2.8). The second main microwave heating mechanism is ionic conduction. During microwave radiation, the ionic species in solution instantaneously respond to the applied electric field, i.e., they continuously oscillate around an equilibrium point. This periodic movement promotes several collisions between the charged species with neighboring molecules or atoms, causing rapid heating of the system. In terms of effectiveness in producing heat, ionic conduction has a superior effect on dipolar polarization; however, the synergism of both mechanisms is responsible for the overall increase of temperature (Bilecka & Niederberger, 2010; Gude et al., 2013; Kappe, 2004; Kappe et al., 2009). Unlike conventional heating, the microwave heating of a given substance (material or solvent) is dependent on its intrinsic dielectric properties, including the dielectric constant (ε´-the ability of a material to be polarized by an external electric field) and dielectric loss (ε00 -effectiveness of a material to convert electromagnetic energy into thermal energy). The ratio between these two dielectric variables is the so-called loss tangent ( tan δ ¼ ε 00 /ε0 ), which can estimate the ability of a material to convert electromagnetic energy into heat at a given frequency and temperature. According to the tan value, the microwave absorption properties of a substance to produce a rapid heating are categorized as high (tan δ > 0.5), medium (tan δ  0.1–0.5), and low (tan δ < 0.1) (Bilecka & Niederberger, 2010; Kappe, 2004; Kappe et al., 2009). Some of the most common solvents used for the preparation of inorganic materials under microwave-assisted hydro
/solvothermal reactions are water ( tan δ  0.123), dimethyl sulfoxide (tan δ  0.285), ethylene glycol (tan δ  1.350), ethanol (tan δ  0.941) and N-methyl-2-pyrrolidone (Li et al., 2013; da Pereira et al., 2019; Y. Wang, et al., 2020; Yoon et al., 2012). The empirical data on the tan of these different solvents were recorded at a microwave frequency of 2.45 GHz at room temperature (Hayes, 2002). In addition, polytetrafluoroethylene (commercially known as Teflon) is classified as a microwave transparent material (i.e., microwaves pass through without any losses) due to its low tan ( 0.0003) (Rajesh et al., 2009). Therefore, this polymer is widely chosen to manufacture autoclaves dedicated to microwave-aided reactions. Fig. 2.1 shows two typical standard autoclaves employed in microwaveassisted hydrothermal systems. Good electrical conductors such as metal alloys must be avoided in microwave-assisted reactions since the microwaves can produce a

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concentrated electric field at corners and/or edges in the materials, resulting in electrical sparks and even fires (Chandra et al., 2011). Given the importance of the interaction of microwaves with the substances, published studies have described that the pronounced acceleration in reaction rates commonly perceived in microwave-assisted synthesis is due to a combination of thermal and nonthermal effects (de la Hoz et al., 2005; Herrero et al., 2008). Thermal effects arise from microwave dielectric heating mechanisms such as superheating, hot spots, and selective absorption of microwave radiation by polar substances. These phenomena are not evidenced in a traditional heating process (i.e., involving thermal conduction and convection) (Galema, 1997). On the other hand, nonthermal effects are not entirely known yet, being a controversial issue and constantly debated in the scientific community. Published studies have postulated that such effects are triggered by the direct interaction of the microwave electric field with specific molecules, intermediates, or even transition states in the reaction medium, but without affecting the overall temperature (Kappe, 2004). Hence, nonthermal effects are related to the inherent features of microwaves and not too different temperature regimes (Galema, 1997). Although the microwaves do not have sufficient energy to break chemical bonds, there is an assumption that nonthermal effects can increase polar molecules’ collision frequency and efficiency (through mutual orientation) in microwave-assisted reactions (Perreux & Loupy, 2001; Zito et al., 2018). In this regard, experimentally identifying and separating the contribution of thermal and nonthermal effects is a complex challenge to describe any mechanistic model. Some experimental issues to prove the existence of nonthermal effects are mainly caused by inaccurate temperature reading, thermal runaway, nonuniform heating of the medium, and other experimental errors (de la Hoz et al., 2005; Nozariasbmarz et al., 2018). The use of microwave-assisted hydrothermal/solvothermal routes has been successfully employed to synthesize perovskites (Da Silva et al., 2019; Moreira et al., 2009). Moreira et al. (2009) reported forming a monodisperse system composed of microcube-shaped CaTiO3 mesocrystals via a microwave-assisted hydrothermal method. They demonstrated that a short time of 10 min was enough to obtain these superstructures with a high degree of crystallinity. Based on both experimental data and theoretical calculations, the broad visible photoluminescence emissions were explained by the appearance of intermediary energy levels in the band gap due to intrinsic defects arising from tilting in TiO6 octahedra.

2.6 Chemical vapor deposition (CVD) Chemical vapor deposition (CVD) is a technique where a thin film is deposited onto a solid substrate from one or more gas-phase precursors. CVD usually needs a reaction activator, such as light, heat, high-energy radiation, or plasma (Raiford et al., 2020). The CVD process starts with the introduction of the gas-phase precursors. According to how the gas phase is input in the system, CVD can usually be classified as aerosol-assisted CVD (AACVD), direct liquid injection CVD (DLICVD), metalorganic CVD (MOCVD), and hybrid physical CVD (HPCVD) (Liu et al., 2019).

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Then, the gas-phase precursors adsorb onto the substrate surface. Once it happens, chemical reactions will convert these adsorbed precursors into solid deposits onto the already present substrate. In this way, the solid layer on the substrate progressively grows. At the same time, these chemical reactions can generate byproducts, which can desorb from the substrate and be eliminated from the thin film as gaseous species (Apreutesei et al., 2011). One conceptual difference between CVD and its counterpart physical vapor deposition (PVD) is that the precursor adsorption onto the substrate relies solely on physical interactions, without the occurrence of chemical reactions (Apreutesei et al., 2011). Different precursors, such as halides, hydrides, and metalorganic compounds, for instance, metal alkyls, metal alkoxides, and metal carboxylates, can be used in the CVD process (Apreutesei et al., 2011). Some advantages of CVD in relation to other deposition techniques are the relatively lower operating temperature and the possibility of preparing different materials such as metals and non-metal elements, oxides, nitrides, hydrides, sulfides, and polymers. Furthermore, CVD allows deposition in complex substrates, large-size, and multisubstrate deposition (Apreutesei et al., 2011; Liu et al., 2019). Despite all these advantages, CVD also presents some disadvantages, such as the high cost and toxicity of some precursors; the reaction byproducts usually can be toxic, corrosive, or harmful, such as CO, HCl, and H2 (Apreutesei et al., 2011). Throughout the literature, it has been discussed that CVD has been used to deposit thin films of different perovskite materials such as titanate, ferrites, and manganites. Although the recent literature has concentrated most of its efforts on metal halide perovskites for solar cell applications (D. Li, et al., 2021), there is still space and interest in metal oxide perovskite materials, which are the foci of this chapter. For instance, manganites are a class of appealing materials for the current literature due to the correlation of structural, magnetic, and transport properties (Pla et al., 2017). In this sense, Catalano et al. prepared Pr0.7Ca0.3MnO3 using MOCVD; the precursors used were Pr(hfa)3diglyme, Ca(hfa)2tetraglyme and Mn(tmhd)3, where hfa ¼ hexafluoroacetylacetonate, tmhd ¼ tetramethylheptandionate, diglyme ¼ CH3(OCH2CH2)2OCH3, and tetraglyme ¼ CH3(OCH2CH2)4OCH3 (Catalano et al., 2015). These precursors were vaporized at 130°C. Deposition was carried out at 900°C for 60min, yielding a film thickness of 250 nm onto the SrTiO3 (100) and SrTiO3 (110) substrates. XRD and TEM confirmed the epitaxial growth of the Pr0.7Ca0.3MnO3. In the same paper, magnetic field and zero-field cooling measurements indicated a ferromagnetic transition temperature. Huertas-Flores et al. used laser-assisted CVD to prepare NaTaO3 and SrTiO3 onto stainless steel substrates (Huerta-Flores et al., 2017). These materials were prepared to be used as photocatalysts for H2 evolution under UV light. Regarding the XRD results, the authors could prepare phase pure samples for the orthorhombic and monoclinic phases of NaTaO3. For SrTiO3, they prepared a phase-pure cubic phase and tetragonal phase with Sr2TiO4 as a secondary phase. For both materials, the morphology presented was very uncommon; NaTaO3 had a roof-like morphology, whereas SrTiO3 had a cauliflower morphology, as named by the authors and shown in Fig. 2.9.

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Fig. 2.9 (A) Microstructure of orthorhombic NaTaO3 thin films, (B) cross section of the monoclinic NaTaO3 thin films deposited by LCVD, (C) microstructure of the tetragonal SrTiO3 thin films, and (D) cross section of the cubic SrTiO3-tetragonal Sr2 TiO4 thin films. From Huerta-Flores, A. M., Chen, J., Torres-Martı´nez, L. M., Ito, A., Moctezuma, E., & Goto, T. (2017). Laser assisted chemical vapor deposition of nanostructured NaTaO3 and SrTiO3 thin films for efficient photocatalytic hydrogen evolution. Fuel, 197, 174–185. https://doi.org/10.1016/j.fuel.2017.02.016, with the permission of Elsevier.

The images shown in Fig. 2.9 revealed the capacity of the laser-assisted CVD to produce unusual particle morphology and produce well-covered thin films from these particles. Regarding the catalytic performance, the catalyst was the orthorhombic NaTaO3, with an H2 evolution activity equal to 5672 μmol g
1 h
1. This activity outperforms, by at least four times, some other ones presented by NaTaO3 synthesized via sol-gel, solid-state reaction, or solvocombustion methods (Huerta-Flores et al., 2017). CVD is beneficial for thin film deposition, but it can also be used to prepare powders. For example, Xu et al. prepared LaCoO3 powders with diameters around 100 nm, using rotary CVD (Xu et al., 2018). These powders were precipitated and dispersed in α-Al2O3 and were used as catalysts for NO decomposition. The precursors used were La(dpm)3 and Co(dpm)3, where dpm is dipivaloylmethanate. The precursors were vaporized at temperatures around 250°C. The deposition was carried out at 700°C for 1 h. According to XRD, the samples presented some La2O3 and Co3O4 as secondary phases, but the FE-SEM elemental mapping indicated a homogeneous distribution

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of lanthanum, cobalt, and oxygen throughout the samples. Regarding the catalytic tests, the LaCoO3 catalysts presented higher NO conversion and activity per unit than other catalysts used for NO decomposition, such as La2O3, Co3O4, α-Al2O3, La1.2Ba0.8NiO4, and Ba0.8La0.2Mn0.8Mg0.2O3.

2.7 Electron-beam physical vapor deposition (EBPVD) The electron-beam-physical vapor deposition (EB-PVD) technique is versatile in fabricating thin films with controlled microstructure and thick coatings on a suitable substrate (Gong & Wu, 2011; Hkiri et al., 2021). In EB-PVD, a vapor cloud is created when the target material is bombarded with a high-energy electron beam source under a magnetic or electric field. The high temperature generated on the target causes melting of target material resulting in a vapor cloud. This process occurs in a highly evacuated chamber. Subsequently, the vapor cloud condenses on the surface of the substrate resulting in a thick film or thin film in a low-pressure environment. Likewise, since the evaporating rate is high, almost all the materials can be evaporated. During deposition onto the substrate external heat is typically applied. The application of heat (activating the surface of the substrate) improves the bonding between the film and the substrate (Qiu et al., 2021). Five important issues pertain to films deposited by the EB-PVD process. First, the electron beam source can be self-accelerated straight or electromagnetic deflected through 180° or 270°. Second, the evaporated material is placed in a water-cooled copper crucible or tantalum crucibles that withstand high temperature during evaporation. The substrate temperature during the vapor cloud condensation is a third issue. Because during the vapor cloud condensation, two different sets of phenomena can happen. For instance, the vapor cloud can deposit, adhere, and then diffuse into the substrate. Or, the vapor cloud can deposit but debonds from the substrate. In this sense, the substrate temperature is an essential factor for increasing the adhesion between coatings and the substrate and affecting the crystal structure. Experimental study (Clabel, Nazrin, et al., 2021) and a model created (Chevallier et al., 2021) investigated the relationship of substrate temperature and coating crystal structure. Fourth, in the vacuum chamber, if O2 exists in the chamber, it will decrease the background vacuum. This is because O2 can result in filament oxidation of the electron gun, which will affect its lifetime, and avoid impurities present in the vacuum chamber, which can be incorporated in the film. The background vacuum pressure should be evacuated in 10
2–10
6 Pa during deposition. Fifth, deposition rate, both the current and voltage, are important parameters of the deposition rate. Deposition rate affects the size of the columnar crystals and the size and number of holes between the grain boundaries. As the deposition rate increases, the size of the columnar crystals and the holes between the grain boundaries also increase. However, both deposition rate and substrate temperature are relations e.g., low temperatures of the substrate, and at higher deposition rates, more defects are incorporated in the film. Contrarily, in the substrate’s high temperature, the defect is avoided and grain growth is enhanced. Thus, the deposition rates influence the type of zone structure (columnar structure, compacted columnar structure, and recrystallization structure) one obtains and the method of growth of the film (Stutzin et al., 1993). Likewise, multiple electron beam guns and target materials may

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be used to enhance coating adhesion, proper chemistry, and microstructural control. In this sense, various sources are evaporated from several crucibles, one for each element, and the vapor is condensed on a substrate simultaneously. The control of the temperature of the individual sources is fundamental. Since the evaporation rate depends exponentially on temperature, a slight fluctuation can result in a significant variation in the composition of the film. Two considerable advantages arise from the use of the EB-PDV technique: higher energies of evaporating atoms are in the creation of dense films and the ability to not only evaporate elements with low melting points (e.g., Zn, Al, Ag, etc.) but also elements with high melting points (e.g., Mo, W, C, etc.). Highly reactive elements (e.g., Ti, Ta, and Nb) can also be evaporated. The electron beam-physical vapor deposition (EB-PVD) technique is widely used in the fabrication of thermal barrier coatings (TBCs), electrodes, optoelectronic devices, thin dielectric deposition, and storage capacitors, among others (Clabel et al., 2017, 2019; Hasan et al., 2020; Peters et al., 2001; Yang et al., 2020). Depositing ceramic thermal barrier coatings (TBCs) are applied to gas turbine engines, consisting of a metallic bond coat and a ceramic topcoat (Clarke et al., 2012). Strunz et al. intensively investigated the evolution of microstructure in the TBC-type perovskite SrZrO3 and pyrochlore La2Zr2O7, to understand the link between exposure microstructure and macroscopic properties of TBC materials (Vassen et al., 2000). The result reveals that through fabrication by EB-PVD of SrZrO3 and La2Zr2O7, and exposed at 1200°C up to 100 h showed a stable microstructure of large pores. Likewise, a larger volume fraction of pores in SrZrO3 and La2Zr2O7 than Y2O3-ZrO2 (Strunz et al., 2006) was shown, which can be due to the fact that the vacancies do not diffuse and cluster at all, as the oxygen vacancy mobility is reduced in SrZrO3 and in La2Zr2O7 than in Y2O3-ZrO2, due to its ordered structure. Mion et al. (2014) fabri˚ layer of Ti, followed by a 6000 A ˚ Au cated electrodes using EB-PVD with a 250 A layer on a perovskite Ba(Ga,Ta)0.05Ti0.90O3 surface. A temperature insensitive dielectric response, dissipation factor, and tunability over the temperature range (
50°C to 125°C) suggests Ba(Ga,Ta)0.05Ti0.90O3 is a plausible candidate for frequency-agile devices. Perovskite nanostructured SrTi1
xFexO3 thin films with different Fe concentrations (0.075, 0.10, and 0.15 mol% Fe) were prepared by EB-PVD (Mastelaro et al., 2013). SrTi1
xFexO3 samples of 70 nm thickness were evaporated by keeping the substrates at around 50°C, with the oxygen pressure in the chamber kept around 2  10
4 mbar. It was revealed that the surface roughness of the films increased slightly (1.4– 9.2 nm) as the amount of iron increased. The perovskite SrTi1
xFexO3, thin-film sensor response exhibits the best ozone sensitivity to the x ¼ 0.075 sample operating at 250°C. Da Silva et al. (2015) report an investigation into the sensitivity of two oxidizing gases (ozone and nitrogen dioxide) for perovskite SrTi0.85Fe0.15O3 thin films deposited by the EB-PVD process. SrTi0.85Fe0.15O3 was deposited in a Balzers BAK600 evaporator on Si (100) and SiO2/Si substrates containing 120-nm-thick Pt electrodes. Electrical measurements revealed that the thin film was sensitive to oxidizing gases, especially to low ozone gas levels, exhibiting a fast response time, a short recovery time, and good reproducibility and reversibility. Clabel et al. (2015, 2019), Rivera et al. (2014), Khaoula Hkiri et al. (2021) utilized the EB-PVD method

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for fabricating Eu:Y2O3, ZrO2, BaTiO3(BT), CaZrO3 (CZ), and Er-doped BaTiO3 (BTE) films, which were deposited on Si(100) substrates and Si(111). Both BT and BTE films were produced in an evaporation chamber at a pressure of 4  10
6 mbar, using an electron beam gun operating at 5 kV and 50 mA. In comparison, CZ films were produced in an evaporation chamber utilizing an electron beam gun operating at 15 kV, with a beam current of 50 mA in a high vacuum (5  10
6 ˚ /s for the coating of CZ. The condimbar) and deposition rate of approximately 0.6 A tions used resulted in the homogeneous deposition of hard transparent films. Likewise, the BT, BTE, and CZ films have shown roughness values from 2 to 6 nm, 2.27 to 12.9 nm, and 6 to 12 nm, respectively. The experimental results demonstrate that extrinsic and intrinsic defects in BT and BTE can induce and enhance luminescence. Likewise, the luminescence spectrum in thin films of CZ exhibit a blue emission band (425 nm). Such a response can be considered a promising material for photonics applications in both cases. Y. Yang et al. (2020) developed ferroelectric thin perovskite SrTiO3 films using the EB-PVD method. Before deposition, a mechanical pump coupled with a molecular pump was utilized to make the pressure below 2  10
3 ˚ /s at room Pa, and the SrTiO3 target was deposited on the substrate at a rate of 0.6 A temperature. The maximum field emission current density reaches 88.96 μA/cm2 at 10.45 V/μm applied electric field, while field emission results showed that the planar SrTiO3 film emitter device showed good field emission stability and repeatability. Such results make SrTiO3 films a more excellent application prospect in optoelectronic fields such as field emission applications.

2.8 Molecular beam epitaxy (MBE) MBE is used to grow structures and allows precise control of doping levels, atomic interfaces, and the growth of single-crystal thin films (Li et al., 2019). The growth chamber of an MBE system is an ultrahigh vacuum (UHV) environment lower than 10
12 Torr, for epitaxial growth (aligned to the crystal structure of the substrate) via interaction of one or several molecular or atomic beams that occur by impingement of the beams onto a surface of a rotating heated crystalline substrate. The UHV is achieved by the use of cryo-, turbo-molecular, and ion pumps to complete this vacuum. These pumps have different strengths and can achieve different levels of vacuum. The system’s vacuum is further improved by cryopanels encasing the growth chamber. Likewise, such cyopanels also remove the heat dissipated by the heater and the individual sources (Arthur, 2002; Mazet et al., 2015). The UHV does not permit the impurities to impinge on the surface of the substrate, preventing gas-phase collisions of the species on the way to the substrate. Subsequently, samples are loaded into the growth chamber through a series of initial chambers to reduce contaminants. In this sense, a small load lock chamber is brought to the air environment to introduce multiple new samples into the system, afterward pumping down and heating the substrate and holder to reduce contamination. Then, samples may be transferred into a buffer chamber through a gate valve, where they may be stored until growth. This chamber typically contains a heated stage on which individual samples may be heated to a higher temperature to remove contaminants further. The beams are generated by

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high-purity elemental evaporation sources, where the cell temperatures and valves control the source fluxes. The sources are arranged in a ring and aligned to a direct line of sight between the source and the substrate holder. Gas sources can also be used in molecular form or as excited as plasma. To ensure uniformity in how the flux of each material is distributed around a wafer, it is necessary to rotate the sample during growth. Different beams can be turned on and off with less than a monolayer of development by mechanical shutters in front of each source, allowing for multiple layers with abrupt interfaces to be grown. As mentioned before, in MBE growth, working in UHV enables us to monitor the growth procedure. Depending on the needs, a variety of characterization apparatuses can be assembled inside the MBE chamber. Among them are reflection high energy electron diffraction (RHEED) and quadrupole mass spectrometer (QMS) (Brown et al., 2006). RHEED is essential for in situ and real-time surface study on flat surfaces. The main advantage of MBE is the operation of growth under UHV conditions, which gives us the possibility of doing in situ characterizations such as RHEED to monitor the growth of thin films. It allows getting a live picture of the reciprocal space lattice of the sample surface. It is susceptible to surface changes, either by structural changes or due to chemical adsorption (Arthur, 2002). QMS allows determining the type and amount of molecules and therefore is used to determine the residual gas species within the growth chamber. QMS benefits from using a quadrupole radiofrequency electric field (dynamic mass analyzer) that forces ions of a particular mass onto stable oscillatory trajectories. Such a feature present allows us to apply it as a mass spectrometer. Another application of the QMS is adsorbing/desorbing flux measurements (Herman & Sitter, 1996). MBE techniques have most frequently been used for photonics, electronic devices, energy harvesting, electromechanical systems, or sensors. Darby et al. revealed a symmetry variation in BiFeO3 and Bi1
xFexO3 (x ¼ 0.08 to 0.24) films, induced to an enhanced piezoelectric response (Darby et al., 2013). Using the MBE technique, both perovskite films were fabricated on Si/SiO2/TiO2/Pt and silicon substrates. The increase in the piezoelectric response immediately before the onset of a mixed-phase region is attributed to the softening of the lattice in the vicinity of a phase boundary. Chambers et al. used MBE-deposited epitaxial undoped BaSnO3 films on SrTiO3(001) and LaAlO3(001) substrates, and determined the optical bandgaps and band offsets. The perovskite BaSnO3 film has high optical transparency because of its wide bandgap, and exhibits direct and indirect bandgaps of 3.56  0.05 eV and 2.93  0.05 eV. Raghavan et al. (2016) and Prakash et al. (2017) studied perovskite BaSnO3 films fabricated by MBE as a functional material and showed that the BaSnO3 films had high electron mobility. Likewise, Raghavan et al. (2016) demonstrated a modified oxide molecular beam epitaxy approach, which supplies preoxidized SnOx, leading to electron mobilities of 150 cm2 V
1 s
1 in films grown on PrScO3; similar results were obtained for SrTiO3 (Cain et al., 2013). Cheng et al. (2020) fabricated devices of high efficiency, of heterostructure field-effect transistors (HFETs) type BaTiO3/BaSnO3 were MBE grown on SrTiO3 substrates. The device demonstrated a record high current density of 406.7 mA/mm and a maximum transconductance of 72.3 mS/mm. The device has a low threshold voltage of
4.5 V

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Perovskite Ceramics

and can modulate and deplete  5.7  1013 cm
2 electron density. The oxygen evolution reaction (OER) process is a key to emerging energy conversion technologies such as electrochemical energy storage devices and water splitting. In this sense, Tang et al. (2016) report the OER electrocatalytic activity on a SrIrO3(100)p film grown on a DyScO3(110) substrate using MBE. SrIrO3 exhibits more than an order of magnitude activity higher than IrO2 despite having the same nominal valency (Ir4+).

2.9 RF magnetron sputtering Magnetron sputtering is a variant of the widely used physical vapor deposition (PVD) technique applied to deposit thin films. It can be operated in a wide range of discharge modes (Gudmundsson, 2020), such as direct current (dc), pulsed-dc, radio-frequency (rf), and ionized sputter-deposition techniques (Sarakinos et al., 2010) such as highpower impulse magnetron sputtering (HiPIMS) (Rofifah, 2020). The driving force has been aimed at better magnetic, electric, ferroelectric, and photonic properties and friction-corrosion resistance and wears resistance properties, thus improving functionality. Such goals will be achieved from its several advantages: having the same composition as the target source, good adhesion to substrates, high density, and homogeneity of as-deposited thin films. Sputtering is a deposition process based on the ejection of atoms, molecules, or molecular fragments from a target that is bombarded by energetic particles (mostly ions) in a glow discharge plasma located in front of the target. After the process of bombardment on the target, the atoms will be removed from the target’s surface (i.e., emitted particles will be generated), which may then condense on a substrate as a thin film. The magnetic field strength and magnetic field configuration variations strongly affect the glow discharge plasma, deposition rates, and energy flux in a magnetron sputtering system (Ekpe et al., 2009). Magnetron sputtering can be performed under two magnetic field configurations to improve the ion bombardment on the substrate during film deposition; they are conventional or balanced magnetron (BM) and unbalanced magnetron (UBM) (Gudmundsson, 2020). In BM sputtering discharge, all the field lines of the magnetic trap form closed loops between the magnetic poles. Thus, the plasma will be confined to the target, as shown in Fig. 2.10A. On the other hand, in UBM, sputtering can exhibit two basic shapes, UBM-type I and UBM-type II. In case of UBM-type I configuration, the central magnets are stronger than the outer magnet. In that case, some of the field lines are directed to the chamber walls, and the plasma density in the substrate vicinity is low. In case of UBM-type II configuration, the outer magnets are stronger than the central magnet. In that case, not all the field lines form closed loops between the magnetic poles, and some field lines are directed toward the substrate, and some of the electrons are channeled from the plasma toward the substrate, as shown in Fig. 2.10A and B. In this situation, gas ionization occurs near the substrate, and the energetic Ar gas ions (Ar+) will bombard the surface of the substrate, causing ion-assisted deposition of the growing films and changing the surface properties of the deposited films. In a DC diode sputtering system, an interesting behavior on the cathode surface can be observed in the presence of oxygen flow (Sarkar, 2013). We take an example of reactive

Fig. 2.10 Schematic diagram of sputtering processes under (A) balanced magnetron (BM), (B) unbalanced magnetron configurations (UBM) and (C) Reactive sputtering hysteresis curves in deposition rate, chamber pressure, and discharge voltage, all as a function of reactive gas flow during the sputtering of a metal target. Adapted from Gangwar, A. K., Godiwal, R., Jaiswal, J., Baloria, V., Pal, P., Gupta, G., & Singh, P. (2020). Magnetron configurations dependent surface properties of SnO2 thin films deposited by sputtering process. Vacuum, 177, 109353. https://doi.org/10.1016/j.vacuum.2020.109353 and Rossnagel, S. M., & Powell, R. A. (1999). PVD for microelectronics: Sputter deposition applied to semiconductor manufacturing. Academic Press, with the permission of Elsevier.

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sputtering of aluminum in oxygen gas. In a low oxygen flow (see Fig. 2.10C, color orange), slight changes are observed. These are there because oxygen is rapidly adsorbed by the freshly sputtered Al films on the walls and the sample. When the oxygen flow is increased (see Fig. 2.10C, color green), slight changes are observed, the deposition rate is similar, the discharge voltage on the cathode is unchanged, and there is no change in the chamber pressure. On further increase of the oxygen flow (see Fig. 2.10C, color red), an irreversible process (hysteresis) is observed as follows: the deposition rate drops rapidly, and the chamber pressure increases quickly. It also turns out that for many gas-cathode systems, the discharge voltage changes significantly also, due to changes in the secondary electron yield going from a clean to a compound surface (Fig. 2.10). Recently, Hao et al. (2021) fabricated K0.5Na0.5NbO3 films on SrRuO3 (SRO)buffered (100) SrTiO3 (STO) substrates with high dielectric tunability, using the RF-magnetron sputtering method, see Fig. 2.11. The plasma is much more confined in magnetron sputtering, leading to higher plasma densities at the target surface and higher sputtering rates. Sputter-deposited films were grown in a vacuum chamber in a mixed Ar/O2 atmosphere (1.2 Pa, Ar/O2 flow ratio ¼ 3:1). Higher crystallinity in the deposited K0.5Na0.5NbO3 films was also reported through the intensity of the reflection peaks in X-ray diffraction; there were no other secondary phases. The dielectric tunability, defined as the relative change (%) of dielectric permittivity under an E field, positively correlates with the tunability factor (which characterizes the rate of the nonlinear change of its dielectric permittivity) of the ferroelectric film. Thus, after a survey, the configuration of the material design showed a significant η  80% value achievable in (001)-oriented tetragonal K0.5Na0.5NbO3 films. Tao Li et al. (2014) also utilized the RF-magnetron sputtering method for fabricating highly (001) oriented K0.55Na0.55NbO3 films on SrRuO3-buffered SrTiO3 and studied the electrical properties. The preferential growth of K0.55Na0.55NbO3 films was induced by the (001)-oriented SrTiO3 and the thickness of K0.55Na0.55NbO3 film was about 1 μm by changing the Ar/O2 ratio as 9:1 under a total pressure of 1.3 Pa. However, the dielectric tunability is reported the lowest: η  69.7% values for K0.55Na0.55NbO3 thin film dielectrics. Recently, Jong-Hyun Kim et al. (2021) utilized the RF-magnetron sputtering method for fabricating (K1
xNax)NbO3 films to develop piezoelectric energy harvesters (PEHs). The (K1
xNax)NbO3 thin films were grown in a vacuum chamber with a mixed atmosphere of Ar and O2 (Ar:O2 ¼ 4:1) under a total pressure of 1.33 Pa. Moreover, to endow the potential application to the piezoelectric energy harvesters (PEHs), they grew (K1
xNax)NbO3 on SN/P-S (SN ¼ Sr2Nb3O10, P-S ¼ Pt/Ti/SiO2/Si), SN/Ni, and SN/I-Q (I-Q ¼ ITO/quartz) substrates at room temperature. The significant d33 value of this highly (001)-oriented crystalline (K1
xNax) NbO3 film on SN/Ni substrate is induced on PEH, showing a significant output power (proportional to the value of d33 2/εr) of 2.9 μW with a power density of 20 μW/mm3 at 1.0 MΩ (Fig. 2.11). The RF-magnetron sputtering method is critical for creating layer-by-layer, bilayer, and multilayer films (Godiwal et al., 2021). Ruyi Zhang et al. (2021) synthesized high-mobility epitaxial La-doped BaSnO3 films on SrTiO3 single crystal substrates using high-pressure magnetron sputtering. Electron mobility can reach

Fig. 2.11 Cross-sectional TEM images of the thin films with the KNN thickness of (A) 150 nm and (E) 500 nm; (B–C) and (F–G) are the high-resolution TEM images near the SRO/STO and KNN/SRO interfaces for the 150- and 500-nm-thick KNN films, respectively; (D) and (H) are the selected area electron diffraction patterns of the boxed and circled regions in (C) and (E), respectively. Adapted from Hao, L., Yang, Y., Huan, Y., Cheng, H., Zhao, Y. Y., Wang, Y., Yan, J., Ren, W., & Ouyang, J. (2021). Achieving a high dielectric tunability in strain-engineered tetragonal K0.5Na0.5NbO3 films. npj Computational Materials, 7(1). https://doi.org/10.1038/s41524-02100528-2, with the permission of Nature.

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121 cm2 V
1 s
1 at the carrier density of  4.0  1020 cm
3 at room temperature. The critical point is in high argon pressure during sputtering, which plays a vital role in stabilizing the films and inducing oxygen vacancies, facilitating high mobility. Tan et al. (2021) fabricated Li0.33La0.55TiO3 with a perovskite structure for further functionalization in flexible all-solid-state batteries, particularly all-solid-state lithium-ion batteries (ASSLB). To improve performance, i.e., their ionic conductivity, layer by layer highly conductive solid-state electrolytes of the type Li0.33La0.55TiO3 were sandwiched between LiNi0.5Co0.3Mn0.2O2 with a volume change rate of 5% as the cathode and Li4Ti5O12 with a volume change rate of 2% as the anode. The sputtering power densities of LiNi0.5Co0.3Mn0.2O2, Li0.33La0.55TiO3, and Li4Ti5O12 were 2.63, 3.07, and 2.63 W cm
2, respectively, and the base pressure of each deposition was less than 5  10
4 Pa. Two types of structures were applied: the 1-stack mono-bipolar-ASSLB cell and the 2-stack bipolar-ASSLB cell. When the length of the battery is compressed to one-third of the original size, the battery’s capacity remains at 89.2% of the unbent state. Thus, the 2-stack bipolar all-solid-state lithium-ion battery cell is an excellent candidate for microwearable energy devices. Sun et al. (2021) developed a ferroelectric thin film capacitor of high-energy storage density using the RF-magnetron sputtering method. Perovskite-structured Smdoped BaZr0.2Ti0.8O3 were successfully fabricated on SrTiO3 substrates with pseudo-cubic phase and strict c-axis epitaxial behavior and fourfold symmetry under a 20-Pa mixed atmosphere of O2:Ar ¼ 1:1 at 800°C and different sputtering times to controlled different thicknesses. Consequently, a better energy storage density of 40.42 J/cm3 with η of 85.03% was obtained in 200 nm film. Although device fabrication using the RF-magnetron sputtering method is one of the most widely used PVD techniques, high-power pulsed magnetron sputtering (ionized PVD technique) results in plasma electron densities as high as 1019 m3 above the target surface during the short high-power pulses compared to conventional systems (Raman et al., 2015). This advantage results in a wide variety of applications like the formation of diffusion barriers, improved adhesion of hard coatings (Raman et al., 2015), optical thin films (Rydosz et al., 2020), electrical applications (Mesˇkinis et al., 2015), solar cells (C. Chen, et al., 2015), depositing wear and corrosion-resistant coatings, and vias in the microelectronics industry (Hrostea et al., 2020).

2.10 Pulsed laser deposition (PLD) Pulsed laser deposition (PLD) is a thin film physical vapor deposition technique that uses intense laser pulses of short duration to ablate materials from the target using a focusing lens, i.e., the interaction between the laser beam and a solid surface. This interaction involves the following steps: (i) absorption of photon energy by the target; (ii) plasma formation (by surface melting of the oxide target), heating, and initial isothermal expansion; (iii) adiabatic expansion; and (iv) nucleation and growth of the thin films on a substrate. It is essential to highlight that evaporation and ionization of the oxide target lead to the formation of a plasma plume in which atoms and molecules are transferred onto a substrate. The PLD process can be done in a high vacuum or in the presence of an ambient gas such as oxygen. The substrate can be heated to

Synthesis of ceramic materials with perovskite structure

61

enhance the growth kinetics of the deposited film. Choosing proper deposition parameters, e.g., mass distribution, ion and atom velocity, and angular distribution of the plume species, enables control of the grain size and stable microstructure and makes it possible to study microstructure property relationships (Sata et al., 2000). It has been demonstrated that ferroelectrics or an oxide conductor (e.g., BaTiO3, BaSnO, or TiO2) can be deposited or epitaxially grown by pulsed laser deposition (Chaluvadi et al., 2021; Jelı´nek et al., 2017; Wang et al., 2015). One of the PLD advantages is the quick deposition of multiple materials (especially oxides) (Orita et al., 2001), providing the stoichiometric material transfer from the target to the film of any desired composition (Groenen et al., 2015) and giving a path to multilayers of different materials by using multiple targets (Blank et al., 2014). Despite these advantages, PLD is still confined to research purposes for three main reasons. First, the deposited area is relatively small. Second, a plasma plume created during ablation is highly directional and induces nonuniform thicknesses over a film. Last, the extended period of ablation of a target sometimes results in formations of particulates in the plume and thus their incorporation in film materials. In addition, several modifications of PLD have been developed for specific deposition requirements. For instance, a supplementary ion or plasma beam may enhance chemical reactivity through the Aurora PLD method, which is characterized by an improved ionization of the ablated particles during transport from the target to the substrate through interaction between the particles’ applied magnetic field (Debnath et al., 2018). Also, a modification of PLD by a synchronized reactive-gas pulse (pulsed reactive crossed-beam laser ablation, PRCLA) was used to deposit thin films. Details of this modified pulsed-laser deposition (PLD) method can be found elsewhere ( J. Chen, et al., 2015). PRCLA is used to obtain oxynitride films, particularly perovskite oxynitrides, whose general formula can be written as ABOxNy (where A ¼ La, Ba, Sr, Ca; B ¼ Ti, Ta, Nb; x + y ¼ 3); these films have gained a lot of attention over the past decade due to their photocatalytic properties using visible light (Lawley et al., 2020). Multitarget carousels facilitate the deposition of film alloys, heterostructures, and new compounds. They are combined with a literal translation of the substrate to obtain spatial variations of the composition on a larger substrate (Christen & Eres, 2008). PLD was recently applied to prepare CsPbBr3 thin films on a substrate of F-doped SnO2 (FTO) glass, for application in perovskite solar cells (Wang et al., 2019). An excimer laser ablated the CsPbBr3 target with a wavelength of 248 nm (the PLD chamber had a base pressure of less than 10
3 Pa), energy of 100 mJ, and a frequency of 5 Hz. The pulse number controlled the thickness of films. By optimizing the thickness of mesoporous TiO2 and CsPbBr3 layers, the device can achieve the highest power conversion efficiency of 6.3%. Many piezoelectric ceramics proposed for implants are based on perovskite oxide ferroelectric barium titanate BaTiO3 (BTO). Jelı´nek et al. (2017) have successfully deposited BTO layers by PLD on TiNb, Pt/TiNb, and Si(100). A high-power KrF excimer laser (COMPexPro™ 205 F, λ ¼ 248 nm, τ ¼ 20 ns) was used as an external source to vaporize the target material. Smirnov et al. (2021) developed a room-temperature pulsed laser deposition process to obtain Zr-doped In2O3 thin films (named transparent conducting oxides) on glass substrates

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with excellent optoelectronic properties. A KrF excimer laser (λ ¼ 248 nm) was used for all experiments with a repetition rate of 20 Hz and a fluence of 1.9 J cm
2. Likewise, deposition of titanium dioxide and bismuth-based perovskite oxide Bi2FeCrO6 thin films on Si(100) and LaAlO3(100) was achieved by a technique that demonstrates the performance and high-quality thin films of the PLD and RF magnetron sputtering combination. The results showed the improved quality of the deposited films and increasing film uniformity and deposition rate when using the hybrid technique (Benetti et al., 2017).

3

Conclusions

Over the years, the scientific community has extensively studied perovskite-type ceramics because of the excellent combination offered by their physicochemical properties (ferroelectric, piezoelectric, optical, magnetic, etc.). Such multifunctional properties arise from the crystalline phase, surface chemistry, microstructure, and morphological aspects. Modifications in these complex features can drastically expand the overall performance of these materials for diverse technological fields. Thus, the production of materials technologies has a crucial role within this context. Based on this strategy, a detailed description of traditional and promising top-down and bottom-up approaches employed in the preparation of perovskites was introduced in this chapter. The fundamental operating principles and some advantages/disadvantages generally identified in each technique were properly reported in the text. The examples found in published studies were focused on the formation and growth of perovskites under the influence of different optimized synthesis parameters (processing time, heat-treatment temperature, pH condition, etc.). Notably, we hope that understanding these synthetic approaches can serve as a general guide to help scientists design other advanced materials and enhance the synthetic routes in future research.

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Antiferrodistortive phase transition in doped strontium titanate ceramics: The role of the perovskite lattice vacancies

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Alexander Tkach Department of Materials and Ceramic Engineering, University of Aveiro, CICECO—Aveiro Institute of Materials, Aveiro, Portugal

1

Introduction

SrTiO3-based compounds have been attracting considerable interest both for a wide range of applications, particularly in tunable electronic devices and from a fundamental point of view (Kleemann et al., 2020; A. Tkach & Vilarinho, 2019). Undoped strontium titanate (SrTiO3—ST), as a member of the perovskite structure material family, has been extensively studied and drawn the attention of the researchers for its distinctive properties, related to both suppression of a ferroelectric state at low temperatures (K. A. M€ uller & Burkard, 1979) and cubic-tetragonal phase transition occurring at Ta
110 K (P. A. Fleury et al., 1968). The ferroelectric phase suppression has been attributed to paraelectric state stabilization due to quantum fluctuations at low temperatures, because of what ST is known as incipient ferroelectric or quantum paraelectric (K. A. M€ uller & Burkard, 1979). Previous studies addressing ST have thus far revealed that whenever electric field (P. A. Fleury et al., 1968; Worlock & Fleury, 1967) or strain (Uwe & Sakudo, 1976) is applied or dopants (V. V. Lemanov, 2002; Porokhonskyy et al., 2004; A. Tkach et al., 2011) are introduced into the ST lattice, large changes are evident in the lattice and hence in physical properties. Dopants in particular may induce a variety of phases at low temperatures, ranging from dipolar glass and relaxor to ferroelectric. This occurs at isovalent substitution for Sr2+ by Ba2+ (V. V. Lemanov, 1999; V. V. Lemanov et al., 1996), Pb2+ (V. uller, 1984), and Mn2+ ions (A. Tkach V. Lemanov, 1999), Ca2+ (Bednorz & M€ et al., 2005a) and in heterovalent substitution for Sr2+ by Bi3+ (Ang & Yu, 2000; Okhay et al., 2019; Porokhonskyy et al., 2004; Z. Yu et al., 1998), Y3+ (A. Tkach, Okhay, et al., 2017; A. Tkach et al., 2015), Dy3+ (A. Tkach et al., 2018), and Gd3+ ions (A. Tkach, Amaral, et al., 2017). However, no polar state but only the lowtemperature dielectric permittivity decrease was reported to be induced by La3+ incorporation on the Sr2+ site (Z. Yu & Ang, 2002) and by Ti4+-site substitutions with Mn4+ (A. Tkach et al., 2004b) or Mg2+ (A. Tkach et al., 2004a).

Perovskite Ceramics. https://doi.org/10.1016/B978-0-323-90586-2.00012-7 Copyright © 2023 Elsevier Ltd. All rights reserved.

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Nevertheless, the effect of quantum fluctuations itself does not fully explain the undoped ST behavior at low temperatures. First-principle calculations and hyperRaman measurements have pointed out that tetragonal distortion, which emerges on entering the antiferrodistortive (or improper ferroelastic) phase transition, also plays an important role in the polar response of SrTiO3 and other perovskites (Schranz et al., 2022; Yamada & Shirane, 1969; Yamanaka et al., 2000). In its turn, the antiferrodistortive phase transition has been understood within the framework of a triply degenerated soft mode at the Brillouin-zone boundary associated with the antiphase rotation of the oxygen octahedra (Shirane & Yamada, 1969). It has been described within the scope of the Landau theory, wherein long-range strain interactions are considered (Hayward & Salje, 1999). Furthermore, Ta can also be drastically affected by isovalent substitution (V. V. Lemanov, 2002) or by heterovalent substitution with associated vacancies (A. Tkach et al., 2011, 2007) that is the subject of this chapter.

2

Phase transition of undoped ST ceramics

As one of the double oxides with a general formula ABO3, strontium titanate, SrTiO3 is a member of the perovskite family. The prototype perovskite structure is cubic, where relatively large A cations (in this case Sr ions) are situated at the cube corners, small B cations (in this case Ti ions) at the body center, and oxygen ions at the face centers, as represented in Fig. 3.1A. The structure can also be viewed as a threedimensional framework of BO6 octahedra, arranged in a simple cubic pattern (see Fig. 3.1B). In the octahedron unit, the B ion is at the center with the oxygen ions at the corners shared by different octahedra and with A cations in the space between (Last, 1957). Thus, the coordination numbers of A and B cations are 12 and 6,

Fig. 3.1 Cubic perovskite structure with A-cation at the origin (A) and with B-cation at the origin (B). No permission required.

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respectively, and the ideal perovskite unit cell contains one formula unit, i.e., one A cation, one B cation, and three oxygen ions. The stability of the perovskite lattice is related to the high degree of ionic bonding and to the proper value of the Goldschmidt tolerance factor t that defines the relation between the ionic radii in the perovskite lattice (Bhalla et al., 2000; Reaney et al., 1994), as r + rO t ¼ pffiffiffi A 2ð r B + r O Þ

(3.1)

where rA and rB stand for the average A- and B-site ionic radius, respectively, and rO ˚ (Shannon, 1976)). Although stands for the ionic radius of the oxygen (equal to 1.40 A recently a new tolerance factor equation was introduced for perovskites (Bartel et al., 2019), the classical one matches very well to the SrTiO3-based materials and is more transparently deduced, being hence used in this chapter. As can be seen from Eq. (3.1), the A ions fluctuate in their sites when t < 1, whereas B ions are closely packed. And vice versa, when t > 1 the A ions are closely packed, while B ions fluctuate. The appro˚ for the Sr2+ ion and 0.605 A ˚ for the priate effective ionic radii for SrTiO3 are: 1.44 A 4+ Ti ion (Shannon, 1976). Therefore, according to Eq. (3.1), the tolerance factor is very close to 1, confirming the very high packing density of SrTiO3 for both A and B ions and making interstitial occupation not likely. At room temperature, strontium titanate has an undistorted cubic perovskite struc˚ ture, which belongs to the Pm-3m space group, with a lattice parameter a ¼ 3.905 A (Mitsui and Westphal, 1961). However, on cooling SrTiO3 undergoes an antiferrodistortive (or improper ferroelastic) phase transition at a temperature Ta ¼ 105–110 K. This close to second-order structural phase transition from a hightemperature cubic to low-temperature tetragonal but centrosymmetric phase (Lytle, 1964; K. A. M€ uller, 1959; Rimai & deMars, 1962) with a doubling of the unit cell (P. A. Fleury et al., 1968; Shirane & Yamada, 1969) is connected to an antiphase tilting of the oxygen octahedra around one of the [100] axes (Unoki & Sakudo, 1967), as presented in Fig. 3.2A. Hence, no polarization is induced in ST below Ta and the transition is a nonferroelectric structural phase transition. The order parameter of such a transition is a rotation angle φ that varies with temperature as shown in Fig. 3.2B for ST slightly doped with Fe for local symmetry analysis using electron paramagnetic resonance (EPR) (K. A. M€ uller et al., 1968). Below the phase transition temperature, the structure is of the space group I4/mcm and has a unit cell of pffiffiffi pffiffiffi 2 a  2 a  2c, where a and c correspond to the tetragonal one-molecule unit (Shirane & Yamada, 1969). Besides EPR (K. A. M€ uller et al., 1968; Rimai & deMars, 1962), the phase transition was studied using a wide range of experimental methods from elastic moduli measurements (Bell & Rupprecht, 1963; V. V. Lemanov, 2002) and inelastic neutron scattering (Cowley et al., 1969; Shirane & Yamada, 1969) to diffraction techniques (Chrosch & Salje, 1998; Kiat & Roisnel, 1996; Lytle, 1964) and specific heat measurements (McCalla et al., 2016; Salje et al., 1998). Among them, such an inelastic light scattering technique as Raman spectroscopy, involving quantitative analysis

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Fig. 3.2 Oxygen octahedra tilting at antiferrodistortive phase transition of SrTiO3 from cubic to tetragonal phase at 105–110 K (A); tetragonal rotation angle φ of oxygen octahedra in SrTiO3 as a function of temperature below Ta (B). (A) Reproduced with permission from Shirane, G., & Yamada, Y. (1969). Lattice-dynamical study of the 110 K phase transition in SrTiO3. Physical Review, 177, 858–863. https://doi.org/ 10.1103/PhysRev.177.858. Copyright (1969) by the American Physical Society. (B) Reproduced with permission from M€uller, K. A., Berlinger, W., & Waldner, F. (1968). Characteristic structural phase transition in perovskite-type compounds. Physical Review Letters, 21, 814–817. https://doi.org/10.1103/PhysRevLett.21.814. Copyright (1968) by the American Physical Society.

of soft phonon modes as a function of temperature, is particularly popular (P. A. Fleury et al., 1968; Petzelt et al., 2001; A. Tkach et al., 2011). The Raman spectra of SrTiO3 single crystals (P. A. Fleury & Worlock, 1968) and ceramics (A. Tkach et al., 2011) at different fixed temperatures are shown in Figs. 3.3 and 3.4, respectively. Second-order features dominate the spectra, particularly at room temperature. With decreasing temperatures, however, the spectra of ceramics present extra Raman-forbidden infrared-active modes that stem from the local loss of the inversion center, apparently associated with the granular micronanostructured nature of the ceramics (Han et al., 2021; Petzelt et al., 2001). Consequently, the three infrared transverse optical modes, TO1, TO2, and TO4, also become Raman active. These modes can be associated with Ti-O-Ti bending, Sr translations against TiO6 octahedra, and Ti-O stretching, respectively (Perry et al., 1964). As opposed to TO2 and TO4 hard modes, TO1 progressively softens as temperature decreases. However, since no transition toward a ferroelectric phase occurs, TO1 is never completely softened. At Ta, a doubling of the lattice unit cell due to frozen oxygen octahedra rotation around the c-axis leads to a folding of the Brillouin zone. Thus, the zone-boundary modes from the R-point, becoming permitted G-point modes, emerge in the Raman spectra as shown in Figs. 3.3A and 3.4A at 448–459, 146–147, and below 50 cm1. The two hard high-frequency R-modes (unresolved doublets) are marked as Eg + B1g, whereas the low-frequency R-modes correspond to the A1g  Eg structural soft-mode

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Fig. 3.3 Raman spectra at 30 K (A) and 300 K (B) as well as temperature dependence of the soft modes in the tetragonal phase (C) of SrTiO3 single crystal. (A, B) Reproduced with permission from Fleury, P. A. & Worlock, J. M. (1968). Electric-fieldinduced Raman scattering in SrTiO3 and KTaO3. Physical Review, 174(2), 613–623. https://doi. org/10.1103/PhysRev.174.613. Copyright (1968) by the American Physical Society. (C) Reproduced with permission from Fleury, P. A., Scott, J. F., & Worlock, J. M. (1968). Soft phonon modes and the 110 K phase transition in SrTiO3. Physical Review Letters, 21(1), 16–19. https://doi.org/10.1103/PhysRevLett.21.16. Copyright (1968) by the American Physical Society.

doublet (P. A. Fleury et al., 1968; Petzelt et al., 2001). These modes have been associated with rotation of the oxygen octahedra (Shirane & Yamada, 1969) and found to follow a temperature dependence of the form ω ¼ A (Ta  T)n, where Ta ¼ 110 K and n ¼ 0.31, for SrTiO3 single crystals as shown in Fig. 3.3B (P. A. Fleury et al., 1968). The activation of Eg + B1g hard doublet modes at 147 and 448 cm1, occurring at Ta ¼ 110 K, as well as the frequency vs temperature dependence of the A1g soft mode, fitted by a power law ωA1g ¼ A (Ta  T)1/3, with A and Ta as fitting parameters (Petzelt et al., 2001), were also reported for SrTiO3 ceramics as shown in Fig. 3.4B. The transition temperature is specified as Ta ¼ 110  3 K with A ¼ 11.2  0.4 cm1 K1/3 in agreement with the Ta value reported for SrTiO3 single crystals (P. A. Fleury et al., 1968), shown in Fig. 3.3B. A similar value of parameter A was obtained previously for another set of ST ceramics (Petzelt et al., 2001). However, in that work, Ta is reported as high as 132 K, which is likely due to surplus calcium contamination during ceramic processing. It means that even slight doping can significantly affect the transition temperature, as will be discussed in the further sections mainly for ceramics, since they have wider possibility for doping than single crystals.

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Fig. 3.4 Raman spectra at selected temperatures (A) as well as temperature dependence of the first-order mode frequencies (B) in SrTiO3 ceramics. From Tkach, A., Correia, T. M., Almeida, A., Agostinho Moreira, J., Chaves, M. R., Okhay, O., Vilarinho, P. M., Gregora, I., & Petzelt, J. (2011). Role of trivalent Sr substituents and Sr vacancies in tetragonal and polar states of SrTiO3. Acta Materialia, 59(14), 5388–5397. https:// doi.org/10.1016/j.actamat.2011.05.011, Copyright (2011), with permission from Elsevier.

3

Effect of isovalent doping on the phase transition

Raman spectroscopy used for characterization of the Sr1 xCaxTiO3 system with x ¼ 0.007 has shown that such a small amount of Ca2+ substituting for Sr2+ is enough to increase the Ta value to 125 K (Bianchi et al., 1994). This and other results on the onset of the antiferrodistortive phase transition in Ca-doped ST combined with those detected by acoustic methods on Ba-, Pb-, and (Mg1/3Nb2/3)-doped ST are shown in Fig. 3.5 as a function of dopant content (V. V. Lemanov et al., 2004). For A-site doping, Ta was found to correlate with the lattice parameter (V. V. Lemanov, 2002). The transition temperature Ta decreases with increasing lattice parameter, when larger Ba2+ and Pb2+ ions substitute for Sr2+, and increases with decreasing lattice parameter, when smaller Ca2+ ions (Shannon, 1976) occupy the Sr site. On the other hand, an opposite dependence is observed in Fig. 3.5 for B-site doped ST composition SrTi1 x(Mg1/3Nb2/3)xO3, where the transition temperature Ta increases together with the lattice parameter (V. V. Lemanov et al., 2004). To check the effect of smaller B-site dopants, a substitution for Ti4+ ions in ST by Mn4+ ones was performed, resulting in a significant lattice parameter decrease (A. Tkach et al., 2005b). About 5% substitution led to the detection of R-modes in the Raman spectra

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83

Fig. 3.5 Antiferrodistortive phase transition temperature Ta plotted vs Ba, Pb, Ca A-site, and (Mg1/3Nb2/3) B-site dopant concentration x in SrTiO3-based solid solutions. Reproduced with permission from Springer Nature: Lemanov, V. V., Smirnova, E. P., & Ukhin, E. V. (2004). Structural phase transition in (1 x)SrTiO3+xSrMg1/3Nb2/3O3 solid solutions. Physics of the Solid State, 46(7), 1323–1326. https://doi.org/10.1134/1.1778459, Copyright (2004).

only around 80 K and below, as shown in Fig. 3.6A. Moreover, Fig. 3.6B shows that a mode at 437 cm1 was observed in infrared (IR) reflectivity spectra to appear on cooling around 80 K as well (A. Tkach et al., 2010). Furthermore, zone axis electron diffraction patterns (ZADPs) obtained for SrTi0.95Mn0.05O3 with the electron beam parallel to a pseudocubic direction at 16, 80, and 120 K are shown in Fig. 3.6C–E, respectively. ½{ooo} superlattice reflections, characteristic to the tetragonal phase (Reaney et al., 1994; Ubic et al., 2015), are more diffuse than those of undoped ST and could be only detected up to
80 K (A. Tkach et al., 2007). Therefore, Ta was concluded to be suppressed by Mn4+ substitution for Ti4+ on the B-site of ST together with the lattice parameter decrease. At the same time, using the fact that Mn is a multivalent element, ceramics with nominal composition Sr0.975Mn0.025TiO3 were prepared and confirmed to have a majority of Mn ions in 2 + state on Sr sites by extended X-ray absorption fine structure analysis (Levin et al., 2010). That resulted as well in the lattice parameter decrease but with lower slope compared to Ti-site Mn-doped ST ceramics (A. Tkach et al., 2005b). However, Fig. 3.7 shows that the R-modes started to become visible in Raman and IR spectra, while ½{ooo} superlattice reflections became detectable in ZADPs of Sr0.975Mn0.025TiO3 ceramics, at temperatures up to 140 K, i.e., higher than that for undoped ST and much higher than that for SrTi0.95Mn0.05O3 ceramics (A. Tkach et al., 2007, 2010). It was concluded, therefore, that Ta is enhanced while the lattice parameter is decreased by Mn2+ substitution for Sr2+ on the A-site, similarly to the case of Sr-site substitution with Ca2+.

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Fig. 3.6 Raman (A) and infrared reflectivity (B) spectra as well as p zone axis electron diffraction patterns (C–E) of SrTi0.95Mn0.05O3 ceramics at selected temperatures. (A, C–E) Reproduced with permission from Tkach, A., Vilarinho, P. M., Kholkin, A. L., Reaney, I. M., Pokorny, J., & Petzelt, J. (2007). Mechanisms of the effect of dopants and P(O2) on the improper ferroelastic phase transition in SrTiO3. Chemistry of Materials, 19(26), 6471– 6477. https://doi.org/10.1021/cm071795c. Copyright 2007 American Chemical Society. (B) From Tkach, A., Vilarinho, P. M., Nuzhnyy, D., & Petzelt, J. (2010). Sr- and Ti-site substitution, lattice dynamics, and octahedral tilt transition relationship in SrTiO3:Mn ceramics. Acta Materialia, 58(2), 577–582. https://doi.org/10.1016/j.actamat.2009.09.036, Copyright (2010), with permission from Elsevier.

However, there is still an apparent contradiction between the transition temperature—lattice parameter correlation for A-site doped ST and that for B-site doped ST systems. It can be resolved by relating the change in Ta not to the lattice parameter but to the perovskite tolerance factor t, introduced in Eq. (3.1). It has already been demonstrated for other perovskite systems that, as t increases, Ta decreases and vice versa ˚ large ˚ large Ba2+ or 1.49 A (Reaney et al., 1994; Woodward, 1997). In the case of 1.6 A 2+ 2+ ˚ large Sr , A-site ionic radius r A increases (Shannon, Pb ions substituting for1.44 A 1976), leading to an increase in t and, therefore, Ta decreases. Similar behavior is observed for SrTi1 xMnxO3, where t increases due to the decrease of B-site ionic radius rB. On the other hand, the tolerance factor in SrTi1 x(Mg1/3Nb2/3)xO3 system ˚ ) by decreases due to an increase in the average rB at substitution for Ti4+ (0.605 A ˚ ) and Nb5+ (0.64 A ˚ ) (Shannon, 1976) ions. That results in an larger Mg2+ (0.72 A increase of Ta similar to the case of Sr1 xCaxTiO3 and Sr1 xMnxTiO3. Thus, all the data points of Ta vs t for isovalently doped ST (Bianchi et al., 1994; De Lima et al., 2015; V. V. Lemanov et al., 2004; Ranjan et al., 2000; Smirnova et al., 2005; A.

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85

Fig. 3.7 Raman (A) and infrared reflectivity (B) spectra as well as p zone axis electron diffraction patterns (C–E) of Sr0.975Mn0.025TiO3 ceramics at selected temperatures. (A, C–E) Reproduced with permission from Tkach, A., Vilarinho, P. M., Kholkin, A. L., Reaney, I. M., Pokorny, J., & Petzelt, J. (2007). Mechanisms of the effect of dopants and P(O2) on the improper ferroelastic phase transition in SrTiO3. Chemistry of Materials, 19(26), 6471– 6477. https://doi.org/10.1021/cm071795c. Copyright 2007 American Chemical Society. (B) From Tkach, A., Vilarinho, P. M., Nuzhnyy, D., & Petzelt, J. (2010). Sr- and Ti-site substitution, lattice dynamics, and octahedral tilt transition relationship in SrTiO3:Mn ceramics. Acta Materialia, 58(2), 577–582. https://doi.org/10.1016/j.actamat.2009.09.036, Copyright (2010), with permission from Elsevier.

Tkach et al., 2007) appear in the first and third quartiles of Fig. 3.8, presenting opposite variation of the transition temperature with the tolerance factor.

4

Acceptor doping/oxygen vacancy effect

Besides electron-hole generation, substitution of Ti4+ ions with acceptor ions AccTi0 or AccTi00 (e.g., Mg2+ on the Ti site) is known to induce oxygen vacancies VO according to the equations in Kr€ oger-Vink notation: ll

TiTi ¼ Acc0Ti +

1 V 2 O ll

TiTi ¼ Acc00Ti + V O ll

(3.2a) (3.2b)

which are particularly important at temperatures below 750–950 K (Waser, 1991). In addition, VO can be formed when titanates are reduced at low oxygen partial pressure ll

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Fig. 3.8 Antiferrodistortive phase transition temperature Ta for ST-based materials with Sr2+ substituted by Ba2+, Pb2+, Ca2+, and Mn2+, and Ti4+ substituted 5+ 4+ by Mg2+ as a 1/3Nb2/3, and Mn function of tolerance factor t. Several data points are included for some substituents since t varies with dopant concentration. No permission required.

P(O2). Their formation was reported to decrease the Ta value as was found by neutron diffraction (Hastings et al., 1978) and ultrasonic velocity measurements (B€auerle & Rehwald, 1978) in SrTiO3 fired at low P(O2). In order to check the effect of oxygen vacancies, experiments were also carried out on SrTi0.95Mn0.05O3δ ceramics sintered in O2 and N2 atmospheres. Since Mn is a multivalent element, sintering in O2 promotes Mn4+, whereas in N2, Mn2+, and Mn3+ form (Tkach, Vilarinho, & Kholkin, 2006), typically compensated by VO according to Eq. (3.2). Fig. 3.9 shows p ZADPs taken at different temperatures on the grains of SrTi0.95Mn0.05O3δ ceramics, sintered in O2 (Fig. 3.9A–C) and N2 (Fig. 3.9D and E). In addition to the basic reflections of the cubic perovskite, ½{ooo} superlattice reflections, whose schematic representation is shown in Fig. 3.9F, appear for O2-sintered SrTi0.95Mn0.05O3δ at
80 K, at rather the same temperature as for identical ceramics sintered in air. This confirms that Mn4+ is present not only in O2 sintered but also as the major dopant ion in air sintered SrTi0.95Mn0.05O3δ ceramics. On the other hand, superlattice reflections were not observed at 80 K and were only present at 16 K for ceramics sintered in N2, implying that Ta is suppressed when oxygen vacancies are formed. The basis of tilting in perovskites is a cooperative long-range rotation of octahedra in a ‘cogwheel’ manner. Disruption of the ll

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87

Fig. 3.9 p zone axis electron diffraction patterns obtained at 16 K (A, D), 80 K (B, E), and 120 K (C) from the grains of SrTi0.95Mn0.05O3 ceramics sintered in oxygen (A, B, C) and nitrogen (D, E) atmospheres as well as schematic representation of the reflections in p ZAPDs in I4/mcm phase (F). Reproduced with permission from Tkach, A., Vilarinho, P. M., Kholkin, A. L., Reaney, I. M., Pokorny, J., & Petzelt, J. (2007). Mechanisms of the effect of dopants and P(O2) on the improper ferroelastic phase transition in SrTiO3. Chemistry of Materials, 19(26), 6471–6477. https://doi. org/10.1021/cm071795c. Copyright 2007 American Chemical Society.

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connectivity within the octahedral layers by introducing VO statistically decreases the correlation length of tilting/rotation. In order to compensate for this disruption, a greater driving force is required to induce a tilted structure and, therefore, Ta is suppressed. To further verify the effect of VO on Ta, Mg ions were substituted onto the Ti site. Mg ions have a stable 2 + valence state and charge compensation occurs via VO, according to Eq. (3.2b). Fig. 3.10A and B shows the < 110>p zone axis electron diffraction patterns for SrTi0.95Mg0.05O3δ, recorded at 16 and 80 K. No superlattice reflections are observed at 80 K, but weak spots appear at ½{ooo} positions at
16 K, confirming that B-site acceptor doping in ST suppresses Ta. In the Raman spectra for SrTi0.95Mg0.05O3δ (Fig. 3.10C), the R-modes are detected only below 30 K, consistent with the in situ electron diffraction data, for which superlattice reflections were observed at 16 K. It is concluded therefore that Mg2+ substitution for Ti4+, strongly suppresses Ta due to the formation of VO. ll

ll

ll

ll

Fig. 3.10 p zone axis electron diffraction patterns (A, B) as well as Raman spectra (C) of SrTi0.95Mg0.05O3δ ceramics at selected temperatures. Reproduced with permission from Tkach, A., Vilarinho, P. M., Kholkin, A. L., Reaney, I. M., Pokorny, J., & Petzelt, J. (2007). Mechanisms of the effect of dopants and P(O2) on the improper ferroelastic phase transition in SrTiO3. Chemistry of Materials, 19(26), 6471–6477. https://doi. org/10.1021/cm071795c. Copyright 2007 American Chemical Society.

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Fig. 3.11 Antiferrodistortive phase transition temperature Ta for ST-based materials with Sr2+ substituted by Ba2+, Pb2+, Ca2+, and Mn2+, and Ti4+ 5+ 4+ substituted by Mg2+ 1/3Nb2/3, Mn , 2+ and Mg as a function of tolerance factor t. Several data points are included for some substituents since t varies with dopant concentration. No Permission Required.

Concerning the tolerance factor contribution, there is an apparent decrease in ˚ large Mg2+ ions with coordination number of 6 are assumed to substitute t when 0.72 A ˚ for 0.605 A large Ti4+ (Shannon, 1976) as seen in Fig. 3.11. However, one should note that compensation of each Mg2+ at the Ti4+ site requires one oxygen vacancy VO according to Eq. (3.2b), lowering the number of coordinating O2 octahedra ions ˚ large Mg2+ with coorfor Mg2+ in ST lattice to 5. Following such an approach, 0.66 A dination number of 5 (Shannon, 1976) instead of 6 has to be considered for the t-value calculations. Then, the tolerance factor becomes higher, as shown by the dash arrow in Fig. 3.11. Moreover, the O vacancy results in the coordination number of 5 for the Ti4+ ˚ (Shannon, ion adjacent to the Mg2+ one, also decreasing its ionic radius to 0.51 A 1976) and thus further increasing the tolerance factor, as also shown in Fig. 3.11 by solid arrow. As a result, the data point appears in quadrant I of Fig. 3.11 and the rule that a decrease in Ta corresponds to an increase in the tolerance factor can be applied as well to Mg-doped ST. However, the point is still below the general trend due to the VO disrupted correlation between the O octahedra tilting. In addition, McCalla et al. have suggested lowering the number of coordinating O ions for Mg in ST lattice to 4 (McCalla et al., 2016). However, then the Mg2+ ions with a radius ll

ll

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Perovskite Ceramics

˚ (Shannon, 1976) should be smaller than Ti4+ ions that contradicts the of 0.57 A experimentally observed linear lattice parameter increase with Mg content in SrTi1 xMgxO3δ system (A. Tkach et al., 2004a).

5

Donor doping/strontium vacancy effect

In contrast to acceptor doping, donor dopants DSr and DTi should generate electrons. Moreover, low oxygen pressure at preparation stage supports the reduction of Ti4+ into Ti3+ triggering the electronic charge compensation of donor doping in ST, whereas that is not the case for ceramics sintered in air or oxygen (Fu et al., 2008; Kovalevsky et al., 2017; Lu et al., 2016). In the latter case, the donor charge can be compensated ionically in two possible ways for each site: l

l

Sr Sr ¼ DSr +

1 ⁗ 1 V , Ti ¼ DTi + V ⁗ 4 Ti Ti 4 Ti

(3.3a)

Sr Sr ¼ DSr +

1 00 1 V , Ti ¼ DTi + V 00Sr 2 Sr Ti 2

(3.3b)

l

l

l

l

00 The compensation by titanium vacancies V⁗ Ti and strontium vacancies V Sr does not induce V O and will only affect the onset of tilting by considering how cation vacancies influence t. In ST, compensation for donor doping on the A-site is generally considered to occur via V 00Sr (Moos et al., 1997). The Raman spectra of Sr1–1.5xLaxTiO3 and Sr1–1.5xBixTiO3ceramics with different dopant contents recorded at 50 K are shown in Fig. 3.12A and B, respectively. Intensity of the R-mode highlighted in Fig. 3.12 increases with La content (A. Tkach et al., 2011) but decreases with Bi content increase (Ang & Yu, 2000). Therefore, La doping increases Ta, while Bi doping suppresses its value. Once again, such a variation can be ˚ at a coordination explained by the fact that La3+ ions with an ionic radius of 1.36 A 2+ number of 12 are smaller than Sr ions (Shannon, 1976). Moreover, from Fig. 3.13 for Sr1–1.5xGdxTiO3 and Sr1–1.5xYxTiO3 ceramics with the same dopant content x ¼ 0.01, it is apparent that Gd or Y doping also increases Ta to 177 or 203 K, respectively. The increase is larger with a higher difference between the Sr2+ host ˚ and the dopant ion size, extrapolated to be 1.28 A ˚ for Gd3+ and ion size of 1.44 A 3+ ˚ 1.25 A for Y for coordination number of 12 (Shannon, 1976; A. Tkach et al., 2011). As seen from Fig. 3.14A, the increase of Ta with dopant content can be fitted with a linear relation within the corresponding range, wherein the slope is different for each dopant type. The average raising rate of Ta is about 16 K per 1% La, 67 K per 1% Gd, and 93 K per 1% Y (A. Tkach et al., 2011). Hence, Ta is definitely dependent on the type and concentration of dopant used to substitute Sr2+ ions in the ST lattice. However, despite no Ta value was reported for Sr1–1.5xBixTiO3 system and hence not shown in Fig. 3.14A, it evidently decreases with bismuth content according to Fig. 3.12B, whereas all the reported Bi3+ radius values are smaller than that of Sr2+ (Shannon, 1976). One of the explanations can be the fact that the Bi3+ size for coordination number of 12 is not reported by Shannon, but linear extrapolation of the data for lower ll

a -c d e

Raman Intensity (arb. units)

Sr1-1.5xLaxTiO3

Sr1-1.5xBixTiO3

f

j

i

x = 0.0533

x = 0.0533

a -c

(a)

(b)

e f

d

j

gh i

x = 0.0133

x = 0.0133

ab

50 K

100 200 300 400 500 600 –1

Wavenumber (cm )

d e

gh

50 K

i

j

x=0

x=0 0

c

0

100

200

300

400

500

–1

Wavenumber (cm )

600

Fig. 3.12 Raman spectra of Sr1–1.5xLaxTiO3 with x ¼ 0, 0.0133 and 0.0533 (A) Sr1–1.5xBixTiO3 with x ¼ 0, 0.0133 and 0.04 (B) recorded at 50 K. (B) Reproduced with permission from Ang, C. & Yu, Z. (2000). Phonon-coupled impurity dielectric modes in Sr1–1.5xBixTiO3. Physical Review B, 61(17), 11,363–11,366. https://doi.org/ 10.1103/physrevb.61.11363] Copyright (2000) by the American Physical Society.

Fig. 3.13 Temperature dependence of the first-order Raman mode frequencies for Sr0.985Gd0.01TiO3 (A) and Sr0.985Y0.01TiO3 (B) ceramics. From Tkach, A., Correia, T. M., Almeida, A., Agostinho Moreira, J., Chaves, M. R., Okhay, O., Vilarinho, P. M., Gregora, I., & Petzelt, J. (2011). Role of trivalent Sr substituents and Sr vacancies in tetragonal and polar states of SrTiO3. Acta Materialia, 59(14), 5388–5397. https:// doi.org/10.1016/j.actamat.2011.05.011, Copyright (2011), with permission from Elsevier.

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Fig. 3.14 Antiferrodistortive phase transition temperature of Y-, Gd-, and La-doped ST ceramics vs dopant content x (A) and as a function of tolerance factor t (B) with corresponding linear fittings. Modified from Tkach, A., Correia, T. M., Almeida, A., Agostinho Moreira, J., Chaves, M. R., Okhay, O., Vilarinho, P. M., Gregora, I., & Petzelt, J. (2011). Role of trivalent Sr substituents and Sr vacancies in tetragonal and polar states of SrTiO3. Acta Materialia, 59(14), 5388–5397. https://doi.org/10.1016/j.actamat.2011.05.011, Copyright (2011), with permission from Elsevier.

coordination numbers (Shannon, 1976) results in the 12-coordinated ionic radius ˚ (A. Tkach & Okhay, 2021). Therefore, Bi3+ ions should be considered value of 1.45 A as those of similar ionic size to Sr2+. However, then R-mode intensity and Ta value should not vary significantly with Bi content, but they do as shown in Fig. 3.12B. Moreover, even the lattice parameter shows a significant increase with the Bi content (Z. Yu et al., 1998) in contrast to all other ST-based systems with trivalent dopants on Sr site (A. Tkach, Amaral, et al., 2017; A. Tkach et al., 2018; A. Tkach, Okhay, et al., 2017; A. Tkach et al., 2015). Thus, there should be another explanation for that. Then, it is important to remind that, along with the substitution of Sr2+ by trivalent ions, strontium vacancies are formed in order to maintain charge neutrality. Whenever substitution occurs, three Sr2+ ions are randomly removed and then replaced by two trivalent ions and one strontium vacancy. In addition, these vacancies were reported to be a key factor for the record thermoelectric figure of merit obtained in La-doped ST ceramics and demonstrated to order at high enough concentration using highresolution transmission electron microscopy (Lu et al., 2016). Thus, to satisfy charge neutrality in Sr1–1.5xMxTiO3 ceramics, strontium vacancies VSr have to be created along with substitution of divalent Sr2+ by trivalent M3+ ions and the “extended” chemical formula may be written as [Sr1–1.5x(VSr)0.5xMx]TiO3. Therefore, while in ˚ and rB ¼ rTi ¼ 0.605 A ˚ (Shannon, 1976), rA term should be Eq. (3.1) rO ¼ 1.4 A

Antiferrodistortive phase transition

93

Fig. 3.15 Antiferrodistortive phase transition temperature Ta for ST-based materials with Sr2+ substituted by Ba2+, Pb2+, Ca2+, Mn2+, La3+, Gd3+, and Y3+ as well as Ti4+ substituted by Mg2+ 1/ 5+ 5+ 4+ 2+ 3Nb2/3, Nb , Mn , and Mg as a function of tolerance factor t. Several data points are included for some substituents since t varies with dopant concentration. No permission required.

weighted by the concentrations of Sr ions, dopant ions, and Sr vacancies. Fig. 3.14B presents the transition temperature Ta as a function of tolerance factor t for Sr1–1.5xMxTiO3. The tolerance factor was calculated using reported ionic radius values or values linearly extrapolated from the reported ones (Shannon, 1976). Since the actual size of the strontium vacancy is not known, its value was estimated, assuming that a simple linear relation is able to describe the t dependence of Ta, regardless of the type of dopant ion. The Sr vacancy size within the rA term of tolerance factor ˚ . As shown in Fig. 3.14B, all the data equation was incremented with a step of 0.001 A points for three dopants under study were found to follow the linear equation Ta ¼ 176,894–176,501 t with the maximum R2 ¼ 0.9938, when strontium ˚ was applied. It is higher than that of the Sr2+ ion, but such vacancy size of 1.55 A increase can be explained by electrostatic repulsion of the neighboring oxygen anions in the absence of the Sr2+ cation. Therefore, the obtained value of the strontium vacancy size yields an increase in the tolerance factor and hence decrease in the transition temperature Ta for Sr1–1.5xBixTiO3 ceramics. Moreover, our finding that the Sr vacancy is larger than Sr2+ ion is in agreement with the analysis of perovskite systems by Ubic et al.

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(2015), although McCalla et al., suggesting more complicated and less available bond-valence-based quantitative empirical approach for Ta variation not supporting the compositions with stoichiometric vacancies, rejected to consider Sr vacancies (McCalla et al., 2016). On the other hand, they also reported Ta values for SrTi1 xNbxO3 single crystals (McCalla et al., 2016), following the trend previously reported for the same system by Tao et al. (2016). As seen in the summary Fig. 3.15, these points, taken assuming Ti3+ compensation for Nb5+ by analogy to the (Mg,Nb)-doped ST system (V. V. Lemanov et al., 2004), appear as well in the third quadrant of the Ta—t plot.

6

Conclusions

The perovskite-structured SrTiO3 undergoes a cubic to tetragonal phase transition at Ta
108 K associated with rotations of the O octahedra in antiphase around the [001] direction. Using inelastic light scattering as well as electron diffraction to study the lattice dynamics of SrTiO3 with isovalent and heterovalent dopants, substituting for Sr2+ or Ti4+ ions, three main conclusions can be made. First, even for a small percentage of isovalent dopant ion substitution for host ions, Ta is significantly altered, varying oppositely to the tolerance factor t. Second, oxygen vacancies significantly reduce Ta, breaking long-range coupling of rotating octahedra and reducing cation coordination numbers. Third, strontium vacancy size in donor-doped ST is larger than that of Sr ions, thus increasing t and decreasing Ta in ST-based systems doped by donor ions, particularly Bi3+ on Sr site.

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Polaronic effects in perovskite oxides

4

Marius Adrian Husanu and Dana Georgeta Popescu Surface Science, National Institute of Materials Physics, Magurele, Romania

1

Introduction

The discovery of superconductivity and its formulation in terms of electron pairs (Cooper pairs) held together by some sort of attractive field boosted the study of the electron-phonon coupling mechanism. Seen as a possible candidate to mediate the Cooper pair formation, the electron-phonon interaction was the starting point for the elaboration of sophisticated field theories (R.P. Feynman et al., 2010) with dramatic consequences in the future evolution of condensed matter. For historical reasons and consistency of the theoretical concepts, we will introduce first the concept of the large polaron, where the coupling of the excitations extends over several unit cells. Intuitively, polaronic excitations can be understood as conduction charges (electrons or holes) together with the polarization induced in the material by their movement. Originally defined for polar materials and ionic crystals (Landau & Pekar, 1948), the first expression of the renormalized effective mass of the polaron was derived when the coupling involves the longitudinal optical phonons. Its binding energy, the effective (renormalized) mass, and the particular response under external fields depend on the different coupling regimes with the lattice. The interaction between the electrons and the lattice excitations can be expressed by the third term of the standard Fr€ ohlich Hamiltonian (Fr€ohlich, 1954), where the electron-phonon interaction is added at the electron and lattice excitation terms.

 ℌ ¼ ℏ2 k2 =2 m + ⅀q ℏωq dq ✝ dq + 1=2 + + ⅀q Vq dq eiqr + H:c

(4.1)

In Eq. (4.1), an electron with momentum k interacts with a longitudinal optical phonon of wave vector q described in the second quantization form by dq (d✝q), the bosonic annihilation (creation) operators. Vq is the Fourier component of the electron-phonon interaction term, ℏ the reduced Planck’s constant, ω LO the frequency of the phonon, and m is the electron mass. V q ¼ iℏωLO =qð4πα=V Þ½ ðℏ=2mωLO Þ1=4

(4.2)



 α ¼ e2 =ℏc  mc2 =2ℏωLO ½ ð1=ε∞  1=ε0 Þ

(4.3)

Perovskite Ceramics. https://doi.org/10.1016/B978-0-323-90586-2.00007-3 Copyright © 2023 Elsevier Ltd. All rights reserved.

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α is the Fr€ ohlich coupling constant and ε∞,ε0 are, respectively, the electronic and static dielectric constants of the crystal accounting for the response of the system at higher and lower frequencies. We can distinguish two extreme cases, namely: When the electrons adiabatically follow the atomic motion, the Born-Oppenheimer approximation captures well the dynamic of the system with localized electrons. In such a case, the factorization of the total wave function into the “electronic” (φ) and “field” (f) part is a good “ansatz”: |ϕi ¼ |φi  | fi

(4.4)

This case corresponds to the so-called strong coupling limit. Calculation of relevant quantities is straightforward by using Eqs. (4.1), (4.4). For example, the expectation value for the energy is simply hℌi. With the notations from Eq. (4.1), it becomes E ¼   
 ¼ hϕℏ2 k2 =2 mϕi + h f j ⅀q ℏωq dq + dq þ 1=2 
 þ þ ⅀q Vq dq eiqr þH:c jii

(4.5)

yielding the ground state energy of the polaron (Evrard, 1965): E0 ¼ 0:106α2 ℏωLO

(4.6)

And a renormalization of the polaron mass. m∗ ¼ 0:0200α4 mb

(4.7)

These values are not far from the values deduced in more accurate treatments (Miyake, 1975). When α is small, the weak coupling regime allows to treat the interaction term, proportional to α as a perturbation with the polarization or the phonon cloud closely following the slow motion of the polarons. This was shown to result in (Fr€ ohlich, 1954). E0 ¼ αℏω0 ¼ ð1=ε∞  1=ε0 Þe2 ðmω=2ℏÞ1=2

(4.8)

and the effective mass of the electron dressed with the lattice interaction. m∗ ¼ m=ð1  α=6Þ ffi mð1 + α=6Þ

(4.9)

However, the all-coupling continuum polaron theory developed by Richard Feynman, which relies on the path-integral formalism (R.P. Feynman, 1955), delivers results independent of the strength of the coupling.

Polaronic effects in perovskite oxides

101

The main idea of the approach is to reformulate the electron-phonon problem in a Lagrange formalism, where the evolution over time of the system is encoded into the action functional S and the equations of motion are derived by the stationary conditions of S. This is a natural formulation for an electron which polarizes the lattice around and the interaction changes both the energy and the apparent mass of the electron. Furthermore, it travels accompanied by the exponentially decaying polarization which it creates. In the path integral formulation, the transition amplitude between two states is represented by the sum over all histories of exp(iS) with the appropriate boundary conditions for the initial j ii and final j fi states. Within this model, both weak and strong coupling limits are reproduced remarkably, so that in the weak coupling regime. E0 ¼ αℏωLO  0:0123α2 ℏωLO

(4.10)

m∗ =m ¼ 1 + α=6 + 0:025α2

(4.11)

and

while the energy and mass in the SC limit are: E0 ¼ 0:106α2 ℏωLO  2:83ℏωLO

(4.12)

m∗ =m ¼ 0:0204α4

(4.13)

and

Although Feynman used a variational approach starting from an approximate but exactly solvable solution to solve the path integral analytically, the problem can be solved numerically by directly summing the Feynman diagrams for Green’s functions in momentum space. By using the diagrammatic Monte-Carlo technique (A. Mishchenko et al., 2000; Prokof’ev & Svistunov, 1998) it was shown that qualitatively the crossover from weak to strong coupling regimes occurs for α ffi 6 (Hahn et al., 2018; A. Mishchenko et al., 2000; A.S. Mishchenko et al., 2003, 2014), while for α ffi 1 the perturbation theory used to describe the weak coupling regime becomes inadequate. Compared to Eq. (4.1), the Holstein Hamiltonian describing small polarons reads:
 ℌ ¼ t⅀i,j ci ✝ cj + ω⅀i bi ✝ bi + g⅀i ni bi ✝ + bi

(4.14)

where t is the hopping amplitude between neighboring sites, ci(bi) and c†i (b†i) are the fermionic (bosonic) creation and annihilation operators acting on the site i, ω is the phonon frequency, g is the electron-phonon coupling strength, and ni ¼ c†ici is the fermionic particle number operator.

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At this point, it is worth observing that formally the Holstein and Frohlich hamiltonians are not much different, including the same kinetic, the phonon term, and the interaction between the bosonic and fermionic excitation. The distinct features, fundamentally different stems in their different representation: site representation, underline the discrete nature of the Holstein polarons, localized within the lattice constant, while the mode representation derived from a continuum approach illustrates the delocalized nature of the Frohlich polarons. However, despite the simplicity of the Holstein model, there are only very limited cases relying on approximations which allow for an analytical treatment of the polaronic problem.

2

Response of Frohlich and Holstein polarons to external fields

The relevance of the electron-phonon interaction is evident from the changes which it introduces in the electronic structure, which consequently alters the way the system behaves when exposed to an external perturbation. It is understandable, even expected, that the electrons coupled with lattice vibrations manifest differently compared to free electrons. When, for example, the system is under some bias and we measure conductivity or mobility, there will be certain fingerprints pointing at the scattering on lattice vibrations, dependent on the temperature, the character of the external stimulus (continuous or periodic), and the frequency of the incoming perturbation. Similarly, when X-ray or optical experiments are performed, the response of the coupled electron-phonon system, in particular the absorption, will exhibit distinctive features. In the following, we will focus on some experimental techniques and on their theoretical background allowing the detection of polaronic signatures. We will be focusing: (i) on ways to establish the intrinsic electronic properties and how they are influenced by the coupling with the lattice vibrations and (ii) on the vibrational properties. The first category (i) involves understanding the response of the coupled electron-phonon system in optical, electrical, and photoemission measurements, while the second category (ii) involves understanding the fingerprints of the coupled el-ph on the vibrational structure through Raman, IR, and neutron scattering experiments. For each case, we will underline the specific signatures of the polaron size: small vs large polarons and coupling strength going from the weak to the strong coupling limits. This resumes with expressing the response of the system to some external excitation: electric field, thermal, optical, or X-ray radiation.

2.1 Probing the electronic system In a simple picture, we are interested in sketching the general formalism which further allows us to understand the way the coupled electron-lattice vibrations system reacts under an applied field, acting as a perturbation.

Polaronic effects in perovskite oxides

103

An elegant and general formulation is in terms of response function and their associated susceptibilities, which express the proportionality between the applied perturbation and measured response. One assumes first that the change in the expectation value of an operator is linear in the perturbation: ð (4.15) hδOi ðtÞi ¼ dt0 χ ij ðt, t0 Þ ϕj ðt0 Þ where hδOi(t)i is the expectation value of the operator, including its time dependence, ϕj(t0 ) is our considered perturbation, and χ is the response function. It is more practical to use the Fourier transform (Eq. 4.15) and switch to frequency representation in the form: Oi ðωÞ ¼ 0 χ ij ðωÞ ϕj ðωÞ

(4.16)

It is also instructive to split the response function in the real and imaginary parts. χ ðωÞ ¼ Re χ ðωÞ + iIm χ ðωÞ

(4.17)

Then the real and imaginary parts of χ can be written as Re χ ðωÞ ¼ 1

ð +∞

1 2

½χ ðtÞ + χ ðtÞ eiωt dt

(4.18a)

½χ ðtÞ  χ ðtÞ eiωt dt

(4.18b)

∞

and i Im χ ðωÞ ¼  2

ð +∞ ∞

One may notice that the real part of the response function is even, hence invariant over time inversion, Re χ(ω) ¼ Re χ( ω) , while the imaginary part is an odd function Im χ(ω) ¼  Im χ( ω) . The time invariance of Re χ indicates that it is the reactive part while Im χ is the dissipative component involved, for example, in absorption processes, and also named the spectral function. For example, in the case of electric transport, the response function is the conductivity σ tensor and in the linear response theory, the relevant information is its real part. Conversely, the imaginary part is probed in optical absorption and photoemission measurements. The experimentally derived quantity is the spectral function, which in turn can be expressed as a function of the imaginary part of the response function (see Sections 2.1.2 and 2.1.3). Given the analytic nature of the χ(ω) function, the real and imaginary parts of χ are connected through the Kramer-Kronig relation: Re χ ðωÞ ¼

1 P π

ð +∞

Im χ ðω0 Þ dω0 0 ∞ ω  ω

(4.19)

104

Perovskite Ceramics

Im χ ðωÞ ¼ 

1 P π

ð +∞ ∞

Re χ ðω0 Þ dω0 ω0  ω

(4.20)

They derive from the causality condition, equivalent to time ordering in such a way that the evolution in time of the system from t1 to t2 means that t1 < t2. The condition also involves the analytic nature of χ , and that means that the dissipative, imaginary part of the response function is determined in terms of the reactive, real part, and vice versa. However, the relationship is not local in frequency space, and the integration in Eq. (4.18) means that one needs to know Re χ(ω) for all frequencies in order to reconstruct Im χ(ω) for a well-defined frequency.

2.1.1 Electrical and optical measurements: Optical conductivity and light absorption For an electronic system experiencing some external perturbation, the conductivity within the linear response theory reads (Kubo, 1956): σ αβ ðωÞ ¼

ine2 2π + mω hωV

ð∞

dt eiωt

0



j^α ðtÞ, j^β ð0Þ



(4.21)

and the quantity of interest in static conductivity becomes (Degiorgi et al., 1987; Halder et al., 2017; Vashishta et al., 2008) 

βhω

1  e 2πkB T Re σ αβ ðωÞ ¼ π hωV

ð∞

dt eiωt

0

 

j^α ðtÞ, j^β ð0Þ

(4.22)

In Eqs. (4.21), (4.22), n is the carrier density, e is the elementary charge, m is the electron mass, ω is the frequency of the perturbation, and j^α ðtÞ is the current operator. The dc conductivity of the electron gas when ω ! 0 is simply Re σ αβ ¼

1 VkB T

ð∞ dt 0



j^α ðtÞ, j^β ð0Þ



(4.23)

and written in terms of the Fermi-Dirac distribution f using the Boltzmann equation reduces to a simple form σ xx ¼

  e2 X 2 ∂f n ¼ vkx τk  e2 τ V k ∂εk m eff

(4.24)

where conductivity is expressed as a function of velocity v, scattering time τ indicates how long the particle moves until it loses coherence and energy at the Fermi level. This form is very useful as, on the one hand, it makes the connection with the band picture and, on the other, introduces the effective mass, modified by the scattering with impurities, by electron-electron interaction or by coupling with other elementary

Polaronic effects in perovskite oxides

105

excitations. It is hence a natural point to introduce the concept of “quasiparticle,” which allows us to consistently treat the polaronic problem. Its origins are in the formal development of the Fermi liquid theory (Abrikosov et al., 2012; Brown, 1971; Landau, 1957; Venema et al., 2016) where the interaction between particles is absorbed into a single entity at the cost of turning the electron mass into an effective mass and the energy into the self-energy Σ(k,ω). For small polarons and weak coupling, conductivity takes this particular form (D. M. Eagles, 1963; Klinger, 1963; Reik, 1963)  pffiffiffi 2    π t 1  eω=T ðω  4Ea Þ2 pffiffiffiffiffiffiffiffi exp  σ ¼ np e a 16Ea T 2ω Ea T 2

2

(4.25)

while for large polarons the real part of the conductivity can be written as (Alexandrov & Devreese, 2010; J. Devreese et al., 1972) Re σ ðωÞ ¼

πe2 2e2 ω0 α pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi ω  ω0 θðω  ω0 Þ δ ð ωÞ + 2meff 3m ω3

(4.26)

2.1.2 Optical methods: Absorption The absorption coefficient Γ(Ω) of a radiation with frequency Ω was written in its most general form as ( J. Devreese et al., 1972) Γ ð ωÞ ¼ 

jImΣ ðωÞj 1 e2 nεv m ½Ω  Re Σ ðωÞ2 + ½ImΣ ðωÞ2

(4.27)

n is the refractive index of the system, εv is the dielectric permittivity of the vacuum, and Σ(Ω) is the so-called “memory function,” which includes the dynamics of the polaron and depends on Ω, α, temperature, and applied external fields. Expression (4.14) is independent of the coupling strength regime and reflects the fact that under a radiation field, the electron coupled with the induced polarization—the polaron will undergo a transition to an excited state and subsequently relax with the emission of a photon. Expression (4.27) allows one to express mobility for systems of coupled electrons and phonons as well:  ImΣ ðΩÞ Ω Ω!0

μ1 ¼ lim

The physical mechanism of light absorption free polarons can be understood in a twostep model ( J. Devreese et al., 1974; Devreese & Evrard, 1966; Evrard, 1965; Kartheuser et al., 1969). Initially, the electron oscillates in the potential well of the polarization induced by itself and at the ground-state level E0. Under the action of the electromagnetic wave, the electron undergoes first a transition from the ground

106

Perovskite Ceramics

state toward a Frank-Condon (FC)—like excited state characterized by an energy EFC. This sudden transition occurs in time of the order τ  ℏ(EFC  E0) ¼ 1/ω so that the shape of the polarization cloud does not change during the process. However, the FC state is unstable so that the ionic polarization relaxes and adapts itself to the final electronic configuration, and the FC state decays in a second step to an “internal relaxed state” by emission of one or more phonons. The hypothesis of Kartheuser et al. (1969) is that for sufficiently large α (α  3), the (first) relaxed excited state (RES) of a polaron is a relatively stable state, which gives rise to a “resonance” in the polaron optical absorption spectrum. Consequently, a transition leading to a “zero-phonon” peak in the absorption by a strong-coupling polaron was suggested. If the frequency of the incoming photon is equal to ΩRES ¼ 0.065α2ω0, the electron jumps from the ground state (which at large coupling is well characterized by “s”-symmetry for the electron) to an excited state (“2p”), while the lattice polarization in the final state is adapted to the “2p” electronic state of the polaron. Intuitively, the significance of the FC states is evident in the absorption of a photon of appropriate energy, ΩRES when a self-trapped carrier can be freed from the potential well. Since the motions of the atoms that produce the potential well that self-traps the carrier are slower than the photoionization process, the self-trapping potential well is treated as fixed during the photoionization process (Emin, 1993). For photon energies larger than ΩRES + ωLO, a transition of the polaron toward the first scattering state, belonging to the RES, becomes possible. The final state of the optical absorption process then consists of a polaron in its lowest RES plus a free phonon. This is accompanied by a “one-phonon sideband” which appears in the polaron absorption spectrum (De Filippis et al., 2006; Kartheuser et al., 1969): For photon energies larger than ΩRES + ωLO, a transition of the polaron toward the first scattering state, belonging to the RES, becomes possible. The final state of the optical absorption process then consists of a polaron in its lowest RES plus a free phonon. A “one-phonon sideband” then appears in the polaron absorption spectrum. (a) An absorption peak (zero-phonon line) appears, which corresponds to a transition from the ground state to the first RES at ΩRES. (b) For Ω > ΩRES + ω0, a phonon sideband structure arises. This sideband structure peaks around ΩFC. Even when the zero-phonon line becomes weak, and the most oscillator strength is in the LO-phonon sidebands, the zero-phonon line continues to determine the onset of the phonon sideband structure.

In the weak-coupling limit, on the other hand, the absorption coefficient (Eq. 4.15) becomes (Devreese, Huybrechts, & Lemmeks, 1971) Γ ð ωÞ ¼ 

1 2np e2 αω0 2 nεv 3mΩ2

qffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi
 ω=ω  1 θðω  ω0 Þ 0

(4.28)

with n the refractive index of the system, εv the dielectric permittivity of the vacuum, np the carrier concentration. Such features were experimentally identified in oxides such as in manganites (C. Hartinger et al., 2004; C. Hartinger et al., 2006; J.H. Jung et al., 2000; Kim et al., 1998;

Polaronic effects in perovskite oxides

107

Lee et al., 2002; Mercier, 2021; Paraskevopoulos et al., 2000; M.A. Quijada et al., 2001), nickelates (Bi & Eklund, 1993; D.M. Eagles et al., 1995; J.H. Jung et al., 2001), cobaltates (Huang et al., 2019), or titanates (Calvani et al., 1993; D.M. Eagles et al., 1996; Schooley et al., 1964; van Mechelen et al., 2008). The step function in Eq. (4.28) reflects the fact that at zero temperature, the absorption of light accompanied by the emission of a phonon can occur only if the energy of the incident photon is larger than that of a phonon (Ω > ω0) and the absorption spectrum consists of a “one-phonon line,” whereas at nonzero temperature the absorption of a photon can be accompanied not only by emission, but also by absorption of one or more phonons.

2.1.3 X-ray spectroscopies X-ray photoelectron spectroscopy is a photon in - electron out technique where one excites a valence or deep core level electron and measures its kinetic energy. Based on the one-to-one correspondence between binding energy and wave vector with the kinetic energy and emission angle of the outgoing electron, a detailed image of the electronic structure, including correlations and coupling with various excitations, can be achieved. In addition to chemical selectivity, the strength of this technique lies in the ability to directly access the particular changes of the self-energy Σ(k, ω) through the spectral function. The electron dynamics as the result of electron interaction with other quasiparticles can be described by the Green function: GðE, kÞ ¼

1 E  ξk  Σ ðk, ωÞ

(4.29)

related to the spectral function, which is the essential result in angle resolved photoelectron spectroscopy (ARPES) by Aðk, ωÞ ¼

jImΣ ðk, ωÞj ½ω  ξk  Re Σ ðk, ωÞ2 + ½ImΣ ðk, ωÞ2

(4.30)

In Eqs. (4.29), (4.30), ξk is the bare band dispersion and Σ(k, ω) is the complex selfenergy which captures effects going beyond the single particle picture. The real part of the self-energy captures the effect of energy renormalization as the result of interactions and corresponds to the difference between the parabolic bare band dispersion and the renormalized one extracted from the momentum distribution curves maxima. The imaginary part is related to the finite lifetime of the excited states and depends on the correlated nature of the system and/or scattering with impurities. Although a very powerful technique, its limitation is that such experiments have to be carried out in an ultrahigh vacuum, whereas conventional UV radiation sources are only surface sensitive. This underlines the challenging investigation of 3D materials due to a broadening in the kz direction of the reciprocal space.

108

Perovskite Ceramics

Soft X-ray photoelectron spectroscopy (SX-ARPES) is naturally suited to penetrate through the FE capping layer and reveal the electronic properties at the buried interface, while the sharp kz resolution allows to precisely navigate in the k-space, resolve the Fermi surfaces in both kjj and kz directions, and resolve the many-body effects in the E(k) band structure (V.N. Strocov et al., 2012, 2014). For lattice polarons, the spectral function has the following expression: Aðk, ωÞ ¼ Z0 ½δðω  EðkÞ + AH ðk, ωÞ

(4.31)

The first term represents the sharp QP peak with the dispersion E(k) and effective mass m∗ resulting from the renormalization of the noninteracting single-particle band with dispersion ε(k) and mass m0. The second term arises due to phonons coupling to the excited photohole. The effect of the electron-phonon interaction on A(ω,k) is that it reduces Z0 to below unity and builds up, with its dispersion following ε(k), as a FrankCondon series of phonon peak replicas at energy separations ωn ¼ nω0 from the QP peak, where ω0 is the phonon frequency and n indexes the replicas.

3

Electron-phonon interaction (EPI) in transition metal oxides

3.1 EPI in manganites Physics with correlated electrons in transition metal oxides (TMOs) emerges as a promising perspective for transistors (Huang et al., 2019; Mannella et al., 2007; Misewich & Schrott, 2000; Xiang et al., 2011; Zhong et al., 2015) or information storage and processing (Garcia et al., 2009; Yamada et al., 2013; Yin et al., 2013). A promising strategy to develop new functionalities is to work with materials at the verge of a phase transition. This approach is especially suitable to realize giant variations of the conductivity across the metal-to-insulator transition (MIT) in correlated oxides ( Ji et al., 2012; Yajima et al., 2015), with the change driven by structural distortions (Mercy et al., 2017), substrate-induced strain (Yang et al., 2016), thickness (Aetukuri et al., 2013; P.D.C. King, Wei, et al., 2014), bias (Yajima et al., 2015), as well as to switch the spin order between different magnetic ground states (Pantel et al., 2012; Yin et al., 2013). Although the most practical control route is electrical biasing or gating as it offers direct integration in devices, interfacing TMOs is an additional pathway to trigger MIT. For example, electron density in TMO heterostructures can be locally altered by charge transfer or charge reorganization close to the contact region with a dissimilar material (Cao et al., 2016; Luo et al., 2013; Marinova et al., 2015; Salafranca et al., 2014; Vaz et al., 2010; Zhong & Hansmann, 2017). The strength of this approach depends on the sensitivity of the active compound to charge doping. A prototypical system suited for applications in spin-dependent transport (Burton & Tsymbal, 2011; L€ u et al., 2016), or devices with magnetocrystalline anisotropy controlled by electric fields (Chaluvadi et al., 2018; Mannhart & Schlom, 2010; Mathews, 1997), is the hole-doped manganite La1 xSrxMnO3 (LSMO). Due to its mixed Mn3+/Mn4+ valencies and their accompanying double-exchange interaction, it

Polaronic effects in perovskite oxides

109

exhibits a rich phase diagram depending on the doping level (x) including metal-toinsulator transition and colossal magnetoresistance (CMR) effect accompanied by transition between A-type antiferromagnetic (AFM)(x  0.5) and ferromagnetic (FM) states (Fujishiro et al., 1998; Hemberger et al., 2002). Around the ideal doping x  ⅓, La1 xSrxMnO3 is a half metal with fully spin-polarized spin states at the Fermi level having the Mn t2g orbitals completely and one electron in the crystal-field split eg orbitals. In bulk, the eg(x2  y2) and eg(3z2  r2) states, lying in the (001) plane and, respectively, perpendicular to the (001) plane have the same energy, i.e., are degenerate (Tokura, 2000). Strain, in conjunction with breaking of symmetry (Pesquera et al., 2012; Tebano et al., 2008) near the surface/interface, lifts the degeneracy, thus their energy and occupation probability, translating into different magnetic and transport properties. In particular, a compressive strain of 0.3% when using, for example, NdGaO3 as a substrate has been shown to favor the occupation of eg(3z2  r2) orbitals close to the surface (Pesquera et al., 2012; Tebano et al., 2008) or interface with a ferroelectric material (H. Chen et al., 2014; Preziosi et al., 2015), while a tensile strain of 0.5% in the case of SrTiO3 lowers the energy of the eg(x2  y2) states, hence becoming preferentially occupied. Such an approach illustrates the critical role of the coupled or reacted interface to modifying the electronic structure close to the contact region and thus defines new functionality (Bocirnea et al., 2020; Gheorghe et al., 2012; D.G. Popescu et al., 2015, 2019, 2020; Popescu & Husanu, 2014). Remarkably, the same subtle balance between the tunable electron and hole occupations of TMO orbitals with different (x2  y2) and (3z2  r2) symmetries has been shown to open new avenues to engineer exotic quantum states such as metal-toinsulator transitions (Aetukuri et al., 2013), superconductivity (D. Li, Lee, et al., 2019; W.M. Li, Zhao, et al., 2019; Liao et al., 2019) or Dirac-type quasiparticles (Horio et al., 2018; Tao & Tsymbal, 2018). In addition, the interplay of the lattice degree of freedom with the electronic structure through coupled charge-lattice bosonic excitations (Merten, BruchmannBamberg, Damaschke, Samwer, & Moshnyaga, 2019; Quijada et al., 1998), which is a common aspect to many TMOs and to their interfaces, translates into enhanced m∗, stronger carrier localization, and lower electron mobility. Their general signature in ARPES is clear through humps in the spectral function near the quasiparticle peak (QP) and kinks in the experimental band structure at the energy corresponding to the phonon mode (Cancellieri et al., 2016; Verdi et al., 2017; Wang et al., 2016). In two-dimensional, layered manganite La22xSr1+2xMn2O7-2D LSMO, the polaronic state condensate which forms coherent states was clearly revealed in a number of studies (Li et al., 2016; Mannella et al., 2007; Massee et al., 2011; Rønnow et al., 2006) confirming the well-established fact that for Frohlich polarons, the lowtemperature conduction is coherent in nature (Mahan, 1966), whereas at higher temperature the conductivity drops into incoherent mechanisms. The hump peak position in angle resolved photoelectron spectroscopy (ARPES) measurements (Fig. 4.1) is commonly identified with the polaron-binding energy and the identified loss of polaron quantum coherence, signaled as a collapse of the quasiparticle peak at very low energy, and involves considerable transfer of spectral weight from the quasiparticle peak to higher binding energies resulting in the shift of the hump peak position close to the Fermi wave vector kF. Indeed, the temperature range where

Perovskite Ceramics

0.8

0.6

0.6

0.4

0.4

0.0 –0.3

–0.2 –0.1 E – EF (eV)

0.0

0.1

1.0

0.8

0.2

a)

–0.4

TC

1.0 sab/sab30K

106 K 124 K 132 K 149 K 168 K

33 K 48 K 68 K 86 K

0.2

b) 30

0.0 60 90 120 Temperature (K)

QP weight / Qp weight 30K

Intensity (arb. units)

110

150

Fig. 4.1 Polaron coherency. (A) Temperature dependence of the QP peak at kF along the (0,0) to (π,π) “nodal” direction. (B) Integrated QP spectral weight and in-plane DC conductivity σ ab (from Eq. 4.13). Both the QP weight and σ ab have been normalized to their value at 30 K. Open and filled squares denote warm up and cooling down cycles, respectively. From Polaron coherence condensation as the mechanism for colossal magnetoresistance in layered manganites, N. Mannella, et al. Physical Review B 76, 233102 (2007).

the intensity of the quasiparticle peak drops follows remarkably well the variation of the conductivity Fig. 4.1, linking such variation to the colossal magnetoresistance effect. On the other hand, photoemission experiments performed on 3D perovskites such as LaSrMnO3 (Horiba et al., 2016) or CaMnO3 (M.A. Husanu et al., 2020) and focusing on such fingerprints are notoriously challenging due to the narrow probing depths and the inherent kz broadening. Moreover, transport measurements by probing the entire structure can rarely isolate the interface contribution only. Preziosi (Preziosi et al., 2015) and Cui (Cui et al., 2015), exploring LSMOs with different dopings (x ¼ 0.175 respectively x ¼ 0.3), grown on SrTiO3 (STO) and covered by ferroelectric oxides, identified a trend which in both cases is consistent with a change in orbital occupancy induced by the FE layer. More exactly, when the FE polarization is switched between two opposed out-of-plane direction, preferential occupation of either Mn eg (x2–y2) or Mn eg (3z2–r2) states can be achieved. This orbital polarization effect is characteristic to the interface and occurs in accordance with a model elaborated earlier (H. Chen et al., 2014) where the modulation of the Mn-O-Mn bonds close to the interface is triggered by the FE instability propagating several unit cells in LSMO (Peters, Apachitei, Beanland, Alexe, & Sanchez, 2016; Pruneda et al., 2007). However, most of the studies which accessed the interface buried under thin FE layers lacked angle-resolved information on the states at the Fermi energy, focusing on the empty states from an X-ray absorption perspective (Cui et al., 2015; Preziosi et al., 2015) or on deep core levels (Chen & Klein, 2012; D.G. Popescu et al., 2015; Rault et al., 2013) while giving only indirect information on the valence states. k-resolution is critical in identifying the distinct contribution of interface band carriers at the screening of the depolarizing field and compensation of a well-defined orientation of the FE polarization. Consequently, the effect of the FE instability propagating into the joining metallic electrode, lifting the eg orbital degeneracy, and its impact on LSMO carriers with different band character, different effective masses meff, and dimensionality remained unresolved.

Polaronic effects in perovskite oxides

111

Recently, by using soft X-ray photoelectron spectroscopy which enhances the probing depth and the kz resolution, it was shown (M.A. Husanu et al., 2022) how the parabolic band dispersion of both electrons and holes which are featured by kinks indicating renormalized effective mass with respect to their bare surface (Horiba et al., 2016) change upon doping (see Fig. 4.2) of LSMO. For LSMO, which is featured by large polarons and moderate coupling strength (Graziosi et al., 2014; C. Hartinger et al., 2004; C. Hartinger et al., 2006), they identify a more complex scenario with coexisting 3D electron and 2D hole polarons, similar to Ce-doped CaMnO3 (CCMO) (M.A. Husanu et al., 2020).

Fig. 4.2 Experimental electronic structure of Ce-doped CaMnO3. (A,B) Out-of-plane FS maps in the ΓXM plane for 2%Ce-CaMnO3-(CCMO2) (A) and 4%Ce-CaMnO3 (CCMO4) (B) measured with hv at the Mn 2p resonance. The white square designates the cubic-pseudocell surface BZ. The doping clearly increases the experimental Luttinger volume. (C) Out-of-plane FS map for CCMO2 in the ΓZR plane recorded under variation of hv and the trajectory of the 643 eV energy in the k-space (H). (D)–(G) Band dispersions E(k) measured at the Mn 2p resonance for (D,E) CCMO2 and (F,G) CCMO4. E(k) of the 3D bands around the Γ-point (D, F) identifies light electron charge carriers. The blue arrow indicates the threshold energy of EDC maxima deviating in CCMO4 from the parabolic dispersion. ARPES images of the q2D bands around ky ¼ 0.5π/a (E,G) show massive humps extending down in EB which manifest heavy polaronic charge carriers. Also shown through (D–G) are the overlaid DFT-theoretical bands and gradients of the energy-integrated ARPES intensity, identifying kF. Sketch of the 643 eV line cutting the Γ8 point in the second BZ in kjj (H); (I,J) Spectral function A(k,ω) ∝ EDC at kx ¼ kF for the 3D bands (thin lines) and at ky ¼ 0.5π/a for the q2D ones (thick lines) for CCMO2 (I) and CCMO4 (J). For the q2D bands, the whole A(k,ω) is dominated by the polaronic hump. From Electron-polaron dichotomy of charge carriers in perovskite oxides, M.A. Husanu, et al. Communications Physics 3, 52 (2020).

112

Perovskite Ceramics

While for CCMO the additional light electrons introduced by minute Ce doping transition the system across the metal-to-insulator transition by weakening the strength of the EPI (see Fig. 4.2), in LSMO the doping is achieved by interfacing LSMO with a ferroelectric layer which depletes the interface in order to compensate its well-defined ferroelectric polarization.

3.2 EPI in titanates In the cubic phase, stoichiometric SrTiO3 has an empty t2g conduction manifold composed of three dispersive, three-dimensional (3D) bands that are degenerate at the Γ point (Mattheiss, 1972). Its band structure consists of a weakly dispersive (heavymass) band and a pair of degenerate, strongly dispersive (light-mass) bands. They arise, respectively, from the small and large overlaps, along the y axis, of neighboring titanium 3dxz, 3dxy, and 3dyz orbitals. Due to its simplicity, the bulk electronic structure of doped SrTiO3 has been extensively studied, revealing dispersing quasiparticle peaks and in-gap features pointing to different polaronic effects near the Fermi energy (Aiura et al., 2002; Chang et al., 2010; C. Chen et al., 2015; Dudy et al., 2016; Fujimori et al., 1992; Ishida et al., 2008; P.D.C. King, McKeown Walker, et al., 2014; Santander-Syro et al., 2011; Walker et al., 2015; Wang et al., 2016). However, the debate focused on whether the EPI in STO belongs to Frohlich (Cancellieri et al., 2016; J.T. Devreese et al., 2010) or lattice, Holstein-type polaron (Hao et al., 2015) is not entirely resolved and remains still actual. By carefully tuning the carrier concentration by either annealing (C. Chen et al., 2015) (Fig. 4.3) or prolonged exposure to the damaging X-ray radiation (Wang et al., 2016), it was shown that the development of the polaronic replica and its disappearance upon doping can be well captured within a Frohlich model of large polarons with coexisting 2D and 3D dimensionality (see Fig. 4.3). One step closer to pushing forward any device functionality is to extend the range of the active regime from the ideal, carefully prepared surfaces to interfaces. This can be, for example, achieved by exploiting the metallic conductivity which has been shown to emerge at the interface of SrTiO3 with the LaAlO3 insulator (Ohtomo & Hwang, 2004). In a pioneering soft X-ray photoemission study, a large polaron signature featuring the two-dimensional electron gas at the interface has been obtained. By extracting the effective mass renormalization meff/m0  1/Z0, where Z0 is the integral weight of the quasiparticle, the electron-phonon coupling constant ƛ  0.6 has been obtained (Fig. 4.4). Interestingly, the bare SrTiO3 surface at large carrier density does not show any systematic temperature effects in the weight of the quasiparticle peak (Chang et al., 2010), indicating that the interface formation with the LaAlO3 overlayer and the resulting 2DES significantly alter the EPI in the LAO/STO system. Moreover, the exact mechanism of the mobility drop at the LAO/STO interface was identified in coupling to soft phonon modes, different from the hard LO3 vibrations, resulting in a concomitant increase of the electron effective mass meff. This phenomenon provided an insight into the microscopic mechanism of the mobility drop with increasing temperature (A.S. Mishchenko et al., 2015). Also relevant is that the mobility of the two-dimensional gas of electrons at the interface γ-Al2O3/SrTiO3 (GAO/STO) can be enhanced not only by doping but by band

Fig. 4.3 (A, D) Fermi surface of SrTiO3 in kx  ky  kz space with dxy orbital character. The FSs are centred at one Brillouin zone centre, Γ, which is ˚ –1. (B) Evolution of dispersions one high-symmetric point in three-dimensional momentum space with (kx, ky, kz) ¼ 2π/a(1,1, 3)  (1.61, 1.61, 4.83) A from the Γ11 cut. (C) EDC curves (black empty circles) extracted from the momentum window shown in grey in B. Franck–Condon line shape fitted curves are shown with red lines. Blue lines represent individual Gaussian peaks. From Observation of a two-dimensional liquid of Fr€ohlich polarons at the bare SrTiO3 surface, C. Cehn, et al. Nature Communications 6, 8585 (2015).

Fig. 4.4 Polaronic effects at the LaAlO3/SrTiO3 interface. (A,B) High-resolution ARPES images along the ΓX (ky ¼ 0) line at the L3- and L2-edges, respectively, with the superimposed theoretical dxy (pink) and dyz (green) bands. The lower panels show the corresponding second derivative d2I/dE2 > 0 plots, which clearly show both the quasiparticle (QP) peak and the dispersive hump formed by the LO3 phonon. (C,D) A series of (normalized) EDCs extracted from A,B, respectively, at the indicated kx-values through the occupied part of the BZ. The color bar indicates the intensity maximum and minimum. The two curves at the bottom show EDCs integrated over the k-ranges indicated in C,D as well as the whole BZ in the kx-direction. The characteristic PDH spectral structure in C,D manifests a polaronic metal state formed by the hard LO3 phonon, and renormalizing the dyz-band dispersion in A,B and the clear hump dispersion in B identifies a large polaron. From Polaronic metal state at the LaAlO3/SrTiO3 interface, C. Cancellieri, et al. Nature Communications 7, 10286 (2016).

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engineering as well. This goal was achieved by interfacing STO with Al2O3 in spinel form. The interface is altered by the crystalline structure of the top layer, and the energy of the orbitals changes with respect to the bare surface or the LAO/STO interface. Such orbital-dependent additional modulation impacts the electron-phonon interaction as well. The weakening of the EPI, expressed by the Z 0 values in Fig. 4.5, is another aspect of GAO/STO contributing to the μe-boost. For many systems, such weakening is

Intensity (arb.u.)

(A)

(B)

Z0

LAO/STO

ΔE LAO/STO

31%

low-ns GAO/STO

high-ns GAO/STO

–1.0

–0.5 EB (eV)

LO3

EF

62 meV

46%

low-ns GAO/STO

50 meV

44%

high-ns GAO/STO

48 meV

–0.25

EF EB (eV)

dxy dxz

(C)

EF

GAO

dyz

dxy Vos

hv dxz/yz

STO Sr

Ti

Al

O vacancy O

Fig. 4.5 Electronic structure and polaronic effects at the γ-Al2O3/SrTiO3 (GAO/STO) interface. Experimental A(ω,k) for oxygen-deficient LAO/STO and two GAO/STO samples represented by EDCs of the ARPES intensity, integrated over (A) the whole kF interval and ˚ 1 interval. Compared to LAO/STO, in both GAO/STO samples, the QP peak (B) the kF  0.1 A shows larger spectral weight Z0 and smaller energy broadening ΔE (values on the plots), which identify weaker EPI and smaller effective disorder, respectively, (C) represents the sketch of the modified band ordering at the GAO/STO interface. Adapted from Band-order anomaly at the γ-Al2O3/SrTiO3 interface drives the electronmobility boost, A. Chikina, et al. ACS Nano 2021, 15(3), 4347–4356.

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caused by screening of the EPI by mobile electrons. However, at the GAO/STO interface, ARPES shows that, at least near the interface, the dxz/dyz-derived electron density in GAO/STO is smaller than the dxz/dyz (and all the more dxy + dxz/dyz) density in LAO/ STO. In fact, the EPI may be amplified by trapping (localization) of charge carriers on defects (A.S. Mishchenko et al., 2009), and this will lower the crystal symmetry, allowing otherwise prohibited phonon modes and thus additional EPI channels. Such localization-enhanced EPI in LAO/STO is evident from the phonon-mode continuum displayed by A(ω), as opposed to the single mode in GAO/STO. Then the reduction of the effective defect concentration in GAO/STO, expressed by the smaller disorder seen in the QP-peak width, Fig. 4.5, explains the reduction of the EPI. The observed weakening of EPI in GAO/STO propagates into μe not only by polaronic renormalization of the electron effective mass but also by mediating electron scattering on defects, which can somehow be viewed as frozen phonons. Therefore, the defects affect μe not only through their concentration but also through the EPI-mediated electron-defect scattering strength, which in turn increases with the defect concentration.

4

Outlook and perspectives

The impact of the electron-phonon interaction and of the resulting polaronic coupling remains a relevant topic as new limits on material growth and device fabrication are being achieved. Continuous development of advanced deposition techniques such as molecular beam epitaxy, pulsed laser deposition, or atomic layer deposition allows control of unit-cell growth of oxides and device fabrication at the nanometer scale. Such scales may even exceed the size of large Frohlich polarons. However, systems featured by discrete, lattice polarons will still experience the effect of energy and mass renormalization due to EPI. Hence, devices with reduced dimensionality should be the ideal playground to study interacting Holstein polarons, bipolarons, and magnetic polarons while directly connecting the experiments with the theoretical predictions without relying on simplifying approximations. Secondly, further effort should be directed into understanding and separating the contribution of large and small quasiparticles in systems with coexisting Frohlich and Holstein polarons. This is particularly relevant in materials where conduction involves different types of carriers: electrons and holes in light and heavy bands crossing the Fermi energy, as for example in the case of manganites (M.A. Husanu et al., 2020, 2022). Thirdly, the extension of these concepts to other transition metal oxides such as nickelates and vanadates, where EPI was recently observed (Mirjolet et al., 2021), is a direction which is expected to offer new directions for polaron-based functionality controlled by strain, electrostatic doping, gate doping, alloying, or band engineering.

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Acknowledgments This work was funded by the Romanian UEFISCDI Agency under Contracts No. PCE 96/2021, code PN-III-P4-ID-PCE-2020-2540. D.G.P. was supported by the Romanian UEFISCDI Agency under Contract No. TE 50/2022.

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Fujishiro, H., Fukase, T., & Ikebe, M. (1998). Charge ordering and sound velocity anomaly in La1-XSrXMnO3(X 0.5). Journal of the Physical Society of Japan. https://doi.org/ 10.1143/jpsj.67.2582. Garcia, V., Fusil, S., Bouzehouane, K., Enouz-Vedrenne, S., Mathur, N. D., Barthelemy, A., et al. (2009). Giant tunnel electroresistance for non-destructive readout of ferroelectric states. Nature, 460, 81–84. Gheorghe, N. G., Husanu, M. A., Lungu, G. A., Costescu, R. M., Macovei, D., Popescu, D. G., et al. (2012). Reactivity, magnetism and local atomic structure in ferromagnetic fe layers deposited on SI (001). Digest Journal of Nanomaterials & Biostructures (DJNB), 7. Graziosi, P., Gambardella, A., Prezioso, M., Riminucci, A., Bergenti, I., Homonnay, N., et al. (2014). Polaron framework to account for transport properties in metallic epitaxial manganite films. Physical Review B: Condensed Matter, 89, 214411. Hahn, T., Klimin, S., Tempere, J., Devreese, J. T., & Franchini, C. (2018). Diagrammatic Monte Carlo study of Fr€ohlich polaron dispersion in two and three dimensions. Physical Review B: Condensed Matter, 97. https://doi.org/10.1103/physrevb.97.134305. Halder, S., Sheikh, M. S., Ghosh, B., & Sinha, T. P. (2017). Electronic structure and electrical conduction by polaron hopping mechanism in A2LuTaO6 (A ¼ Ba, Sr, Ca) double perovskite oxides. Ceramics International, 43, 11097–11108. Hao, X., Wang, Z., Schmid, M., Diebold, U., & Franchini, C. (2015). Coexistence of trapped and free excess electrons inSrTiO3. Physical Review B. https://doi.org/10.1103/physrevb.91.085204. Hartinger, C., Mayr, F., Deisenhofer, J., Loidl, A., & Kopp, T. (2004). Large and small polaron excitations in La 2/3 (S r/C a) 1/3 MnO 3 films. Physical Review B: Condensed Matter and Materials Physics, 69, 100403. Hartinger, C., Mayr, F., Loidl, A., & Kopp, T. (2006). Polaronic excitations in colossal magnetoresistance manganite films. Physical Review B: Condensed Matter, 73, 024408. Hemberger, J., Krimmel, A., Kurz, T., von Nidda, H.-A. K., Ivanov, V. Y., Mukhin, A. A., et al. (2002). Structural, magnetic, and electrical properties of single-crystallineLa1  xSrxMnO3(0.4 Fferroel

P(+)

P(+)

P(–)

P(+)

Fig. 5.3 An overview of the cases investigated in terms of work-function differences between the substrate and the ferroelectric film, and also as a function of the conduction properties of the substrate. The cases with yellow background and red text are from this work and previous work, as follows: PZT (Ф ¼ 5.3 eV)/SRO (Ф ¼ 4.9 eV) (Maksymovych et al., 2009; Tanase et al., 2016); PZT (Ф ¼ 5.3 eV)/LSMO (Ф ¼ 5.2 eV), inward polarization (Maksymovych et al., 2009; Teodorescu et al., 2017; C.-L. Wu et al., 2011); PbTiO3/DyScO3 with a similar work function, resulting polarization modulated (Catalan et al., 2006); BTO (Ф ¼ 4.8 eV)/SRO (Ф ¼ 4.9 eV), outward polarization (Shin et al., 2008); PZT (Ф ¼ 5.3 eV)/Pt (Ф ¼ 5.65 eV), outward polarization (Apostol et al., 2016). From Tanase, L. C., Abramiuc, L. E., Popescu, D. G., Trandafir, A. M., Apostol, N. G., Bucur, I. C., Hrib, L., Pintilie, L., Pasuk, I., Trupina, L., & Teodorescu, C. M. (2018). Polarization orientation in lead zirconate titanate (001) thin films driven by the interface with the substrate. Physical Review Applied, 10, 034020.

These scenarios are compatible with other experiments carried out on similar systems as well. For example, in the case of 5-nm-thick BaTiO3—BTO layers grown on LaSrMnO3—LSMO characterized by a single FE domain state, the FE polarization stabilizes toward the interface (Popescu, Barrett, et al., 2015), and this orientation of the polarization reflects in the gradual upshift of the Fermi level of LSMO as a result of the hole-depletion effect near the interface and the resulting n-type doping. The shift depends on the exponentially decreasing hole-depletion state in the first 1–2 unit cells of the LSMO slab close to the interface with BTO (Ma et al., 2014). The effect of ferroelectric-assisted hole depletion of the bottom LSMO electrode induced by the particular orientation of P modifies the band alignment close to the interface of the two materials and this is seen when probing the interface in more bulk-sensitive or interface-sensitive measurements. The consequent variation of the relative shifts between the surface and bulk core-level binding energies within the buried LSMO (blue and green components, respectively (Fig. 5.4b) derives from the ferroelectricinduced band bending and local hole depletion induced in the bottom electrode by the top ferroelectric layer (Fig. 5.4c).

2.2 Formation of charged defects Other mechanisms of screening of the depolarization charges and consequent compensation of the ferroelectric state may be needed, besides screening, when the contacting materials are nonideally metals or insulators (Bocirnea et al., 2020; Popescu et al., 2019, 2020). Sai et al. (2005) suggested ionic screening as a possibility to control the ferroelectric properties. Chisholm et al. (2010) added that ionic displacements in the PZT are uniformly bulk like to its interface with SRO, and no dead layer and no

Fig. 5.4 Sr 3d XPS spectra recorded on the two samples at normal incidence and at 60 degrees (a) on bare LSMO and (b) on BTO-covered LSMO. (c) The valence-band region of the LSMO/BTO heterostructure recorded at two collecting angles infers the effects of the built-in potential in the BTO on the valence-band maximum. The inset shows the band alignment and Schottky barrier formation. From Popescu, D. G., Barrett, N., Chirila, C., Pasuk, I., & Husanu, M. A. (2015). Influence of hole depletion and depolarizing field on the BaTiO3/ La0.6Sr0.4MnO3 interface electronic structure revealed by photoelectron spectroscopy and first-principles calculations. Physical Review B, 92, 235442.

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reduction in the polarization were found. It was established that the most efficient way to lower the energy of the PZT/SRO system, which has some degree of metallic screening, is a combination of metallic screening, ionic screening, and domain formation. On the other hand, in the monodomain PZT/STO system which lacks the metallic screening, the energy minimization is due to charge compensation at the interface through O vacancies (Wang et al., 2009). The calculations performed on the PbTiO3/STO superlattice revealed that oxygen vacancies prefer to be located in the ferroelectric PTO layer over STO, more exactly in the Pb-O planes in the PTO layer close to the interface and show that the switching of the ferroelectric polarization due to the modified chemical potential of the surface follows the electric potential (Wang et al., 2009). Wang et al. affirm that the chemical environment can play a dominant role in the behavior of nanoscale ferroelectrics and through the equilibration at high or low pO2, assisted by the piezoelectric effect, the reversal of polarization takes place. Remarkably, in 2010, Highlandet et al. (2010) showed that in thin films, the switching occurs by a continuous mechanism without domain formation at the intrinsic coercive field, EIC. If the initial monodomain polar state is subjected to reverse fields just below EIC, the switch domains do not nucleate even on the long timescales of the measurements. This is due to equilibrium or kinetic considerations, considering the ionic compensation at the surface. The explication resides in the second-order transition at zero fields (Pertsev et al., 1998) of PbTiO3 coherently strained to SrTiO3—when cooled under a nonzero applied electric field the continuously increasing polarization would not generate the crossing between nonpolar and polar states (Stephenson & Highland, 2011). The PRL 107 paper also confirms the origin of the charge used in the switching mechanism owed to the equilibrium behavior. It has been shown that when the polar phase loses stability, the paraelectric phase is the stable phase and the internal field is predicted to reach the intrinsic coercive field at these instabilities. On the other hand, Gao et al. (2017) used transmission electron microscopy to show that neither significant atomic displacements in the STO substrate nor 180-degree stripe domains in PZT were observed in a PZT/STO sample (Chisholm et al., 2010). They suggest as possible screening charges the positive charges provided by STO or the internal charges within the PZT thin film at the interface (Streiffer et al., 2002). For under 3 nm thickness, the compensation of polarization may be governed by the internal screening in ultrathin PZT films on an insulating substrate and metallic bottom electrode, with both the interface and surface containing ionic charges (Chisholm et al., 2010; Fong et al., 2006; Gao et al., 2016; Stephenson & Highland, 2011).

2.3 Adsorption of charged species Surface ionic charges may result from the adsorption processes (Tanase et al., 2016), and interact with the ferroelectric bound charges and can either screen (Chisholm et al., 2010; Gao et al., 2017; Levchenko & Rappe, 2008) or enhance the depolarizing field (Tian et al., 2018; Wang et al., 2009). Driven by the improvement in CO2 and H2O catalysis to address the current energy and pollution challenges, Garrity et al. (2013) investigated the variation in the

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properties of the PbTiO3 surface with polarization, which affects CO2 and H2O adsorption. Their ab-initio calculation suggested that a switchable stoichiometric TiO2-terminated PbTiO3 surface would be very useful for CO2 catalysis. They showed the formation of covalent bonds in the paraelectric and negatively poled surface and an ionic one for the positive case regarding CO2 adsorption, which comes from the screening electrons in the Ti conduction band states at the surface. The transfer of charge to the antibonding states of CO2 is associated with low dissociation energy for CO2 on the positively poled surface, which could initiate the catalytic reaction. Then, flipping the polarization on the negatively poled surface would release the products. Switching the polarization in thermodynamically stable PbTiO3 surfaces, one can bind and release CO2. They also confirmed the PbO surface termination. To eliminate the depolarization field and compensate for the surface charge, the stoichiometry of the top-most PbO layer changes in response to the polarization. In addition to CO2 binding, both the binding energy and binding mode of H2O depend on the polarization direction and oxygen coverage. Flipping the polarization, one can change the binding energy of the dissociated H2O and, by switching between a nonpolar and polar surface one can dissociate H2O into OH. Adding a full layer of SrRuO3 to a TiO2-terminated substrate, with the RuO2 layer on the surface, one could bind CO2 with the negatively poled surface, and then release it by switching the polarization, causing oxygen to replace CO2. Gattinoni et al. (2020) observed that in a PTO system with the polarization P pointing away from the substrate, its direction arises from the synergy of the surface and interface chemistry with the electrostatic properties of the ferroelectric. They have shown that the chemistry of the interface can favor a preferential direction of the polarization at the interface, which does not affect the overall direction of the polarization. However, it leads to a reduction of P at the interface in the PbOint system (RuO2-PbO interface), while the other studied system, TiOint (SrO-TiO2 interface) presents only a very small preference for an upward-pointing direction. They concluded that the interface sets the polarization direction for PTO thin films on SRO through its influence on the overall band alignment and through the local chemistry and bonding. The calculations revealed that adsorbates such as OH species and O2 adatoms, or defects such as cation vacancies, negatively charge and stabilize Pup, while the positive ones such as H+ adsorbates, Pb2+, and Ti4+ adatoms, or O vacancies favor Pdown. The conclusion was that surface and interface nonstoichiometry promote stabilization of the noncooperative defects accompanied by a reduction of the polarization in the thin film with respect to the bulk. So, by the engineering of surface and interface screening one can control the polarization (Wang et al., 2009; Xie et al., 2017) through defects or adsorbates. The compensation of the ferroelectric domains by means of ambient carbon adsorbate and its signature in the evolution of the XPS spectra has been established by Stoflea et al. (2014), and shows that molecules of contaminants (fatty acids, alcohols, and esters) are adsorbed preferentially on areas exhibiting outward polarization. These examples illustrate not only the fundamental mechanisms of compensation of the ferroelectric states but also put forward the perspective that these mechanisms create new functionality, which certainly deserves attention in the future.

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A case that illustrates the opposite scenario and the involved consequences in photoemission is that of completely uncompensated ferroelectrics, revealing that the strong internal field of the ferroelectrics as discussed in a study by Popescu, Husanu, et al. (2015). The experiments carried out in ultrahigh vacuum and under X-ray synchrotron radiation on a 250-nm PZT (001) sample after in situ annealing in an oxygen environment revealed that the FE sample becomes atomically clean and the stoichiometry of the layer is rebuilt to preserve both the FE and the insulating character. Hence, it cannot provide enough charge carriers anymore to compensate the depolarization field. When irradiating with soft X-rays, the electrons generated in the bulk migrate near the surface and the holes near the bottom interface. These leads to a decrease in time of the overall band bending of the film. Fig. 5.5a and c evidences the formation of partially compensated areas on the ferroelectric, where the “uncompensated” charge density yields a value of about 5 μC cm2. In Fig. 5.5b, in some areas the “compensated” situation (the red photoemission curve) can be obtained, but there are also areas that present concomitantly two “uncompensated” situations. The representations in Fig. 5.5d–g suggest that the sharp component shifted toward lower binding energies is reduced lead. The area that ejected Pb is oxygendepleted; thus, the Fermi level is pinned close to the conduction band, owing to a strong local n-type doping. Such mechanisms depend, however, on the crystalline structure of the system, with the (111) orientation preventing the formation and migration of Pb in PZT(111) (Husanu, Popescu, Tache, et al., 2015). This suggests that for catalysis applications, certain facets of the grains are more active in the adsorption and desorption of molecular species (Takata et al., 2020).

3

Improper ferroelectrics

The appearance of the ferroelectric state in perovskite materials is at its most fundamental level a result of the particular hybridization of the electron clouds corresponding to neighboring ions in the unit cell. This leads to asymmetrical shifts in the equilibrium ion positions, which promote a permanent dipole moment, also called displacive ferroelectricity. The ionic displacement is particularly energetically favorable when the 3d shell is empty. On the other hand, the magnetic transition-metal ordering needs partially filled 3d shells (Hill, 2000). This incompatibility induced a new search direction for new materials with ferroelectricity guided by nondisplaced mechanisms adapted to magnetic order so-called improper ferroelectrics. In contrast to the ordinary ones, the order parameter of the phase transition in improper ferroelectrics is not the polarization (spontaneous polarization arises only as a secondary effect), but another leading-order parameter whose transformation properties are different from those of the polarization (Indenbom, 1960a, 1960b, 1960c). Here, the temperature dependence of the permittivity does not obey the Curie-Weiss law. Also, an electric field does not suppress the phase transition. More precisely, improper ferroelectricity offers ways to achieve coexistence between electric and magnetic order (Benedek & Fennie, 2011; Cheong & Mostovoy, 2007; Kimura et al., 2003; Sim et al., 2016; Zhou et al., 2019). The domain structure of

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Fig. 5.5 Spectromicroscopic data obtained on the 250-nm PZT(001) layer after annealing in oxygen: (a) Spectromicroscopic image obtained on annealed PZT(001). Superposed: time evolution of the Pb 5d signal. (b) Spectromicroscopic image obtained over the whole energy range designed in the superposed graph, with two typical spectra obtained on areas outlined with blue and red circles. (c) The analysis of the time evolution of the Pb 5d signal from (a). Inset in the upper right corner: (c-i1) binding energies and (c-i2) integral amplitudes resulted from deconvolution. (d) An area plot of the total intensity, with the regions where the Pb 5d time evolution was recorded outlined by the red circle, and the time evolution of the valence band represented in (e) (just the starting and the ending situations) recorded in the area outlined by the blue circle. (f) Time evolution of the Pb 5d signal, together with deconvolutions by using three components, whose integral amplitudes are represented in (g). The binding energies of individual components do not vary considerably in this case and their evolution is not represented. From Popescu, D. G., Husanu, M. A., Trupina, L., Hrib, L., Pintilie, L., Barinov, A., Lizzit, S., Lacovig, P., & Teodorescu, C. M. (2015). Spectro-microscopic photoemission evidence of charge uncompensated areas in Pb(Zr,Ti)O3(001) layers. Physical Chemistry Chemical Physics, 17, 509.

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the improper ferroelectrics provides domains that do not differ in their polarizations. In the infrared spectrum, the soft mode in the nonpolar phase is inactive, as the polarization is not related to the loss of stability in an improper ferroelectric phase transition (Levanyuk & Sannikov, 1974). Moreover, unlike the polarization in thin proper ferroelectric layers, improper ferroelectricity is expected to be less influenced by the depolarizing-field effects, such as the critical thickness for spontaneous polarization (Fong et al., 2004; Sai et al., 2009; Stengel et al., 2012). More exactly, the spontaneous polarization Ps from improper ferroelectrics still results in the buildup of a depolarizing field, reminiscent of conventional ferroelectrics, which is however stabilized by the primary order parameter. Hence, the leading-order parameter does not have a polarization, which is why it is not strongly influenced by the electrostatics of the heterostructure, making the depolarizing field less effective, supporting a possible small attenuation of Ps, but not involving a critical thickness (Sai et al., 2009; Stengel et al., 2012). One of the well-studied systems with geometrically driven ferroelectric order is YMnO3 (Fennie & Rabe, 2005; Nordlander et al., 2019; Van Aken et al., 2004). In this improper system, the ferroelectric vortex domain pattern (Choi et al., 2010; Jungk et al., 2010), ferroelectric domain walls with tunable conductance (Meier et al., 2012; Mundy et al., 2017; W. Wu et al., 2012), and coexistence with magnetic order (Fiebig et al., 2016) depend on the nature of the spontaneous polarization. Nordlander observed a decrease of TC with YMnO3 film thickness. Also, the spontaneous polarization reaches zero for the two-unit-cell film, establishing a threshold thickness for room-temperature ferroelectricity, similar to the depolarizing-field effects in conventional ferroelectric thin films ( Junquera & Ghosez, 2003). They suggest two possible explanations, one related to the decoupling in the thin-film limit of the order parameters Q and Ps (the polarization becomes proper), and the other considers an unidentified mechanism responsible for the suppression of Ps in the first two-unit cells of the YMnO3 thin films with coupled Q and Ps. Using in situ optical second-harmonic generation (ISHG) correlated with transmission electron microscopy at the atomic level, the preserved improper relation between Ps and Q in the bulk was confirmed (Lilienblum et al., 2015) down to the ultrathin limit, where the behavior of Ps in the thin films is directly guided to Q and any mechanism acting on it and not to the depolarization field (Sai et al., 2009).

3.1 Multiferroics Historically, magnets were explored for far many years compared to their more recent ferroelectrics counterpart. Many attempts were made to create a so-called multiferroic material combining ferromagnetism and ferroelectricity in the same phase of a material (Smolenskii et al., 1961). The most intensely investigated compounds were boracites (for example, Ni3B7O13I) due to their pronounced linear magnetoelectric effect with hysteretic switching of multiferroic state by either electric or magnetic fields (Ascher, 1966). Smolenskii and Chupis (1982) and Schmid (2012) summarize well the experimental, theoretical, and applied achievements of these early days in the field. The first work on multiferroic domain walls and magnetoelectric domain

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coupling effects was published by Fiebig et al. (2002). They reported spatial maps of coupled antiferromagnetic and ferroelectric domains in YMnO3 revealed by the spatially resolved second harmonic light. In 2003, Kimura et al. observed giant magnetoelectric coupling effects in a type II multiferroic material (TbMnO3), attributed to switching of the electric polarization induced by magnetic fields. The induction of magnetization by means of an electric field and induction of polarization by means of a magnetic field, also known as a magnetoelectric effect, originate from offcentered spin textures that induce a magnetically controllable electric polarization. Multiferroics are interesting materials because they may lead to faster, smaller, more energy-efficient data-storage technologies, making it possible to exploit the functionalities of both ferroic orders. A four-state memory element could in principle be achieved when an electric bit complements a magnetic one (see Fig. 5.6). The waste heat and relatively long buildup time associated with electric currents are avoided when generating electric current using a voltage pulse rather than a magnetic field. In this context, the coupling of the ferromagnetic and the ferroelectric states might induce novel functionalities not present in either state alone. Based on the mechanisms inducing multiferroicity, we may distinguish four main classes of these materials, which permit the coexistence of ferroelectric and magnetic order. In the first three classes: ferroelectricity driven by electronic lone pairs, charge ordering, and geometric effects, the magnetic and ferroelectric orders occur independently (type I—Fig. 5.6e) and in the last one: magnetism guided ferroelectricity, the magnetic and ferroelectric transitions arise simultaneously (type II—Fig. 5.6f) (Fiebig et al., 2016). Ferroelectricity due to lone pair develops in the presence of spatial asymmetry initiated by the anisotropic distribution of unbonded valence electrons around the host ion (Fig. 5.6e) (Barone & Picozzi, 2015; Fiebig et al., 2016). Ferroelectricity due to charge ordering refers to the noncentrosymmetric distribution of valence electrons around their host ions in the crystal structure to form periodic superlattices (Fig. 5.6e) (Barone & Picozzi, 2015; Fiebig et al., 2016). Hybrid improper ferroelectricity arising from a symbiotic interplay of rotation and tilting distortions leads to ionic shifts that result in the formation of a polar state. Octahedron rotations in perovskites, known to strongly couple to magnetic properties, can induce local polar displacements of A-site cations through efficient force fields applied by O ions (Imada et al., 1998) (Fig. 5.6e). For example, in the geometric ferroelectric AMnO3 hexagonal manganite, the magnetic ordering comes from the electronic d4 configuration of Mn3+ on the B-site, or from a magnetic rare-earth ion on the A-site. The Mn3+ is located close to the barycenter of the bipyramid (van Aken et al., 2001) and the shifts of the A3+ modulate the ferroelectric moment. The stability of the structure preserves only for small rare-earth ions on the A-site, like yttrium, and for Mn3+ on the B-site. When considering the other transition metals on the B-site and large ionic radius elements on the A-site, the perovskite system is favorable (Abrahams, 2001; Bos et al., 2001). Magnetically induced ferroelectricity manifests as a secondary transition induced by the magnetic ordering, which can break inversion symmetry (Tokura et al., 2014). The symmetry breaking is induced by the magnetic electronic ground state or by

Fig. 5.6 Example of new functionalities: One could control the magnetic properties by electric fields instead of magnetic fields. (a) Ferromagnetic hysteresis loop; (b) the relationship between multiferroic and magnetoelectric materials; types I and II transitions; (c) ferroelectric hysteresis loop; (d) multiferroic hysteresis loop; (e) type I multiferroics; and (f) type II multiferroics. No Permission Required.

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structural distortions, which depend on the magnetic moments. So far, it was shown that improper ferroelectricity emerging as a consequence of spin order rely on a couple of distinct magnetoelectric coupling mechanisms (Fig. 5.6f), as discussed in specific reviews on the topic (Kimura, 2007; Picozzi & Stroppa, 2012; Tokura et al., 2014). Among most intensely studied are the materials with noncollinear magnetic order, where antisymmetric exchange striction induces the electrical polarization. The polarization in cycloidal spin spirals (RMnO3) is driven by the competing ferromagnetic nearest-neighbor and antiferromagnetic next-nearest-neighbor interactions, which stabilize the spin spiral state (Kenzelmann et al., 2005). In the spin rotation plane, an electric dipole moment develops (Mostovoy, 2006). In the Dzyaloshinskii-Moriya (DM) interaction, D  [Si ×Sj], where D is a constant vector and Si,j are neighboring spins at sites i and j (Dzyaloshinskii, 1964; Moriya, 1960) a noncentrosymmetric crystallographic environment promotes a noncollinear spin ordering. On the other hand, inverse DM interaction is characterized by a noncentrosymmetric displacement of charges by an acentric spin structure (Katsura et al., 2005; Mostovoy, 2006). In both cases of DM interaction, the spin-orbit coupling is vital for the polarization. The favorable spin cycloid rotation plane is influenced by the single-ion anisotropy. The resulting polarization P ∝ eij ×Si ×Sj (eij represents the unit vector connecting neighboring spins) is driven by the optimization of the spin configuration with consideration of antisymmetric exchange. Some typical magnetically induced ferroelectricity was found in Cr2BeO4 (Newnham et al., 1978), TbMnO3 (Kenzelmann et al., 2005), DyMnO3 (Arima et al., 2006), BiFeO3 (Shen et al., 2021), R2BaCuO5 (R ¼ Dy and Ho) (Yanda et al., 2021), and NaMn7O12 (Gastaldo et al., 2020). The second group of materials shows collinear magnetic order and the electrical polarization is driven by the symmetric exchange striction P ∝ Si  Sj (Sergienko & Dagotto, 2006). The neighboring spins shift in order to optimize their exchange energy, resulting in an increased polarization jSi  Sj j > j Si ×Sj j. One of such case was observed in TbMnO3, where a cycloidal order (parameterized by Si  Sj) changes under pressure into a collinear antiferromagnetic order (parameterized by Si  Sj) (Aoyama et al., 2014). Other studies on the nonrelativistic symmetric exchange were on Ca3Co1 xMn1 xO6 (Zubkov et al., 2001) and YMn2O5 (Chapon et al., 2006) systems. The third group includes materials with electrical polarization induced by spin-dependent p-d hybridization with the p-orbitals of the oxygen ligands. This spin-dependent p-d hybridization yields a spontaneous polarization along the bond direction in delafossites, creates different local dipole moments along with the CodO bonds in Ba2CoGe2O7, with a magnetoelectric coupling mechanism based on a single localized spin coordinated to a ligand-mediated via spin-orbit coupling ( Jia et al., 2006, 2007). In other cases, multiple coupling mechanisms are encountered (Tokura, 2006). The polarization is expressed by the relation (5.6), describing the metal-ligand bond (Fig. 5.6f).  2 P∝ Si  eij eij

(5.6)

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3.2 Type III multiferroics: Domains and domain walls Ferroic materials (ferroelectrics, ferromagnetic, ferroelastic) present areas with order parameters oriented in different ways (e.g., ferroelectric polarization in ferroelectrics) named domains separated by domain walls. Improper ferroelectric materials present major importance not present in the ones having as primary order parameter the polarization. They permit the formation of domains and domain walls, which add new properties and functionalities. The difference in size between smaller ferroelectric domain walls and larger magnetic ones is at least one order of magnitude starting from a few nanometers. Even if at the walls charge or spin phenomena similar to the ones appearing at heterointerfaces can rise, they can possess different functional properties than the domains themselves. The possibility to be created, moved, and erased postgrowth represents an advantage useful in device applications, like storage memory nanodevices (Al Bahri, 2022; Meng et al., 2022; Xu et al., 2022). The difference between types I and II multiferroics is shown in Fig. 5.7. Due to the independence of magnetic and electric order, magnetic and electric walls in type I multiferroics can coincide but do not have to. On the other hand, the magnetic order induces the electric order in type II multiferroics with multiferroic domain walls.

4

Rashba ferroelectrics

In noncentrosymmetric crystals, electronic energy bands are split by spin-orbit coupling (SOC) (Dresselhaus, 1955; Rashba, 1960). SOC is characterized by the presence of a magnetic field in the electrons frame of motion, which are moving in an electric field even in the absence of an externally applied magnetic field. This field that couples to the electron’s magnetic moment are the spin-orbit field. SOC couples spin degrees of freedom with electronic orbits linking the spin space and the real space. The importance of SOC appears both in nonmagnetic solids, such

Fig. 5.7 Example of type I (a) and type II (b) domain walls. In nature (a), domain walls are either magnetic (yellow), or electric (blue). A multiferroic domain wall is formed (red) when the domain design coincides and is influenced by the coupling between magnetic and electric order. In type II multiferroics, the magnetic order promotes the electric order. Furthermore, all ferroelectric domain walls are likewise magnetic, forming this way multiferroic domain walls. No Permission Required.

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as topological matter, spin valley, coupling pseudospin, Rashba-Dresselhaus effects (i.e., k-dependent spin splitting in the band structure), but also in magnetic solids— magnetic anisotropy, Dzyaloshinskii-Moriya weak multiferroics, skyrmions, anisotropic exchange. The increasing interest in SOC phenomena also led to intensive research in spintronics, quantum computing, and cold atom system (Manchon et al., 2015; Varignon et al., 2018). According to the Rashba effect, the dispersion of free-like electrons subject to a potential gradient and in the presence of SOC splits into two branches for oppositely spin-polarized states. The energies of the two branches are described by the relation: E ð k Þ ¼

h2 k 2  αR ðkÞ 2m∗

(5.7)

with k momentum, m* effective mass, αR ∝ λ Ez Rashba parameter—expressing the strength of the Rashba effect, λ is the spin orbit constant, Ez is electric field, resulting a Rashba momentum kR ¼ (j m∗j αR)/ħ2 evaluated as the Rashba-induced momentum offset of the valence band maximum (VBM) or conduction band minimum (CBM) with respect to the a high-symmetry point and αR ¼ 2ER/kR if the the third-order terms in the nearly free-electron approximation is neglected, ER is calculated as the difference between the VBM or CBM estimated at kR and the corresponding energy values at the high-symmetry point considered (Di Sante et al., 2013). Ferroelectric Rashba semiconductors (FERSC) are obtained when combining spinorbit coupling with ferroelectricity (Fu et al., 2021; Picozzi, 2014; Yang et al., 2021), creating the possibility to electrically switch the spin texture related to the Rashba spin splitting (RSS) upon reversal of the ferroelectric polarization (see Fig. 5.8). Di Sante et al. (2013) theoretically revealed a full reversal of the spin texture (i.e., of the Rashba parameter) when switching the ferroelectric polarization in GeTe with applications in Datta-Das spin transistor. The electrical control is made as to the spin direction in each subband changes by 180 degrees upon reversal of polarization. They suggest that the modified “Datta-Das” spin-FET architecture allows the combination of logic and storage functionalities presenting three memory elements (two magnets and one FE) and one logic channel operating with spins and controlled by the FE polarization. An ideal picture of a ferroelectric Rashba semiconductor considers nonmagnetic ferroelectrics insulators with a sizable switchable polarization and a reasonable bandgap. Heavy ions with large SOC exhibiting a significant Rashba spin splitting close to the valence or conduction band edge are involved. The presence of switchability with the polarization and the possibility of doping useful in applications based on spin/charge currents are required. In d0 ABO3 perovskites with a transition metal at the B-site (Hill, 2000), the bandgap is formally between O-2p and B-d states. In this case, one can obtain large Rashba spin splitting around the bandgap using a heavy cation at the B-site with an efficient polarization control due to B-type ferroelectricity. Although the study of reacted interfaces from metals and semiconductors is not a new field (Gheorghe et al., 2011, 2012; Husanu, Ganea, Anghel, et al., 2015; Husanu, Popescu, Ganea, et al., 2015; Lungu et al., 2013; Popescu & Husanu, 2013, 2014), the

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Fig. 5.8 Ferroelectric Rashba semiconductors. The link between materials science and technology. No Permission Required.

exploration of Rashba-related interfaces driven by the breaking of the symmetry is at the beginning. Recently, a giant spin-to-charge conversion at the Fe/Ge(111) interface was obtained by Oyarzun et al. (2016). With a thin MgO barrier used to separate the Fe(111) and Ge(111) showing a drop of the signal that made clear the importance of Fe/Ge(111) interface in obtaining huge spin-to-charge conversion. The Rashba effect into spin-split metallic states made this possible even if the Fe and Ge atoms are relative light. They generated very large charge currents by direct spin pumping into the interface states, which were transformed into a two-dimensional (2D) charge current sustained by the Rashba SOC. Ab-initio calculations assured the existence of metallic exchange and spin-orbit split Fe/Ge(111) interface states. The calculations validated the existence of a large density of metallic states exhibiting both p and d characters originating from hybridized Ge and Fe states in which spin-to-charge conversion might occur. Through calculations, they demonstrate the existence of the Rashba states at the Fe/Ge(111) interface, which explained the large spin-to-charge conversion. Importantly, the study concluded with a giant spin-to-charge conversion at the interface of two light atoms, which are a milestone for developing electronics compatible with silicon-based technology. In 2019, another study performed on Fe films deposited on α-GeTe(111) surfaces by Sławi nska et al. (2019) showed how the Fe capping layers stabilize the GeTe

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surfaces with the two different polar configurations close to the surface. They compared the input of Fe layer on GeTe vs the bare GeTe surface on both cases, with the electric dipole pointing outwards and inwards, where the latter is unstable. A large interdiffusion of Fe ions within the GeTe substrate was observed for ultrathin Fe thicknesses (1ML and 2ML), which altered the GeTe(111) structure. On the other hand, when increasing Fe deposition for more than 3ML of the topmost surface atoms in GeTe are not so affected making it suitable for practical purposes. Taking into account both directions of  P and different thicknesses of the Fe overlayer a study concerning the electronic structures and spin textures was made. It resulted in a suppression of the Rashba surface states due to the hybridization of the Fe states with the GeTe surface was observed. Remarkably, the bulk Rashba bands remain almost electronically unchanged during slight alteration at the interfacial GeTe layer in agreement with the excellent screening properties of GeTe. Rashba bulk bands in FERSC originate from inversion symmetry breaking, associated with the polar axis in ferroelectrics, in contrast with the Rashba effect studied at surfaces where inversion symmetry is intrinsically broken. However, GeTe is known for its small bandgap and related large leakage currents that may prevent polarization switching. Different alternative directions have been explored (da Silveira et al., 2016; Di Sante et al., 2016; Narayan, 2015; Stroppa et al., 2014; Varignon et al., 2019; Yamauchi et al., 2015; H. Zhang et al., 2017; Zhong et al., 2015) to find the best candidate for concrete applications. In 2019, Djani et al. (2019) identified Bi2WO6 Aurivillius crystal as a robust ferroelectric with large and reversible Rashba spin splitting, which can be significant n-type dopped without loss of its ferroelectric properties. The practical ferroelectric Bi2WO6 presents important properties, like a unidirectional spin-orbit field that protects the spin transport from spin dephasing and the persistence of robust polar distortion, large Rashba spin splitting and unidirectional spin-orbit field after n-doping. The spin-orbit field is linked with the in-plane polarization, strong layering-induced anisotropy in the electronic system, and associated symmetry properties mixing. Comparable behavior can be found also in other ferroelectric Aurivillius phases (SrBi2Ta2O9, Bi4Ti3O12) if one takes into account that Rashba spin splitting (Sasmito et al., 2021) relies on the strength of the B-cation SOC that increases with the oxidation state (Dai et al., 2008) and Ps.

5

Conclusions and outlook

Ferroelectric materials are characterized by the presence of an electric dipole that spontaneously lines up in clusters (domains) offering the possibility of orientation in one direction when an external electric field is applied. These materials present a ferroelectric hysteresis, named by analogy with ferromagnetic hysteresis, and presents ferroelectricity only below Curie temperature of the material. Rising the temperature above TC, the spontaneously forces presented in the system are overcome by dipoles agitation due to the heating process. Ferroelectrics are also piezoelectric and pyroelectric with applications from high-energy-density capacitors—compact

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batteries (due to their high dielectric constant, they can store lots of energy) and nightvision devices (some of them have high “pyroelectric coefficient”) to ultrasound medical equipment (due to piezoelectricity, one can generate an electrical voltage in FE when put under pressure by an object that can be used to create an image of the object—imaging unborn babies, the hydrophone, to map the topography of the ocean floor), smart technologies for energy harvesting (harvest the energy from cars and lorries that is otherwise lost as heat or noise, battery-powered mechanical pacemakers recharged by the voltage generated in a ferroelectric material directly from the thrust of the heartbeat), and actuators and translators (used in atomic force microscopes, scanning tunneling microscopes, piezoforce microscopes and magnetic force microscopes). Recently, some exotic physics were obtained in such materials linked to topological defects called “polar skyrmions” and “polar hopfions.” The potential to electrically control magnetism (multiferroic materials) to create new devices within the framework of spintronics, information storage, and communication, forms the catalyst that is energizing worldwide research activity. Also, FERSC can be used as a nonvolatile channel in a spin-FET device with two additional magnetic elements, working as spin injector and detector when in contact with normal magnets.

Acknowledgments We are grateful to the authors who allowed the reproduction of their work here. D.G.P. was supported by the national fellowship program L’Oreal—Unesco “For Women in Science” and by the Romanian UEFISCDI Agency under Contract PN-III-P1-1.1-TE-2021-0136 (TE 50/2022) of the Romanian Ministry of Research. M.A.H. was supported by the Romanian UEFISCDI Agency under Contract PCE 96/2021 of the Romanian Ministry of Research.

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Crystal structures of copper oxide-based perovskite compounds

6

Takeo Oku Department of Materials Science, The University of Shiga Prefecture, Hikone, Japan

1

Introduction

To solve global warming due to carbon dioxide (CO2), development of new energy to surpass conventional fossil fuel is mandatory. Solar cells are one of the important methods for solving the energy problems, and new thin-film solar cells with various perovskite compounds have been developed recently (Han et al., 2019; Lin et al., 2021; Mao et al., 2019; T. Oku, 2020; Tavakoli et al., 2019; Ueoka & Oku, 2020; Wang et al., 2019). However, energy densities of solar cells are not so high. On the other hand, nuclear fusion produces high energy density and forms no CO2. Therefore, the nuclear fusion energy is also considered as one of the substitutes of fossil fuel in the 21st century (Fietz et al., 2017; Fischer et al., 2020; T. Oku, 2018). A proposed nuclear fusion power station called DEMO will be a totally integrated device with various apparatus including a superconducting magnet for plasma confinement, blanket, diverter, plasma-facing wall, heating, diagnostics, cryogenics, and remote maintenance. Especially, the development of superconducting materials for the magnet in DEMO and actual fusion devices is one of the important issues, and copper oxide-based perovskite compounds are one of the promising candidates (Bruzzone et al., 2015; Uglietti et al., 2015). The superconducting properties are strongly dependent on the atomic arrangements, crystal structures, and microstructures of the perovskite compounds (Larbalestier et al., 2001). Transmission electron microscopy (TEM) using high-resolution observation and electron diffraction has been applied for microstructural analysis of various advanced materials. Superconductors with high critical temperatures (TC) based on the perovskite-type structure are suitable subjects for high-resolution electron microscopy (HREM) and electron diffraction, and many TEM studies of these materials have been carried out. In the initial stage of high-TC superconductor studies (Bednorz & M€uller, 1986; Cava et al., 1988; Chu et al., 1993; Maeda et al., 1988; Sheng & Hermann, 1988; M. K. Wu et al., 1987), individual cations are represented as separated dark spots (T. Oku, 2001, 2011) in high-resolution images of Y- (K. Hiraga et al., 1987), Bi- (K. Hiraga, Hirabayashi, et al., 1988; K. Hiraga et al., 1989), and Tl-based (K. Hiraga, Shindo, et al., 1988) superconductors.

Perovskite Ceramics. https://doi.org/10.1016/B978-0-323-90586-2.00011-5 Copyright © 2023 Elsevier Ltd. All rights reserved.

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In the present article, studies on the crystal structures of copper oxide-based perovskite compounds by high-resolution electron microscopy and electron diffraction were reviewed. High-resolution structure images and selected electron diffraction patterns give us valuable information for crystal structure analysis (S. Nakajima et al., 1989; T. Oku, 2012). The main contents of the present chapter are as follows. In the first part, the necessary experimental conditions to observe the high-resolution structure images which can be used for crystal structure analysis were discussed, and information obtained from the structure images was mentioned. In the main part, crystal structures and microstructures of Tl-, Pb-, Ln-, Y-, CO3-, BO3-, Hg-, and Bisystem superconducting copper oxide-based perovskite compounds were presented, which were determined from the observed high-resolution structure images and electron diffraction patterns. Defects, intergrowth, interfaces, and surfaces of the perovskite-type superconducting copper oxides are also described.

2

Atomic observation by HREM

HREM images are phase contrast due to interference of electron waves, and direct information on atomic arrangements in crystals can be obtained. HREM images are sensitive to defocus values, crystal thickness, crystal tilting, and other parameters. Present electron microscopes have enough resolutions to observe atomic columns in simple crystal structures. Therefore, high-TC superconductors based on the perovskite-type structures are suitable for HREM, and a number of structural studies

Fig. 6.1 (A) HREM image of TlBa2Ca3Cu4O11 taken with the incident beam parallel to the a-axis. (b) HREM image after crystallographic image processing of (a), together with a structure model of TlBa2Ca3Cu4O11. Ov corresponds to oxygen vacant positions. No permission required.

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157

of these materials have been carried out. Fig. 6.1A is an HREM image of TlBa2Ca3Cu4O11 (Liang et al., 1988; S. Nakajima et al., 1989), which has the almost highest transition temperature of 123 K, taken with the incident beam parallel to the a-axis together with a structure model (JEM-4000EX, accelerating voltage of 400 kV), and the upper right side of the image is the crystal edge. Fig. 6.1B is an HREM image modified by Fourier transform and crystallographic image processing, which is described in detail in the next section. In Fig. 6.1, the images of the crystal thinner than ca. 2 nm faithfully represent the projection of the crystal structure. Darkness and sizes of dark spots corresponding to Tl, Ba, Cu, and Ca positions can be identified as being nearly proportional to their atomic numbers (M. Kikuchi et al., 1990; S. Nakajima et al., 1990, 1991; D. Shindo et al., 1991). Such a high-resolution image is called structure images. However, in images thicker than 3 nm, the cation positions cannot be identified as dark spots. Image calculations showed that HREM images taken at nearly Scherzer defocus and taken from crystals thinner than 2 nm can be used for structure analysis of high-TC superconducting oxides. The dark spot positions corresponding to the cations in the structure images of Fig. 6.1 faithfully reflect their real atomic positions. Atomic coordinates of the cations measured from the dark spot positions in Fig. 6.1B are listed in Table 6.1, in comparison with atomic coordinates determined by X-ray diffraction (XRD) (Liang et al., 1988). As listed in Table 6.1, the dark spot positions in the observed structure image reflect real positions of the cations within an error of 0.01nm, which may correspond to a measured error. This result indicates that the cation positions can be determined with the precision of 0.01 nm from the structure images. Although oxygen atoms are not represented as dark spots in Fig. 6.1, information on the oxygen atom positions is contained in the structure image. It was also found that the observed image have valuable information not only on accurate coordinates of the cations but also on ordered arrangements of oxygen vacancies, as indicated by Ov in Fig. 6.1B. Unfortunately, oxygen atoms cannot be represented as dark spots in the structure images taken by the present microscope. However, information on the oxygen atom positions is contained in the structure images. Fig. 6.2 is the calculated images of TlBa2Ca3Cu4Ox, and two types of structure models with oxygen vacancies and with oxygen atoms on the Ca layers are shown in Fig. 6.2A and B, respectively. The image Table 6.1 Structural parameters of cations in TlBa2Ca3Cu4O11, measured from the HREM images and XRD. Atom

x

y

z (HREM)

z (XRD)

Tl Ba Cu Cu Ca Ca

0 0.5 0 0 0.5 0.5

0 0.5 0 0 0.5 0.5

0 0.14 0.245 0.413 0.33 0.5

0 0.147 0.25 0.417 0.334 0.5

Space group P4/mmm. a ¼ 0.3848 nm. c ¼ 1.908 nm.

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Perovskite Ceramics

Fig. 6.2 Structure models and calculated images of TlBa2Ca3Cu4O11 with oxygen vacancies on the Ca layers (A), and those of TlBa2Ca3Cu4O14 with oxygen atoms on the Ca layers (B). Intensity profiles along the lines through Tl and Cu atoms parallel to the c-axis are also shown at the right side of the image. No permission required.

calculations were performed under the condition of a defocus value of 45 nm and crystal thickness of 1.92 nm. Intensity profiles along the lines through Tl and Cu atoms parallel to the c-axis are also shown at the right side of the calculated images. Comparing Fig. 6.2A and B, it is noticed that the oxygen vacant positions indicated by Ov show brighter contrast than the oxygen occupied positions, and the intensity of the vacant positions is higher in the intensity profile. To confirm the above results of image calculations, the observed structure image of TlBa2Ca3Cu4O11 in Fig. 6.1B is examined (T. Oku, 2014a). In the observed image, the oxygen vacant positions indicated by Ov show brighter contrast than the other bright regions corresponding to oxygen positions. Therefore, the image clearly shows the existence of oxygen vacancy ordering on the Ca layers. The present methods for nanostructural analysis will also contribute to the development of perovskite solar cell materials (T. Oku, 2020; T. Oku et al., 2018, 2014).

3

Crystal structures of Tl-based copper oxides

HREM structure determination of Tl-based superconducting copper oxide perovskites is described here (T. Oku, 2014a). Bulk samples of Tl-based copper oxides were prepared by reacting Tl2O3, CaO, BaO2, CuO, and BaCuO2 (S. Nakajima et al., 1990, 1991, 1989). The samples of Tl2Ba2CuO6 and TlBa2CaCu2O7 were reheated at temperatures in the range 500–850°C, and then quenched in liquid nitrogen in order to control oxygen contents and to obtain the highest TC. The Tl2Ba2CuO6, Tl2Ba2CaCu2O8, and TlBa2CaCu2O7 samples provided superconductivity at several transition temperatures. Samples for HREM observation were prepared by dispersing crushed materials on holey carbon films. Fig. 6.3A and B are the high-resolution structure images of Tl2Ba2CuO6 (T C ¼ 80 K) taken with the incident beam parallel to the [100] and [110] directions, respectively (T. Oku, 2014a). To observe the atomic arrangements clearly, crystallographic image processing was carried out. The digital images were masked and fast Fourier transformed. Reciprocal lattices were indexed according

Crystal structures of copper oxide-based perovskite compounds

159

Fig. 6.3 High-resolution structure images of Tl2Ba2CuO6 (TC ¼ 80 K) taken with the incident beam parallel to the (A) [100] and (B) [110] directions. (C, D) HREM images after crystallographic image processing of (A) and (B), respectively. No permission required.

to the unit cell, and the lattice parameters were determined using the positions of the strongest peaks in the analysis. The local background was subtracted, and the amplitudes and phases of the peaks were corrected using symmetrization (Hovm€oller et al., 1984; Weirich et al., 1996). Before correcting the phases, the phase origin was determined by investigating the origin shift that gave the best accordance with the phase conditions for the two-dimensional plane group. Averaged symmetrized images were reconstructed from the corrected Fourier transforms. Crystallographic symmetrization based on the two-dimensional space group was used for the reconstruction of the Fourier transform, as shown in Fig. 6.3C and D. Tl, Ba, and Cu atoms are clearly observed, and the darkness and sizes of the dark spots corresponding to the metal atom positions can be identified as being nearly proportional to their atomic numbers. A high-resolution structure image of Tl2Ba2CaCu2O8 (TC ¼ 114 K) is shown in Fig. 6.4A, which was taken with the incident beam parallel to the a-axis. Fig. 6.4B is an HREM image after crystallographic image processing of Fig. 6.4A, and metal atom positions are clearly observed. In addition to the metal atom positions, oxygen vacancies are clearly observed as bright white dots on the Ca layers. Fig. 6.5 are high-resolution structure images of TlBa2CaCu2O7, Tl2Ba2Ca2Cu3O10, TlBa2Ca4Cu5O13, and Tl2Ba2Ca3Cu4O12, taken with the incident beam parallel to the

Fig. 6.4 (A) High-resolution structure image of Tl2Ba2CaCu2O8 (TC ¼ 114 K) taken with the incident beam parallel to the a-axis. (B) HREM image after crystallographic image processing of (A). No permission required.

Fig. 6.5 High-resolution structure images of (A) TlBa2CaCu2O7 (TC ¼ 110 K), (B) Tl2Ba2Ca2Cu3O10 (TC ¼ 122 K), (C) TlBa2Ca4Cu5O13, and (D) Tl2Ba2Ca3Cu4O12 taken with the incident beam parallel to the a-axis, together with projected structure models. Brighter spots indicated by asterisks correspond to the oxygen vacant positions. No permission required.

Crystal structures of copper oxide-based perovskite compounds

161

a-axis, together with projected structure models (M. Kikuchi, Nakajima, et al., 1989). Metal atom positions are clearly observed, and brighter spots indicated by asterisks correspond to the oxygen vacant positions. It is well known that the Tl-based superconductors have various layered structures such as TlBa2Can1CunO2n+3 (n ¼ 1–5) with single Tl layers and Tl2Ba2CanlCunO2n+4 (n ¼ 1–4) with double Tl layers. Their structure models were proposed, as shown in Fig. 6.6. For the TlBa2CanlCunO2n+3 (n ¼ 1–5) system, structure parameters of TlBa2CaCu2O7 (Tl-1212) (Herview et al., 1988), TlBa2Ca2Cu3O9 (Tl-1223) (S. Nakajima et al., 1989), and TlBa2Ca3Cu4O11 (Tl-1234) have been determined by X-ray diffraction (Liang et al., 1988). On the other hand, for the Tl2Ba2CanlCunO2n+4 (n ¼ 1–4) system, Tl2Ba2CuO6 (Tl-2201), Tl2Ba2CaCu2O8 (Tl-2212), and Tl2Ba2Ca2Cu3O10 (Tl-2223) have been examined in detail by X-ray diffraction and neutron diffraction (Torardi et al., 1988). Although structure models of TlBa2Ca4Cu5O13 (S. Nakajima et al., 1989) and Tl2Ba2Ca3Cu4O12 (M. Kikuchi, Nakajima, et al., 1989) were proposed, their structure parameters have not been determined from diffraction analysis because they are always formed as the intergrowth with other structures. Here, the structure models of TlBa2CaCu2O7, Tl2Ba2CuO6, Tl2Ba2CaCu2O8, and Tl2Ba2Ca2Cu3O10 proposed by X-ray diffraction were examined, and the structure parameters of TlBa2Ca4Cu5O13 and Tl2Ba2Ca3Cu4O12 were determined from observed high-resolution structure images. Figs. 6.3–6.5A and B are high-resolution structure images of Tl2Ba2CuO6, Tl2Ba2CaCu2O8, TlBa2CaCu2O7, and Tl2Ba2Ca2Cu3O10, respectively, taken with the incident beam parallel to the a-axis. Arrangement of dark spots in the HREM images agrees well with those of cations in the projected structure models proposed by X-ray diffraction. The larger black spots correspond to the Tl atoms and Ba atoms, and Cu and Ca atoms with smaller atomic numbers are represented as small dark spots. Oxygen atom positions are located on bright regions between the dark spots of the cations. In addition, positions of the dark spots corresponding to cations faithfully reflect atomic coordinates of the cations, which were determined by X-ray diffraction, as listed in Table 6.2. The z coordinates of cations, measured from the observed HREM images, are listed in Table 6.2, together with values determined by diffraction methods (Herview et al., 1988; Torardi et al., 1988). The z coordinates of cations determined from the observed HREM images agree well with those determined from diffraction methods. As can be seen in Table 6.1, the dark spot positions in the structure images reflect the real positions of cations within an error of 0.01 nm, which may correspond to a measured error. This result indicates that the cation positions can be determined with the precision of 0.01 nm from the structure images. The oxygen vacant positions, which are located on the Ca layers, can be distinguished as brighter regions compared with the other bright regions corresponding to the oxygen sites. Fig. 6.5C and D are high-resolution structure images of TlBa2Ca4Cu5O13 and Tl2Ba2Ca3Cu4O12, respectively. They were observed as the intergrowth together with other layered structures, which have smaller numbers of Cu layers. The z coordinates of cations determined from the HREM images of TlBa2Ca4Cu5O13 and Tl2Ba2Ca3Cu4O12 are listed in Tables 6.3 and 6.4, respectively. Oxygen vacancies on the Ca layers can be observed in the images. In Tables 6.2 and 6.3, the

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Perovskite Ceramics

Fig. 6.6 Structure models of TlBa2Can1CunO2n+3 (n ¼ 1–5) and Tl2Ba2Can1CunO2n+4 (n ¼ 1–4). No permission required.

z coordinates of oxygen atoms in Tl-O and Cu-O layers, respectively, were assumed to be the same as those of cations, and oxygen positions in the Ba layer were assumed to be shifted by 0.07 nm along the c-axis from the results of other Tl-based superconductors (Herview et al., 1988; M. Kikuchi, Kajitani, et al., 1989; Liang et al., 1988; Torardi et al., 1988). Tl-1201 structures and other substitution-type structures had also

Crystal structures of copper oxide-based perovskite compounds

163

Table 6.2 Structural parameters (z coordinates) of cations in TlBa2CaCu2O7 and Tl2Ba2CuO6 determined from structure images of Figs.6.5A and 6.3C, and diffraction methods. TlBa2CaCu2O7

z (HREM)

z (XRD)a

Tl Ba Cu Ca

0 0.207 0.371 0.5

0 0.213 0.37 0.5

Tl Ba Cu a b

z (XRD)b

z (HREM)

Tl2Ba2CuO6

0.206 0.084 0

z (Neutron)b

0.20265 0.08301 0

0.2018 0.0822 0

Herview et al. (1988). Torardi et al. (1988).

Table 6.3 Structural parameters of TlBa2Ca4Cu5O13 determined from the high-resolution image of Fig. 6.5C. Atom

Site

x

y

z

Tl Ba Ca(1) Ca(2) Cu(1) Cu(2) Cu(3) O(1) O(2) O(3) O(4) O(5)

1d 2g 2g 2g 1c 2h 2h 4f 4i 4i 2h 1b

0.5 0 0 0 0.5 0.5 0.5 0 0 0 0.5 0

0.5 0 0 0 0.5 0.5 0.5 0.5 0.5 0.5 0.5 0

0.5 0.38 0.072 0.218 0 0.143 0.292 0 0.143 0.292 0.411 0.5

See the text about the z coordinates of oxygen sites. Space group P4/mmm. a ¼ 0.3847 nm and c ¼ 2.225 nm. TC ¼ 106 K.

been reported for the Tl-based superconductors (M. Kikuchi et al., 1994; Kunii et al., 2003; S. Nakajima et al., 1993; Ohshima et al., 1994; Yahya et al., 2007). Summarized structure models of Tl-based Cu oxide superconductor are shown in Fig. 6.6.

4

Modulated superstructures of Tl-based copper oxides

In addition to the basic layer structures of Tl-based copper oxides, modulated structures accompanied with the satellite spots have also been observed (Dmowksi et al., 1988; E. A. Hewat, Bordet, et al., 1988; A. W. Hewat, Hewat, et al., 1988; Iijima et al., 1988; T. Oku, Hiraga, et al., 1991; Parkin et al., 1988;

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Perovskite Ceramics

Table 6.4 Structural parameters of Tl2Ba2Ca3Cu4O12 determined from the high-resolution image of Fig. 6.5D. Atom

Site

x

y

z

Tl Ba Ca(1) Ca(2) Cu(1) Cu(2) O(1) O(2) O(3) O(4)

4e 4e 2b 4e 4e 4e 8g 8g 4e 4e

0 0 0 0 0 0 0 0 0 0

0 0 0 0 0 0 0.5 0.5 0 0

0.226 0.335 0.5 0.42 0.039 0.121 0.039 0.121 0.182 0.274

Space group I4/mmm. a ¼ 0.3853 nm and c ¼ 4.198 nm. TC ¼ 112 K.

H. W. Zandbergen, Groen, et al., 1988). Fig. 6.7 shows the electron diffraction patterns of Tl2Ba2CuO6 taken along the various directions of the crystal (T. Oku, 2014b; T. Oku, Hiraga, et al., 1991). In addition to the fundamental reflections, sharp satellite spots are observed in Fig. 6.7A–C, which indicate the modulated superstructure. The electron diffraction pattern of Fig. 6.7C can be obtained from Fig. 6.7D by 18 degrees rotation along the c-axis. There is no satellite spot in Fig. 6.7D, and the satellite spots along the [310] appear by the 18 degrees rotation along the c-axis as observed in Fig. 6.7C. Tl2Ba2CuO6 has both tetragonal and orthorhombic structures, and the modulated structure is observed in the orthorhombic phase. The fundamental structure with the modulated structure is distorted a little, and the indices are those of an orthorhombic unit cell (a ¼ 0.545 nm, b ¼ 0.549 nm, c ¼ 2.318 nm) as observed in the electron diffraction pattern of Fig. 6.7A. The fundamental lattice has a twin structure with a twin plane of {110} as observed in Fig. 6.7B. The superstructure reflections are observed along , and the modulated wave vector was determined as q ¼ [0.07 0.22 1] ¼ 1/6.2 , i.e., the modulation is incommensurate. Fig. 6.8A is an HREM image of Tl2Ba2CuO6 taken along the c-axis. An enlarged image of Fig. 6.8A is shown in Fig. 6.8B. In addition to the fundamental lattice fringes, dark and bright contrasts with a distance of  1.2 nm and their twin relations can be seen. This HREM images show that the direction of the modulated structure is near [130]. Although twinning of the modulated structure appears on both {110} and near {100} planes, the twinning of fundamental lattice appears only on {110} planes, as indicated by the arrows in Fig. 6.8A and B. Fig. 6.8C is a high-resolution image of Tl2Ba2CuO6 taken with the [1-30] incidence. Dark and bright contrasts with a distance of 2.4 nm are observed, and the modulated region (a) and the nonmodulated region (b) are clearly distinguishable in Fig. 6.8C. By careful observation of the modulation contrast, changes in both the darkness and position of Tl atoms can be seen, as indicated by arrows in the region (a) of Fig. 6.8C. This indicates that the origin of modulated structure would exist on the Tl-O planes.

Crystal structures of copper oxide-based perovskite compounds

165

Fig. 6.7 Electron diffraction patterns of Tl2Ba2CuO6 taken with the incident beam parallel to the (A) [001], (B) [001], (C) [1-30] and (D) [1-10] directions. (B) The electron diffraction pattern showing the modulated structure and its twinning. No permission required.

From detailed composition analysis of the Tl2Ba2CuOx phase, the orthorhombic phase with modulated structure and the tetragonal phase without the modulation had the composition of Tl1.7Ba2CuO5.7 and Tl1.6Ba2CuO5.6, respectively. The modulated structure has 0.3 oxygen and 0.3 Tl deficiencies per unit cell. The electron diffraction and high-resolution observation showed 6.2 times superstructure along the [130] direction. These results indicate that the modulation is due to the atomic ordering of oxygen and Tl in the Tl-O layers along the [130] with a period of 6.2 times. If the oxygen and Tl deficiencies are assumed along the [130] direction, the deficiencies are calculated as 0.36 per the unit cell, which agree well with the composition analysis. Therefore, the modulated superstructure is believed to be due to Tl and oxygen vacancies in the Tl-O layers along the [130] with a period of 6.2 times (2.4 nm). On the

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Perovskite Ceramics

Fig. 6.8 (A) HREM image of Tl2Ba2CuO6 taken along the [3-10] direction. (B) Enlarged image of (A). (C) HREM image of with the incident beam parallel to the [3-10] direction. No permission required.

other hand, the tetragonal phase had more atomic deficiencies randomly, and did not show the modulated structure. Electron diffraction patterns of Tl2BaSrCuO6 taken with the incident beam parallel to the [001] and [010] directions are shown in Fig. 6.9A and B, respectively. When Sr atoms are doped at the Ba sites, the fundamental structure has a tetragonal structure. Satellite reflections due to a modulated structure are observed, which are weak and diffuse compared to the Tl-2201 phase. In addition, the modulation wave vector was changed as q ¼ < 1/6 0 1>. For Tl2Ba2CaCu2O8 (Tl-2212) and Tl2Ba2Ca2Cu3O10 (Tl-2223) (Dmowksi et al., 1988; H. W. Zandbergen, Van Tendeloo, et al., 1988), weak diffuse scatterings were also observed, and the modulation wave vector is determined to be q ¼ < l/6 0 1 > as observed in Fig. 6.9C. This modulation shows two-dimensional character, and the symmetry of the fundamental lattice remains tetragonal (a ¼ 0.385 nm, c ¼ 2.92 nm). A HREM image corresponding to Fig. 6.9C is shown in Fig. 6.9D. Modulation contrast with a distance of 2.3 nm is observed in the Tl-O layer along the a-axis. Models for the modulated superstructures were reported as follows: shortrange ordering due to displacements of Tl and O in the Tl-O planes (Dmowksi et al., 1988), extra oxygen in the Tl-O planes (H. W. Zandbergen, Van Tendeloo, et al., 1988), a partial substitution of Tl3+ by Tl+ (H. W. Zandbergen, Van Tendeloo, et al., 1988), and the mutual substitution of Tl and Ca atoms (A. W. Hewat, Hewat, et al., 1988; K. Hiraga, Shindo, et al., 1988). The compositional analysis of the present samples showed that the Tl-2212 and Tl-2223 phases had compositions of Tl1.7Ba2Ca1.3Cu2O8 and Tl1.7Ba2Ca2.3Cu3O10, respectively. This implies that the excess 0.3 Ca atoms are doped at the Tl sites per the unit cell. The electron diffraction and HREM observation showed 6 times superstructure along the a-axis. If the Tl atoms are substituted by Ca atoms with a period of 6 times along the axis, the substitution atoms are 0.33 per the unit cell, which agreed well with the measured

Fig. 6.9 Electron diffraction patterns of Tl2BaSrCuO6 taken with the incident beam parallel to (A) [001] and (B) [010] directions. (C) Electron diffraction pattern and (D) HREM image of Tl2Ba2CaCu2O8 taken with the incident beam parallel to the [010] direction. Electron diffraction patterns of (E) TlBa2CaCu2O7 and (F) TlBa2Ca2Cu3O9 taken along the c-axis. No permission required.

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Perovskite Ceramics

Table 6.5 Oxygen deficiencies δ of Tl-based superconductors calculated from electron diffractions of Fig. 6.9 and iodimetric measurements of oxygen contents. Structure

Reflection at x 0 0.5

Averaged δ

Iodimetric measurement

Averaged δ

TlBa2CaCu2O7 δ TlBa2Ca2Cu3O9 δ TlBa2Ca3Cu4O11 δ

0.24–0.29 0.26–0.31 0.25–0.32

0.27 0.29 0.29

0.279–0.309 0.277–0.311 0.254–0.326

0.29 0.29 0.29

composition of the samples. Therefore, the modulated superstructure is believed to be due to the Tl substitution by Ca atoms along the a-axis with a period of 6 times (2.3 nm). Fig. 6.9E and F shows the electron diffraction patterns of TlBa2CaCu2O7 (Tl-1212) and TlBa2Ca2Cu3O9 (Tl-1223), respectively, taken along the c-axis. Weak, diffuse satellite scatterings are observed as indicated by arrows, and the observed modulation wave vector is approximately q ¼ < 0.28 0 0.5 > (T. Oku, 2014b; T. Oku, Hiraga, et al., 1991). Almost the same incommensurate diffuse scattering was observed for the TlBa2Ca3Cu4O11 (Tl-1234) phase. As the modulation shows a two-dimensional character, the symmetry of the fundamental lattice remains tetragonal. The diffuse scattering becomes stronger as the oxygen loss is increased and also TC increases. Oxygen deficiencies δ of Tl-based superconductors calculated from electron diffractions of Fig. 6.9 and iodimetric measurements of oxygen contents are summarized in Table 6.5. From the compositional analysis for the Tl-1212, 1223, and 1234 phases, the atomic ratios of Tl:Ba:Ca:Cu were determined to be 1:2:1:2, 1:2:2:3, and 1:2:3:4 (stoichiometry), respectively. However, 0.29 oxygen atoms are deficient per unit cell, as summarized in Table 6.5, which implies that the oxygen vacant positions are the same for these structures. In the present work, oxygen atoms in the Tl-O layers would be deficient, and the measured modulation from the electron diffraction patterns are summarized in Table 6.5. As listed in Table 6.5, assumed oxygen vacancies in the Tl-O layers measured by electron diffraction agreed well with the measured oxygen vacancies by iodimetric measurements. Therefore, it is believed that the modulation superstructure is due to oxygen vacancy ordering in the Tl-O layer along the a-axis with a period of 3.6 times (¼ 0.281) and 2 times along the c-axis. The period of 3.6 times is incommensurate, which implies the mixture of 3 and 4 times superstructures. In fact, modulations with periods of 3.1–4.2 times (¼ 0.321–0.241) are observed for electron diffraction patterns, as listed in Table 6.5.

5

Crystal structures of Pb-based copper oxides

Pb-based superconductors and related oxides have been discovered and investigated (Cava et al., 1988), and various new types of crystal structures of PbBaSr(Y,Ca)Cu3Oy (y ¼ 7–8.4) and Pb2(Ba,Sr)2(Ln,Ce)2Cu3Oy (Ln: Lanthanoid, y ¼ 9–10.4) (A. Tokiwa, Nagoshi, et al., 1990; A. Tokiwa et al., 1989; A. Tokiwa, Oku, et al., 1990) were

Crystal structures of copper oxide-based perovskite compounds

169

determined by high-resolution electron microscopy, with the aid of electron and X-ray diffraction, and quantitative analysis of compositions of cations and oxygen. Some results of the determination of crystal structures are shown here. Samples were prepared from a mixture of PbO, BaO2, Sr2CuO3, Y2O3, Eu2O3, CeO2, and CuO by solid-state reaction. After heating at 830°C in a 1% O2-N2 mixed gas, the specimens were slowly cooled in a furnace, quenched into liquid nitrogen or annealed in a flowing O2 gas at 400°C. Electron probe microanalysis and iodide titration method are used for quantitative analysis of the element. Fig. 6.10A and B shows the high-resolution structure images of PbBaSrYCu3O7 (Tokiwa et al., 1989) and Pb2Sr2Y0.5Ca0.5Cu3O8 taken with the incident beam parallel to the a-axis. The PbBaSrYCu3O7 shows superconductivity

Fig. 6.10 HREM images of (A) PbBaSrYCu3O7 and (B) Pb2Sr2Y0.5Ca0.5Cu3O8 taken with the incident beam parallel to the a-axis. (C, D) HREM images after crystallographic image processing of (A, B), respectively. No permission required.

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Perovskite Ceramics

at 65 K by Ca substitution for Y sites (A. Tokiwa, Nagoshi, et al., 1990), and the Pb2Sr2YCu3O8 also showed superconductivity at 75 K by Ca substitution for Y sites (Masuzawa et al., 1989). HREM images after crystallographic image processing of Fig. 6.10A and B are shown in Fig. 6.10C and D, respectively. PC, BS, YC, and Ov represent (Pb,Cu), (Ba,Sr), (Y,Ca), and oxygen vacancy, respectively, and the image directly shows an arrangement of cations. The indicated atomic arrangements were proposed from the structure image, X-ray diffraction, and energy dispersive spectroscopy (EDS) analysis. In Fig. 6.10C, zigzag arrays consisting of larger black spots correspond to (Pb,Cu) layers, and on both sides of the (Pb,Cu) layers, there are layers of a mixture of Ba and Sr atoms (A. Tokiwa et al., 1989). In this image, Cu and Y atoms, with smaller atomic numbers, are represented as small dark spots. Oxygen atom positions are located at bright regions between the cation sites of the dark spots. In addition, the oxygen vacant positions on the Y layers, indicated by Ov, can be distinguished as brighter regions from the other bright regions corresponding to the oxygen occupied sites in Fig. 6.10C and D. Darkness of the spots in the double (Pb,Cu) layers in Fig. 6.10A suggests that the zigzag double layers contain some light element, Cu in addition to Pb. A random mixture of Pb and Cu in the double Pb layers is unlikely because of their different ionic radii and coordination schemes. Careful examination of contrast for the double Pb layers in a thin part near the crystal edges, as indicated by arrows in Fig. 6.10A, suggests the pairwise distribution of Pb and Cu layers (A. Tokiwa-Yamamoto et al., 1993). Some of these spots are large and dark, while some others are small and faint, which would correspond to Pb and Cu, respectively. The observed HREM image indicates that small domains of Pb and Cu layers alternate every several unit cells. This short-range ordering seems to be two-dimensional in the Pb and Cu layers, and Pb and Cu layers are not distinguishable due to an averaging effect in a thicker part of the specimen. Atomic coordinates of PbBaSrYCu3O7 determined from the high-resolution structure image of Fig. 6.10C are shown in Table 6.6. The z coordinates of oxygen atoms in (Pb,Cu) and Cu-O layers were assumed to be the same as those of cations, and oxygen positions in (Ba, Sr) layers were assumed to be shifted by 0.03 nm along the c-axis from the results of other Pb-based oxides (Cava et al., 1988). Fig. 6.11A and B shows the electron diffraction patterns of PbBaSrY0.7Ca0.3Cu3O7 taken with the incident beam parallel to the [001] and [110] directions, respectively. Satellite reflections at 1/2 1/2 0 are observed, and weak reflections at 1/4 1/4 0 are also observed in Fig. 6.11A, which indicates a modulated superstructure with a modulation wave vector of q ¼ 1/4 , and the modulation is due to oxygen ordering. Weak streaks at 1/2 1/2 0 along c-axis are also observed in Fig. 6.11B, which indicates stacking faults of the superstructure along c-axis. The similar modulated superstructure with a modulation wave vector of q ¼ 1/4 was also observed in PbBa0.7Sr1.3EuCeCu3O9 (A. Tokiwa, Oku, et al., 1990). A high-resolution structure image of PbBa0.7Sr1.3EuCeCu3O9, which was taken with the incident beam parallel to the a-axis, is shown in Fig. 6.12A. Here, the foil thickness increases from the top to the bottom in the image. Image contrast of the thin region directly represents the projected potential. A structure model of

Crystal structures of copper oxide-based perovskite compounds

171

Table 6.6 Structural parameters of cations and oxygen atoms in PbBaSrYCu3O7, determined from the high-resolution structure image of Fig. 6.10C. Atom

Site

x

y

z

Occupancy

Pb Ba Sr Y Cu(1) Cu(2) O(1) O(2) O(3)

4e 4e 4e 2a 4e 4e 8g 4e 4e

0 0 0 0 0 0 0 0 0

0 0 0 0 0 0 0.5 0 0

0.283 0.13 0.13 0 0.283 0.438 0.062 0.359 0.217

0.5 0.5 0.5 1 0.5 1 1 1 0.5

See the text about the z coordinates of oxygen sites. Space group I4/mmm. a ¼ 0.3847 nm and c ¼ 2.748 nm.

Fig. 6.11 Electron diffraction patterns of PbBaSrY0.7Ca0.3Cu3O7 taken with the incident beam parallel to the (A) [001] and (B) [110] directions. No permission required.

PbBa0.7Sr1.3EuCeCu3O9 was determined from the HREM image by the aid of X-ray diffraction. The structure has a characteristic layer structure formed by alternate stacking of double (Eu, Ce) layers and double (Pb, Cu) layers, which are separated by (Ba, Sr) and Cu layers. In Fig. 6.12A, (Eu, Ce) atoms are represented as the largest dark spots. Fig. 6.12B is a high-resolution structure image of PbBa0.7Sr1.3YCe2Cu3O11, which was taken with the incident beam parallel to the a-axis, together with a structure model. A good correspondence between the arrangement of dark spots in the image and that of cations in the projected structure model of PbBa0.7Sr1.3YCe2Cu3O11 is clearly observed. The image shows a layer structure with three (Y, Ce) layers between Cu-O layers. The atomic coordinates were directly determined from the observed

172

Perovskite Ceramics

Fig. 6.12 HREM images of (A) PbBa0.7Sr1.3EuCeCu3O9, (B) PbBa0.7Sr1.3YCe2Cu3O11, (C) Pb2Sr2(YCe2)5/3Cu3O16, and (D) Pb2Sr2(YCe2)7/3Cu3O20 taken with the incident beam parallel to the a-axis, together with projected structure models. No permission required.

HREM image as summarized in Table 6.7. Structure models of Pb(Ba,Sr)2(Ln,Ce)nCu3O5+2n (n ¼ 1–5) determined by the present work are summarized in Fig. 6.13A. In addition to Pb(Ba,Sr)2(Ln,Ce)nCu3O5+2n (n ¼ 1–5) with double (Pb, Cu) layers, Pb2Sr2(Ln,Ce)nCu3O6+2n (n ¼ 1–7) compounds with triple (Pb, Cu) layers were investigated by high-resolution electron microscopy. Fig. 6.12C is an HREM image of the

Crystal structures of copper oxide-based perovskite compounds

173

Table 6.7 Structural parameters of cations and oxygen atoms in PbBa0.7Sr1.3YCe2Cu3O11, which were determined from the high-resolution image of Fig. 6.12B. Atom

Site

x

y

z

Occupancy

Pb Ba Sr Ce(1) Ce(2) Y(1) Y(2) Cu(1) Cu(2) O(1) O(2) O(3) O(4)

4e 4e 4e 2a 4e 2a 4e 4e 4e 8g 8g 4e 4e

0 0 0 0 0 0 0 0 0 0 0 0 0

0 0 0 0 0 0 0 0 0 0.5 0.5 0 0

0.225 0.337 0.337 0 0.434 0 0.434 0.225 0.111 0.033 0.111 0.171 0.275

0.5 0.35 0.65 0.67 0.67 0.33 0.33 0.5 1 1 1 1 0.5

Oxygen sites were assumed. Space group I4/mmm. a ¼ 0.385 nm and c ¼ 3.80 nm.

Fig. 6.13 Structure models of (A) Pb(Ba,Sr)2(Ln,Ce)nCu3O5+2n (n ¼ 3,5) and (B) Pb2Sr2(Ln, Ce)nCu3O6+2n (n ¼ 1–6) determined from high-resolution electron microscopy. No permission required.

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Perovskite Ceramics

Pb2Sr2(YCe2)5/3Cu3O16 (A. Tokiwa et al., 1991) taken with the incident beam parallel to the a-axis, together with a determined structure model, as listed in Table 6.7. This structure can be characterized as a layer structure formed by stacking of fivefold (Y, Ce) fluorite layers between (PbO-Cu-PbO) blocks. Structure models of Pb2Sr2(Ln, Ce)nCu3O6+2n (n ¼ 1–5) determined by the present work are summarized in Fig. 6.13B. In addition, single Pb-layer structures were synthesized and reported (Karimoto & Naito, 2000; H. Sasakura et al., 2014, 2009; Tamura et al., 1997; X. J. Wu et al., 1998), and new types of structures with (Ce, Y) fluorite layers were also reported (Chmaissem et al., 2010; Grigoraviciute et al., 2007). Atomic coordinates were also obtained from HREM images for Pb2Sr2(YCe2)5/3Cu3O16 and Pb2Sr2(YCe2)7/3Cu3O20, as listed in Tables 6.8 and 6.9, respectively. The observed image of Fig. 6.12A was examined in detail with the aid of computer simulations. Multislice calculations were carried out with the z coordinates obtained from the HREM image. Two types of oxygen positions (left: tetrahedral sites, right: octahedral sites) between (Eu, Ce) layers were assumed, as shown in Fig. 6.14A. The calculated image of Fig. 6.14A with oxygen positions at tetrahedral sites, under the condition of a defocus of 35 nm and a crystal thickness of 1.93 nm, agree well with the observed image of Fig. 6.12A (T. Oku, 2014a). This result indicates that the positions of oxygen atoms can be determined by comparing calculated images with observed images. The observed HREM image of Pb2Sr2(YCe2)5/3Cu3O16 in Fig. 6.12C was also investigated by computer simulations. The z coordinates obtained from the HREM image of Fig. 6.12C were used, as listed in Table 6.8, and two types of oxygen positions at tetrahedral sites (left) and octahedral

Table 6.8 Structural parameters of cations and oxygen atoms in Pb2Sr2(YCe2)5/3Cu3O16, which were determined from the high-resolution image of Fig. 6.12C. Atom

Site

x

y

z

Occupancy

Pb Sr Ce(1) Ce(2) Ce(3) Y(1) Y(2) Y(3) Cu(1) Cu(2) O(1) O(2) O(3) O(4) O(5)

2h 2g 2g 2h 1b 2g 2h 1b 1a 2h 2g 2h 4i 4i 4i

0.5 0 0 0.5 0 0 0.5 0 0 0.5 0 0.5 0 0 0

0.5 0 0 0.5 0 0 0.5 0 0 0.5 0 0.5 0.5 0.5 0.5

0.068 0.168 0.305 0.398 0.5 0.305 0.398 0.5 0 0.237 0.068 0.157 0.237 0.352 0.449

1 1 0.8 0.8 0.8 0.2 0.2 0.2 1 1 1 1 1 1 1

Oxygen sites were assumed. Space group P4/mmm. a ¼ 0.383 nm and c ¼ 2.65 nm.

Crystal structures of copper oxide-based perovskite compounds

175

Table 6.9 Structural parameters of Pb2Sr2(YCe2)7/3Cu3O20, determined from the highresolution structure image of Fig. 6.12D. Atom

Site

x

y

z

Occupancy

Pb Sr Ce(1) Ce(2) Ce(3) Ce(4) Y(1) Y(2) Y(3) Y(4) Cu(1) Cu(2) O(1) O(2) O(3) O(4) O(5) O(6)

2h 2g 2g 2h 2g 1d 2g 2h 2g 1d 1a 2h 2g 2h 4i 4i 4i 4i

0.5 0 0 0.5 0 0.5 0 0.5 0 0.5 0 0.5 0 0.5 0 0 0 0

0.5 0 0 0.5 0 0.5 0 0.5 0 0.5 0 0.5 0 0.5 0.5 0.5 0.5 0.5

0.055 0.139 0.254 0.33 0.415 0.5 0.254 0.33 0.415 0.5 0 0.193 0.055 0.13 0.193 0.292 0.373 0.458

1 1 0.667 0.667 0.667 0.667 0.333 0.333 0.333 0.333 1 1 1 1 1 1 1 1

Space group P4/mmm. a ¼ 0.383 nm and c ¼ 3.17 nm.

sites (right) between (Y, Ce) layers were assumed. The calculated image of oxygen positions at tetrahedral sites in Fig. 6.14B agrees well with the observed HREM image of Fig. 6.12B.

6

Crystal structures of Hg-based copper oxide

Hg-based superconducting oxides have the highest TC (134 K at atmospheric pressure and 150 K under high pressure) (Chu et al., 1993; Schilling et al., 1993). In addition to the Hg single-layer structure, an HgTlBa2CuOx compound with (Hg,Tl) double-mixed layers was synthesized for the first time (Nakajima et al., 1996). Although the double (Hg,Tl) layer oxide did not show superconductivity, high TC due to the control of carrier concentration can be expected by introducing doping elements. Currently, a slowscan CCD camera with high linearity and electron sensitivity is used to record HREM images and electron diffraction patterns. Although ordinary negative films have no linearity for a strong electron beam, the digital data taken by the CCD camera can be used for quantitative analyses such as residual indices and difference image calculations, and detected intensities can be directly compared with calculated data. Fig. 6.15A is an HREM image (JEM-4000EX, 400 kV) of HgTlBa2CuOx taken along the [100] direction. Darkness or sizes of dots corresponding to (Hg,Tl), Ba, and Cu positions can be identified as being nearly proportional to their atomic

Fig. 6.14 Calculated images of two models of (A) PbBa0.7Sr1.3EuCeCu3O9 and (B) Pb2Sr2(YCe2)5/3Cu3O16 with the a-axis incident, together with projected structure models. No permission required.

Crystal structures of copper oxide-based perovskite compounds

Fig. 6.15 (A) HREM image of HgTlBa2CuOx taken along the [100] direction. Blue lines indicate a unit cell. Structure models of (B) HgTlBa2CuO6 and (C) HgTlBa2CuO5. No permission required.

177

178

Perovskite Ceramics

Table 6.10 Structural parameters used for image calculation of HgTlBa2CuO5. Atom

x

y

z

Occupancy

Hg Tl Ba(1) Ba(2) Cu O(1) O(2) O(3) O(4)

0.5 0 0.5 0 0 0.5 0.5 0 0

0.5 0 0.5 0 0 0.5 0.5 0 0.5

0.299 0.201 0.087 0.413 0 0.201 0.383 0.117 0

1 1 1 1 1 1 1 1 1

Space group P4nc. a ¼ 0.3856 nm and c ¼ 2.329 nm.

numbers (T. Oku et al., 2000). Although oxygen atoms are not represented as dark spots in Fig. 6.15A, information on the oxygen positions should be included in the image (T. Oku & Nakajima, 1998a, 1999). In the (Hg,Tl) layers, oxygen positions indicated by Ov show a brighter contrast than those indicated by O. This result implies the existence of an ordering of oxygen vacancies in the (Hg,Tl) layers. Similar doublelayer structures have been reported in PbBaSrYCu3Ox and Pb(Ba,Sr)2(Eu,Ce)2Cu3Ox, and these compounds have separated double-layer structures of (Pb-O)-(Cu-Ov), as determined by HREM analysis (Ov ¼ oxygen vacancy) (A. Tokiwa, Nagoshi, et al., 1990; A. Tokiwa-Yamamoto et al., 1993). This separated double-layer structure could not be observed by X-ray powder diffraction because of an alternate atomic arrangement in the double layers, and it can be observed only by HREM analysis. From the above HREM results, structure models for HgTlBa2CuOx are constructed, as shown in Fig. 6.15B and C. A structure model for HgTlBa2CuO6 with a double (Hg,Tl) layer has been proposed in a previous work (S. Nakajima et al., 1996) and is shown in Fig. 6.15B. In the present study, a new model for HgTlBa2CuO5 with an ordering of oxygen vacancies in the Hg layers is proposed, as shown in Fig. 6.15C. In this model, the Hg and Tl layers are separated, and the Hg atoms have straight bonding with two oxygen atoms, which is a reasonable coordination for Hg2+ cations (T. Oku & Nakajima, 1998b). The metal atom positions for HgTlBa2CuO5 can be estimated from the observed HREM image as listed in Table 6.10. The oxygen atom positions were assumed and temperature factors were neglected. To investigate the observed HREM image in detail, HREM images were calculated based on the structure models for HgTlBa2CuO6 and HgTlBa2CuO5. Fig. 6.16A calculated HREM images for HgTlBa2CuO5 along the [100] direction. Image calculations were carried out for various defocus values (under defocus) and crystal thickness to determine the imaging conditions of the observed image. Contrast changes of the images are very sensitive to both defocus value and crystal thickness, and the double (Hg,Tl) layers are observed as dark dots around the defocus value of 30 nm. The Hg layers show brighter contrast compared with Tl layers. To compare quantitatively the observed image of HgTlBa2CuOx of Fig. 6.15A with the calculated

Crystal structures of copper oxide-based perovskite compounds

179

Fig. 6.16 (A) Calculated HREM images of HgTlBa2CuO5 along the [100] direction. One slice is 0.386 nm. (B) RHREM values of observed HREM images as a function of defocus values. No permission required.

images of Fig. 6.16A, residual indices (RHREM ¼ Σj Iobs  Ical j/ΣIobs) (D. Shindo et al., 1994; Smith & Eyring, 1982) were calculated, as shown in Fig. 6.16B. The minimum RHREM values of 0.153 and 0.177 were obtained at a defocus of 26 nm for HgTlBa2CuO5 and HgTlBa2CuO6, respectively. This result indicates that the stability of the double (Hg,Tl) layer structure is due to the formation of oxygen vacancies in the Hg layers. The present result also indicates that the local atomic arrangement with light elements such as oxygen can be determined by a combination of HREM and RHREM values. In this study, a residual index between the observed and the calculated images was used because of the simple form and the usefulness in the determination of defocus values and crystal thickness. The crystal structure of high-TC superconductor of Tl2Ba2CuO6 was investigated by these RHREM values, and the occupancy of Tl was estimated after the determination of the crystal thickness and defocus values of the observed images (D. Shindo et al., 1994). To analyze more clearly, through-focus imaging would be useful. In addition to the present RHREM values, several methods for the estimation of image agreement between experimental and simulated images exist, such as fractional mean absolute difference (Smith & Eyring, 1982), cross-correlation function (M€obus & R€uhle, 1994), mean relative difference (M€obus & R€ uhle, 1994), normalized Euclidean distance (Hofmann & Ernst, 1994), and nonlinear least-squares methods (King & Campbell, 1994). These are the methods useful for structural evaluation by image matching. In the present work, the observed image was fixed after image processing in order to get information on the effects of thickness in the calculated images. It is believed that more accurate

180

Perovskite Ceramics

atomic positions can be determined by combining these RHREM values with the observed HREM images recorded under optimum experimental conditions with quantitative devices.

7

Crystal structures of Bi-based copper oxide

Bi-based copper oxides with Ag are considered for wire application (Gencer et al., 2006; G€ om€ ory et al., 2002). A spray-dried aqueous solution of nitrates with an atomic ratio of Bi1.5Pb0.5Sr2Ca2Cu3 was calcined at 650°C for 15 h, resulting in a mixture of oxides (E. Bruneel, Oku, Degrieck, et al., 2001; E. Bruneel, Oku, Penneman, et al., 2001; E. Bruneel et al., 2004). The obtained precursor powder with a grain size of 3 μm was mixed with 30 vol% Ag whiskers with a diameter of 20–50 μm and a length of a few hundred micrometers. The Ag whiskers were synthesized via an electrochemical reduction of a Ag nitrate solution by a copper wire at pH 2. The mixture was pressed into bars and sintered at 853°C for 170 h in air to obtain the Bi-2223/Ag composites. To observe the atomic arrangements of the Bi1.5Pb0.5Sr2Ca2Cu3O10 (Bi-2223) phase more clearly, crystallographic image processing was carried out using the Fourier transform of the HREM image. Fig. 6.17 is an HREM image after crystallographic image processing of Bi1.5Pb0.5Sr2Ca2Cu3O10, together with a projected structure model and a simulated image enclosed by a white square. The HREM image clearly shows the metal atom arrangements in the crystal, and the blue lines indicate the unit cell. The darkness and sizes of the black spots corresponding to (Bi, Pb), Sr, Cu, and Ca positions can be identified as being nearly proportional to their atomic numbers. Although oxygen atoms are not represented as dark spots in Fig. 6.17, information on the oxygen positions should be included in the image (T. Oku, 2012, 2014b). In the Ca layers, oxygen positions indicated by Ov show a contrast brighter than the other white regions, which indicates the existence of an ordering of oxygen vacancies in the (Bi, Pb) layers. A structure model of Bi-2223 determined by X-ray diffraction (Kijima et al., 1989) is inserted into Fig. 6.17, which agrees well with the arrangements of the black dots in the image. To investigate the observed HREM image in detail, HREM images were calculated based on the structure model of Bi-2223. Fig. 6.18 shows calculated HREM images for Bi-2223 along the [110] direction. Image calculations were carried out for various defocus values (under defocus) and crystal thickness to determine the imaging conditions of the observed images. The contrast changes of the images are very sensitive to both defocus value and crystal thickness. The double (Bi, Pb) layers are observed as dark dots around the defocus value of 30 nm. A simulation image calculated at the defocus value of 30 nm and crystal thickness of 2.7 nm is inserted into Fig. 6.17, which agrees well with the observed HREM image, and structure models of BiSr2Can1CunO2n+4 (n ¼ 1–4) are shown in Fig. 6.19. Fig. 6.20A is a TEM image of the Bi-2223/Ag whisker interface in a sintered composite. A thin layer with a different contrast is observed at the Bi-2223/Ag interface. Electron diffractions of Ag and Bi-2223 phase are also shown in Fig. 6.20B and

Crystal structures of copper oxide-based perovskite compounds

181

Fig. 6.17 HREM image after crystallographic image processing of Bi1.5Pb0.5Sr2Ca2Cu3O10, together with a projected structure model and a simulated image enclosed by a white square. No permission required.

C, respectively. Fig. 6.20B is observed along the [1-10] of the face-centered cubic Ag crystal. The diffraction pattern of Bi-2223 phase in Fig. 6.20C was taken along the a-axis, which indicates a modulated structure with a modulation wave vector of q  < 0 1/4 0>. The origin of the modulated superstructure would be metal atom displacements in the crystal (Eibl, 1991; K. Hiraga et al., 1989; Hirotsu et al., 1988; Kaneko et al., 2008; Matsui & Horiuchi, 1988; D. Shindo et al., 1988). Fig. 6.21A is a TEM image of the Bi-2223/Ag whisker composite with Ag-rich phase. Lattice fringes of c-planes of Bi-2223 phase are observed, and an amorphous/nanocrystalline (AM-NC) structure is observed at the Ag/Bi-2223 interface. The white areas are the result of preferential ion milling, probably of amorphous phases, which are more easily removed compared to Bi-2223. The superconducting phase is oriented with the c-axis perpendicular to the interface. An EDX spectrum of the AM-NC phase in Fig. 6.21A is shown in Fig. 6.21B, which indicates a composition of Ag2Bi2.4Pb0.6Sr2Ca0.8Cu4.3, and an increase of Ag, Bi, Pb, and Cu concentrations in the amorphous phase. An enlarged HREM image and an electron diffraction pattern at the interface are shown in Fig. 6.21C and D, respectively.

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Fig. 6.18 Calculated image of Bi-2223 as a function of crystal thickness and defocus values of the TEM objective lens. No permission required.

The Bi-2223/AM-NC interface exhibits small steps of half or one unit cell of Bi-2223, and such intermediary phase was also observed (Kova´c et al., 1997). The diffraction pattern exhibits [110] incident of the Bi-2223 crystal, and the observed streak along the c-axis indicates that there is a small amount of Bi-2212 or Bi-2234 phase to form an intergrowth structure. A diffuse ring is also observed, as indicated by arrows, which exhibits the amorphous-like structure of AM-NC phase. The influence of Ag on the yield in Bi-2223 synthesis can be explained in terms of the shift in the incongruent melting point.

8

Structures of lanthanoid-based copper oxides

Various types of lanthanoid-based copper oxides have been reported (Akimitsu et al., 1988; Takayama-Muromachi et al., 1988; Varela et al., 1997), and electrondoped Nd2 xCexCuO4 superconductors were discovered (Izumi et al., 1989;

Crystal structures of copper oxide-based perovskite compounds

Fig. 6.19 Structure models of BiSr2Can1CunO2n+4 (n ¼ 1–4). No permission required.

Fig. 6.20 (A) TEM image of (Bi,Pb)2Sr2Ca2Cu3Ox/Ag whisker interface in a sintered composite. Electron diffraction of (B) Ag and (C) (Bi,Pb)2Sr2Ca2Cu3Ox. No permission required.

183

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Fig. 6.21 (A) TEM image of the Bi-2223/Ag whisker composite with Ag-rich phase. (B) EDX spectrum of the AM-NC phase in (A). (C) HREM and (D) electron diffraction pattern at the interface. No permission required.

Takagi et al., 1989; Y. Tokura, Takagi, & Uchida, 1989). To clarify the microstructures, single crystals of Ln2CuO4 prepared with various heat treatments are investigated by means of high-resolution electron microscopy and electron diffraction. Single crystals of Ln2CuO4 (Ln ¼ Pr, Nd, Sm) were grown by the traveling-solventfloating-zone technique using an infrared-heating furnace (T. Kajitani et al., 1990; T. Oku, Kajitani, et al., 1991). To reduce oxygen content, parts of the Ln2CuO4 (Ln ¼ Pr, Nd, Sm) samples were annealed at 1100°C for 18 h in air and quenched in liquid nitrogen. The rest of the samples was annealed at 400°C in air for 38–140 h to saturate oxygen content in the crystals. SmLa0.75Sr0.25CuO4 was also synthesized from a mixture of La2O3, Sm2O3, CuO, and SrCO3 (Y. Tokura, Takagi, Watabe, et al., 1989). Mixed powder was first calcined at 950°C in air for 10 h, then pressed into pellets, and finally sintered at 1130°C in air for 15 h. The pellets were quenched to room temperature in air and subsequently annealed at 550°C in the atmosphere with various oxygen pressures. A high-resolution image of Sm2CuO4 taken with the [100] incidence is shown in Fig. 6.22A. The specimen thickness increases from the top to the bottom in this HREM image. The observed image shows the projected arrangement of metal atoms. There

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Fig. 6.22 HREM images of (A) Sm2CuO4 and (B) SmLa0.75Sr0.25CuO3.95 taken along the a-axis. No permission required.

are two kinds of dark spots in the image. The larger ones in the zigzag arrays correspond to the Sm atoms. Cu atoms, with a smaller atomic number, are represented as less dark spots between the double Sm layers. Fig. 6.22B is an HREM image of SmLa0.75Sr0.25CuO4 taken along the a-axis. Fairly large separation between two lines of La ions is clearly observed in the HREM image, which is in accord with X-ray diffraction analysis (Y. Tokura, Takagi, Watabe, et al., 1989). Based on the crystal structure models of Fig. 6.23A and B, image calculations on Sm2CuO4and SmLa0.75Sr0.25CuO4 were carried out to confirm the structures, as shown in Fig. 6.23C and D, respectively. The images of Fig. 6.23C and D calculated at a Scherzer defocus of 45 nm and a crystal thickness of 1.2–2.7 nm, agree well with the observed image contrast for both thin and thick regions. For the SmLa0.75Sr0.25CuO4 structure, the difference of oxygen atom positions in the Sm and La-Sr layers can be clearly observed both in the observed and in the calculated image, which indicates that the HREM image includes information on both metal and oxygen atoms in the crystal. Modulated superstructures are also observed in the lanthanoid-based copper oxides (Chen et al., 1989; Li et al., 1990; van Aken & M€uller, 1991; Williams et al., 1989). The domains of superlattice 5–60 nm in diameter were observed, and the smaller domains (5–10 nm) are observed around a large one, and seem to grow into larger ones (40–60 nm). A representative superstructure domain in Nd2CuO4 is shown in Fig. 6.24A, and Fig. 6.24B is an electron diffraction pattern of Fig. 6.24A. Sharp satellite reflections with a wave vector q ¼ < 1/4 1/4 0> are observed in Fig. 6.24B. In Fig. 6.24A, repeated dark and bright contrasts separated at a distance of 1.1 nm (’ 2 √ 2  a) are observed in the [110] direction. The high-resolution image and the diffraction pattern reveal that the basic lattice spacing of the superlattice lengthen as much as 101.5% and 100.3% in the [110] and [1-10] directions, respectively, as

Fig. 6.23 Structure models of (A) Sm2CuO4 and (B) SmLa0.75Sr0.25CuO3.95. Calculated images of (C) Sm2CuO4 and (D) SmLa0.75Sr0.25CuO3.95. No permission required.

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Fig. 6.24 (A) HREM image and (B) electron diffraction pattern of a single domain of the superlattice in Nd2CuO4, taken with the [001] incidence. (C) HREM image and (D) electron diffraction pattern of superlattice domains along two directions. No permission required.

compared with the fundamental lattice. Therefore, the contrast due to strain field is observed around the domain. Two-directional superlattice domain is also observed as shown in an HREM image in Fig. 6.24C, and an electron diffraction pattern of the superlattice domains along two directions is shown in Fig. 6.24D, which also indicates a modulation wave vector of q ¼ < 1/4 1/4 0>. It can be considered that such domain structure is due to nonuniformity of oxygen content in the specimens. Various types of modulated superstructures were observed in Ln2CuO4, as listed in Table 6.11. Fig. 6.25A–D show the electron diffraction patterns of Nd2CuO4, Pr2CuO4, Pr1.85Ce0.15CuO4, and Sm2CuO4, taken along the [1-10] direction. For the Pr2CuO4, Pr1.85Ce0.15CuO4, and Sm2CuO4 crystals, satellite reflections at 0.24 0.24 0.72, 1/3 1/3 0, and 1/2 1/2 1 are observed, as shown in Fig. 6.25B–D, respectively. Fig. 6.25E and F show the electron diffraction patterns of Pr2CuO4 taken along the [010] and [111] directions, respectively, which also indicates weak diffuse scattering and sharp satellite reflections at 1/2 0 1/2 and 1/4 1/4 1/2, respectively. A superstructure has been observed and characterized for Nd2 xCexCuO4 by a wave vector q ¼ < 1/4 1/4 0 > (T. Oku, 2014a), and it was suggested that the modulation is due to ordering of oxygen vacancy and/or Ce. However, the satellite reflections were

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Table 6.11 Summary of modulated structures of Ln2CuO4. Structure

Modulation wave vector q

Nd2CuO4 Sm2CuO4 Pr2CuO4 Nd1.85Ce0.15CuO4 Sm1.85Ce0.15CuO4 Pr1.85Ce0.15CuO4

< 1/4 1/4 0 >, < 1/2 1/2 1 > < 1/4.2 1/4.2 3/4.2 >, < 1/4 1/4 1/2>, < 1/2 0 1/2 > < 1/4 1/4 0 > < 1/4 1/4 0 >, < 1/2 1/2 1 > < 1/3 1/3 0 >

not observed in the diffraction patterns of Nd2CuO4 and Pr2CuO4 quenched from 1100°C. The result indicates that the appearance of superstructures is sensitive to the oxygen content and unrelated to ordering of Ce atoms. Since neutron diffraction study shows the deficiency of oxygen in Cu-O planes, it can be supposed that the superlattices are due to ordering of oxygen atoms in the Cu-O planes.

9

Oxygen ordering in YBa2Cu3O72x

The crystal structure of YBa2Cu3O7 is based on a triple perovskite structure and is characterized by the ordering of oxygen vacancies (T. Oku, 2014a), that is, the oxygen positions on the Y atom layer and between two Ba atoms are vacant. In addition, oxygen orderings in the Cu-O basal planes were observed (K. Hiraga et al., 1987). Bulk samples of YBa2Cu3O7 x superconductors were prepared by mixing BaCO3, Y2O3 and CuO powders with the composition of YBa2Cu3O7 x phase. The mixture pellets were calcined at 930°C for 12 h in air and then cooled slowly in a furnace. After crushing the pellets to form powders, the process was repeated once more. The obtained pellets were reheated at 500–900°C and quenched into liquid nitrogen, and subsequently annealed at 500–300°C in the vacuum seal. The pellets were also annealed in a flowing N2 gas at various temperatures to control the oxygen contents. The oxygen contents were investigated by iodimetric measurements and mass change. As the consequence of the tetragonal-to-orthorhombic phase transition of YBa2Cu3O7 x at 600°C (Eatough et al., 1987; T. Kajitani et al., 1987), twin boundaries are often observed. Fig. 6.26A and B show the TEM image and lattice image of YBa2Cu3O7 x taken with the incident beam parallel to the c-axis. Twin boundaries (TBs) are coherent and indicated by arrows. The twin boundaries show distinct contrast in the TEM image, and the existence of boundaries is evident from kinks of the lattice fringes. In the oxygen-deficient YBa2Cu3O7 x compounds, oxygen vacancy ordering was observed (Beyers et al., 1989; De Fontaine et al., 1990). Fig. 6.26C and D show the electron diffraction patterns of YBa2Cu3O6.68 taken with the incident beam parallel to the c-axis and b-axis, respectively. The electron diffraction pattern shows orthorhombic structure with a- and b-axis and a twin structure with a {110} twin plane. Both electron diffraction patterns in Figs. 6.26C and D show diffuse satellite reflections at 1/2 0 0 along the a-axis. This indicates the existence of modulated

Fig. 6.25 Electron diffraction patterns of (A) Nd2CuO4, (B) Pr2CuO4, (C) Pr1.85Ce0.15CuO4, and (D) Sm2CuO4, taken with the [11(_)0] incidence. Electron diffraction patterns of Pr2CuO4 taken along the (E) [010] and (F) [1(_)11] directions. No permission required.

Fig. 6.26 (A) TEM image and (B) lattice image of YBa2Cu3O7 x taken with the incident beam parallel to the c-axis. Twin boundaries (TB) are indicated by arrows. Electron diffraction patterns of YBa2Cu3O6.68 taken with the incident beam parallel to the (C) c-axis and (D) b-axis. (E) Fourier transform of (B). (F) Filtered inverse Fourier transform of (E). No permission required.

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191

superstructure with a modulation wave vector q ¼ < 1/2 0 0>, which is be due to ordering of oxygen vacancies on the basal Cu-O planes. Fig. 6.26E is a Fourier transform of HREM image of Fig. 6.26B, and filtered inverse Fourier transform of Fig. 6.26E is shown in Fig. 6.26F, in which linear bright stripes with the distance of 2a are observed along a-axis. The twin boundary can be clearly seen at a glancing view parallel to the a- or b-axis in Fig. 6.26F. Electron diffraction patterns of oxygen-deficient YBa2Cu3O6.80, YBa2Cu3O6.47, YBa2Cu3O6.23, and YBa2Cu3O6.29 taken along the c-axis are shown in Fig. 6.27A– D, respectively. In Fig. 6.27A, weak diffuse streaks are observed along the a-axis, which indicate short-range ordering of oxygen atoms. Fig. 6.27B and C show superstructures with a modulation wave vector q ¼ < 1/3 0 0> and < 0 1/3 0 >, which indicates that the superstructures are formed along a-axis and b-axis of orthorhombic cell, respectively. In addition, a superstructure with a modulation wave vector q ¼ < 1/4 0 0 > is observed along the a-axis, as shown in Fig. 6.27D. These modulated structures are summarized in Table 6.12. Fig. 6.27E and F show a lattice image and an electron diffraction pattern of YBa2Cu3O6.47taken with the [001] incidence. Satellite peaks with a modulation wave vector q ¼ < 1/3 0 0> are observed together with q ¼ < 1/2 0 0 > in the diffraction pattern. Linear bright stripes with the distance of 3a are observed along the two principal lattice directions in the HREM image of Fig. 6.27E. From these observations, a model for the ordered arrangement of oxygen vacancies is proposed, as shown in Fig. 6.28. Basic structure models of YBa2Cu3O7 and YBa2Cu3O6 are shown in Fig. 6.28A and B, respectively. By combining these two models, an oxygen ordered YBa2Cu3O6.5 model was proposed as shown in Fig. 6.28C. The fundamental unit cell is orthorhombic, with a dimension of 2a  b  c. Other oxygen-ordering models of basal planes (Cu-O) of the YBa2Cu3O7 x are also proposed as shown in Fig. 6.28B, which depends on the oxygen content. These phases would correspond to the ortho-II phase and ortho-III phases (Iliev et al., 1993; Manca et al., 2001; Plakhty et al., 1992; Schleger et al., 1995; Stratilatov et al., 1993; Wille et al., 1988; Yasuoka et al., 1989), which results in changes in TC, and the control of oxygen atoms in the oxide crystals is important (Wille et al., 1988; Yasuoka et al., 1989; Yoshinari et al., 1990).

10

Y-based copper oxides with high JC

Superconducting thin films can be used for electronic devices and other applications (Mikheenko et al., 2011; Palonen et al., 2013; Paturi et al., 2008; P. Zhao et al., 2011, 2013). Chemically vapor deposited (CVD) YBa2Cu3O7 thin films with high critical current density (JC) of 6.5  104 A/cm2 at 77.3 K and 27 T were investigated (Yamane et al., 1989, 1988). The films were deposited on the SrTiO3(100) substrate. β-Diketonate complexes (2, 2, 6, 6-tetramethy 1-3, 5-heptanedionate chelated Y3+, Ba2+ and Cu2+) were used as three vapor sources (H. Yamane et al., 1993). The films were deposited in a mixed gas atmosphere or Ar/O2 ¼ 3/1 at 10 Torr on the substrate which was heated at 900 °C close to the melting point of YBa2Cu3O7 y.

Fig. 6.27 Electron diffraction patterns of (A) YBa2Cu3O6.80, (B) YBa2Cu3O6.47, (C) YBa2Cu3O6.23, and (D) YBa2Cu3O6.29 taken along the c-axis. (E) HREM lattice image and (F) electron diffraction pattern of YBa2Cu3O6.47 taken along the c-axis. No permission required.

Table 6.12 Summary of modulated structures of YBa2Cu3Oy. Quenched and annealed in sealed tube

Annealed in N2

y

TC/K

Observed structure

y

Observed structure

6.8 6.68 6.6 6.47 6.29

86 49 48 4 –

Diffuse streaks along a-axis

Diffuse streaks along a-axis < 1/3 0 0 >, < 1/2 0 0 > < 1/3 0 0 >, < 1/4 0 0 >

6.91 6.74 6.5 6.23 6

Perfect orthorhombic < 0 1/3 0>, < 1/2 0 0> < 1/2 0 0> < 1/2 0 0>, < 0 1/3 0> Perfect tetragonal

Fig. 6.28 Models for (A) YBa2Cu3O7 and (B) YBa2Cu3O6. (C) Two times periodic model of ordered arrangement of oxygen vacancies in YBa2Cu3O6.5 along a-axis. (D) Oxygen-ordering models of basal planes (Cu-O) of the YBa2Cu3O7 x. No permission required.

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The postannealing was carried out in the O2 atmosphere at 1 atm. To view the cross section of the CVD film, two pieces of cut substrate and film were pasted together, film to film, then cut again perpendicular to the film and polished in an Ar ionsputtering mill. Fig. 6.29A is a low magnification image obtained with the incident beam perpendicular to the substrate. Relatively large grains 20–100 nm are observed in the image. Many observed precipitates showed Moire fringes, which indicates that the precipitates are semicoherent with the matrix and have very similar crystal lattice. A cross-sectional high-resolution image is shown in Fig. 6.29B. Relatively large precipitates 30 nm in size are embedded parallel to the c-planes. These precipitates could act as the flux-pinning centers, which would result in high JC. Many defect regions are

Fig. 6.29 (A) Low magnification image of CVD-YBa2Cu3O7 film taken with the electron beam perpendicular to the substrate. Round precipitates with Moire fringes on them are observed. (B) Cross-sectional HREM image of CVD-YBa2Cu3O7 film taken with the electron beam parallel to the a-axis. (C) Stacking faults in the film. HREM image of (D) a grain and (E) grain boundary of the film. (F) Structure model and (G) structure image of YBa2Cu3O7 together with a projected structure model. No permission required.

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formed parallel to c-planes, as shown in Fig. 6.29C. The defect regions consist of extra layers 20 nm in diameter with deformed regions adjacent to them. Those faulted regions may also act as the flux-pinning centers. Fig. 6.29D is a cross-sectional HREM image of a grain in the CVD film, and most of the grains have an preferred c-axis orientation perpendicular to the SrTiO3(100) substrate, as indicated by region 1. On the other hands, the c-axis is oriented parallel to the SrTiO3(100) substrate in region 2. An enlarged HREM image at the grain boundary interface is shown in Fig. 6.29E, and c-axis of the YBa2Cu3O7 structure is perpendicular to one another. A structure model and a structure image of YBa2Cu3O7 are shown in Fig. 6.29F and G, respectively. Undeformed YBa2Cu3O7 x domains, interleaved by the faulted regions, are 50 nm in size. The electron diffraction pattern taken with the incident beam parallel to [010] did not change when the specimen was tilted by 5 degrees, which is appreciably wider than the usual crystals. Such morphology indicates that the sample was grown quickly at a temperature just below the melting point or YBa2Cu3O7 x where the homogeneous nucleation process followed by the preferential crystal growth in the c-planes occurs. Bulk types of YBa2Cu3O7 x superconductors with high JC were also investigated. Oxide powders or Y, Ba, and Cu were mixed with the composition of YBa2Cu3O7 x phase. The mixture was calcined at 900°C for 24 h and then reground. After this process was repeated twice, a small amount of additive oxide powder with a perovskite structure BaZrO3 was added, 3 mol%, to the YBa2Cu3O7 x powder (K. Osamura et al., 1994). Again, grinding and calcining was repeated, and the pellets were sintered at 950°C for 24 h and then cooled slowly. The bulk sample showed high critical current density of 2.0  105 A/cm2 at 77.3 K and 0.1 T. Fig. 6.30A is a TEM image of an YBa2Cu3O7 x bulk specimen with BaZrO3 additive, where two types of fine particles with spherical and irregular shapes

Fig. 6.30 (A) TEM image of 3 mol% BaZrO3-dispersed YBa2Cu3O6+x. (B) HREM image at the BaZrO3 nanoparticles and YBa2Cu3O6+x matrix interface. No permission required.

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are lying in the YBa2Cu3O7 x matrix. Their minimum size was 20 nm and the average size was roughly estimated to be 100 nm, and they were confirmed to be BaZrO3 phase from electron diffraction. Fig. 6.30B is an HREM lattice image near the grain boundary between the BaZrO3 phase and the YBa2Cu3O7 x matrix, where both lattice planes generating a lattice fringe are assigned as indicated in the figure. The lattice plane continues across the interface, but the extra half plane is inserted in every 20 fringes within the BaZrO3 phase, while the interval for the extra half plane is theoretically suggested to be every 14 fringes. Therefore, the interface in this region was suggested to be semicoherent. The mechanism of fine dispersion of ABO3 oxides in the matrix is not yet fully understood. It is suggested that the grain growth takes place during sintering and ABO3 phase is dragged into the growing grain as the same species as the matrix, because both phases have a similar crystal structure and have a semicoherent interface as mentioned above. Several types of pinning centers like Y2BaCuO5 (Matsushita et al., 1991), CuO (Watanabe et al., 1990), SnO2 (K. Osamura et al., 1990), ZrO2 (Oka et al., 1992), and irradiation-induced defects (Civale et al., 1991) have been reported to be effective for the YBa2Cu3O7 x superconductor (Gutierrez et al., 2007; Koshelev & Kolton, 2011; Macmanus-Driscoll, 2010; Maiorov et al., 2009; Miura et al., 2011). A perovskite-type oxide ABO3 seems to be a promising candidate as pinning center for the superconducting copper oxides, because the crystal structure is similar to the host superconducting phase. A remarkable flux-pinning effect could be expected by introducing the oxide. Homogeneous distribution of nonsuperconducting fine particles is very effective, when they distribute in the same dimension with the fluxoid lattice. In addition, the matrix is distorted around the particles with different unit volume by the coherent interface with the matrix, and a significant change of the superconducting property is expected around the coherent particles.

11

CO3- and BO3-based copper oxides

Two interesting structures have been prepared and analyzed. Sr2CuO2(CO3) was prepared at 1273 K and 0.01 MPa CO2 partial pressure in a flowing gas of O2/CO2 using a mixture of SrCO3 and CuO powders as a starting material (Miyazaki et al., 1992). The compound has a tetragonal structure with lattice constants a ¼ 7.8045 and ˚ , and the structure model is shown in Fig. 6.31A. The formula per unit c ¼ 14.993 A cell is 8 Sr2CuO2(CO3). The structure consists of deformed [CuO6] octahedrons and layers of ordered triangular CO3 groups. Sr atoms having eight nearest-neighbor oxygen are between [CuO6] octahedrons and the CO3 layers. Superconductivity was also observed for the CO3-based copper oxides (Feenstra et al., 1995; M. Uehara et al., 1993). Another new oxyborate Nd2Sr3Cu3O6(BO3)2 was also prepared at 1030°C and 1 atm of O2 (Amamoto et al., 1994), and examined by means of electron diffraction, HREM, and Rietveld analysis for the X-ray diffraction data. This compound crystal˚ . The crystal structure of lized in the tetragonal system with a ¼ 7.767 and c ¼ 35.828 A Nd2Sr3Cu3O6(BO3)2 was explained as an alternative stacking of (BO3)(Sr, Nd)2CuO2 (B-1201 unit) layers and (BO3)Sr2NdCu2O4 (B-1212 unit) layers, as shown

Crystal structures of copper oxide-based perovskite compounds

197

Fig. 6.31 Structure models of (A) Sr2CuO2(CO3) and (B) Nd2Sr3Cu3O6(BO3)2. No permission required.

in Fig. 6.31B. The BO3 groups, having a triangular planar structure, were ordered between the CuO2 sheets. Superconductivity was also observed for the BO3-based copper oxides (Sato et al., 1996; M. Uehara et al., 1994).

12

Defects, interfaces, and surface structures

As mentioned in the previous sections, the perovskite-type superconducting copper oxides have many types of layer structures with slightly different compositions. Thus, their crystals always include high-density intergrowth with various types of structures, in addition to well-ordered regions. A typical example of such disordered region is

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shown in Fig. 6.32A, which is a structure image of the intergrowth of single and double Tl layers and various numbers of Cu-O layers (T. Oku, 2014a). The upper and lower rectangles correspond to TlBa2Ca3Cu4O11 and Tl2Ba2Ca3Cu4O12, respectively. Fig. 6.32A shows three-, four-, and fivefold Cu sequences between the Ba layers. In addition to the intergrowth of different numbers of Cu-O layers, an intergrowth structure with different numbers of Tl-O layers was observed in Fig. 6.32A. The various intergrowth structures give rise to different transition temperatures TC, and thus, steps or tails were observed in their resistivity-temperature curve (Nakajima et al., 1989). The intergrowth models are shown in Fig. 6.32B. Another typical example of such disordered regions is shown in Fig. 6.32C, which is a one-dimensional lattice image of TlBa2Ca3Cu4O11 taken with the incident beam perpendicular to the c-axis. Although the image does not represent each atom, stacking sequence of the layered structure can be distinguished, that is, thick black lines correspond to Tl and Ba layers, and thin white lines correspond to Ca layers with oxygen vacancy. In Fig. 6.32C, two-, three-, five-, and sixfold Cu sequences (white lines) between Tl and Ba layers (thick black lines) are observed in addition to the usual fourfold Cu sequences. Since the structures with different numbers of Cu-O layers show different transition temperatures, this intergrowth structure caused multiple transition temperatures near TC (S. Nakajima et al., 1989), because the structures with different numbers of Cu-O layers show different transition temperatures. This defect is very sensitive to annealing condition of the sample. The lattice images such as in Fig. 6.32C can be easily observed and have enough information on the intergrowth of various layer structures. Fig. 6.32D is also an HREM image of TlBa2Ca3Cu4O11, which has a higher resolution than that of Fig. 6.32C. Although the HREM image is a two-dimensional image, it is not a structure image. It is interesting that the periodic intergrowth forms a new ordered structure with a long period. It would be difficult to synthesize this type of crystal with a single phase. Tl atoms in the Tl-based superconductors are easily vapored at high temperatures, and Tl content decreases when annealed at high temperatures. Fig. 6.33 shows the onedimensional lattice images of Tl-vaporized regions. Starting composition of the sample was Tl:Ba:Ca:Cu ¼ 2:2:2:3, and was sintered at 890°C for 10 h. In the lattice image in Fig. 6.33A, black strain contrasts are observed, which resulted from lattice deformation around the terminations of T1-O layers, being about to disappear by the vaporization of Tl, as indicated by small arrows. In Fig. 6.33B, the vaporization of Tl progressed furthermore. The regions of Cu-O layers were expanded, and the strain field contrasts become stronger. After Tl-O layers disappear, the remaining Ba-O layers are observed, as indicated by arrows in Fig. 6.33B. Similar structures were often observed in the Pb-based superconductors. An intergrowth of (Pb,Cu) layers is observed in PbBaSrYCu3O7 (Pb-2212), as shown in Fig. 6.34A. In the image, the intergrowth of two types of units, (a) and (b), which have double (Pb,Cu) layers and stacks of two (Pb,Cu) and one Cu layers is observed. A high-density intergrowth in PbBa0.7Sr1.3YCe3Cu3O13 is also shown in Fig. 6.34B. In this image, the intergrowth of various numbers of (Y,Ce) layers is observed, and disappearance of (Pb,Cu) layers is also observed as indicated by arrows, and extended regions of the (Y,Ce) layers can be seen. Observations of the intergrowth give us

Fig. 6.32 (A) Structure image of intergrowth of single and double Tl layers and various numbers of Cu-O layers. Upper and lower rectangles correspond to TlBa2Ca3Cu4O11 and Tl2Ba2Ca3Cu4O12, respectively. (B) Structure models of intergrowth. High-resolution lattice images of TlBa2Ca3Cu4O11 taken with the incident beam (C) perpendicular to the c-axis and (D) parallel to the a-axis. No permission required.

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Fig. 6.33 (A) High-resolution lattice images taken with the incident beam perpendicular to the c-axis. (B) More Tl-vaporized region. No permission required.

Fig. 6.34 HREM images of (A) PbBaSrYCu3O7 and (B) PbBa0.7Sr1.3YCe3Cu3O13 taken with the incident beam parallel to the a-axis. Number of (Y,Ce) layers are shown in (B). No permission required.

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variable information on the possibility of appearance of new unknown structures and on the transitional structures to the equilibrium state. Lattice defects particularly dislocations in the superconductors having complex structures are an interesting subject to study mechanical and electrical properties of these materials. Fig. 6.35A is an end-on view of a dislocation observed in Tl2Ba2CuO6. The dislocation is laid along the [1-10] direction parallel to the incident beam direction. An extra plane can be seen by obliquely viewing Fig. 6.35A along the vertical direction. Fig. 6.35B is a Fourier transform of Fig. 6.35A, and Fig. 6.35C is an image reconstructed by using 110 and -1-10 reflection in the Fourier diffractogram of Fig. 6.35B. Fig. 6.35D is an enlarged image of Fig. 6.35C. In the image of Fig. 6.35D,

Fig. 6.35 (A) High-resolution image of an end-on view of a dislocation in Tl2Ba2CuO6, taken with the incident beam parallel to the [11(_)0] direction. (B) Fourier transform of (A). (C) Image reconstructed by using 110 and -1-10 reflection in the Fourier diffractogram of (A). (D) Enlarged image of (C). No permission required.

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a long-range strain field over 4 nm in diameter is observed around the dislocation core. Although a Burgers vector of this dislocation cannot be determined only from this image, it would be [100] of the perovskite unit. The smallest Burgers vector in the perovskite unit is 1/2 [111] position, but the dislocation with the [100] Burgers vector is possibly stable in this structure, because a different type of atom is placed at the 1/2 [111] position. This dislocation could be considered to have formed during crystal growth, and it would be placed in the most stable state. Understanding of interface and surface structures is important for application of the superconductors to electronic devices such as Josephson junction and superconducting transistor (Anders et al., 2010; Fenton et al., 2007; Mans et al., 2006; Martens et al., 1991; Scherbel et al., 2004), as illustrated in Fig. 6.36A and B, respectively. Recently, the surface structure has been widely studied with scanning tunneling microscopy in the atomic scale. High-resolution electron microscopy has also made possible to observe atomic structures of interfaces and surfaces in the superconductors. Fig. 6.36C is a one-dimensional lattice image of an antiphase boundary in Tl2Ba2CuO6, which was observed perpendicular to the c-axis. Tl layers are indicated by arrows, and the phase displacement is 0.2c along the c-axis at the interface. This indicates that the displacement corresponds to one block of the perovskite cell, and the Tl layers are connected with Cu layer at the interface. Fig. 6.36D is a structure image of an antiphase boundary in TlBa2CaCu2O7 taken along the a-axis. The interface is {103} of the TlBa2CaCu2O7 structure, and the phase displacement is 0.21c + 0.5a. This indicates that the displacement corresponds to half-block of the perovskite cell, and the Tl layers are connected with the Ba layers at the interface. Small disordering of atoms is observed at the interface within 2 nm, which would work as insulator layers between two crystals. These interfacial structures with distances of a few nanometers are suitable for the Josephson junction devices as shown in Fig. 6.36A. Surface structures are important information on the superconducting transistors as illustrated in Fig. 6.36B. Fig. 6.37A is an HREM structure image of TlBa2Ca3Cu4O11 taken with the incident beam parallel to the a-axis. The image is observed without any contamination layers at the sample edge, and an atomic arrangement can be directly observed at the crystal surface. High-resolution images were obtained from thin samples, which were selected from crushed materials dispersed in a solution of n-butanol and dropped on holey carbon films. A mixing of Tl, Ba, Ca, and Cu atoms near the surface is observed. Fig. 6.37B is an HREM structure image of Tl0.5Pb0.5Sr2CuO5 taken along the a-axis together with projected structure model, which indicates no preferential surface structure. Fig. 6.37C is an HREM structure image of Pb2Sr2Y0.5Cu3O8, and a characteristic surface structure with valleys at Pb layers and hills between the Pb layers is observed. In the region of the hills, mixing of atoms appears to occur. An HREM structure image of PbBa0.7Sr1.3EuCeCu3O9 taken along the a-axis is shown in Fig. 6.37D, which indicates preferential atomic arrangements on the c-plane of the crystal. An HREM structure image of TlBa2CaCu2O7 is also shown in Fig. 6.37E. At the crystal edge, preferential fracture occurs between Tl and Ba layers, as indicated by arrowheads, and the fracture surface has steps with Tl layers. This result shows that the fracture surface parallel to the c-plane in the superconductor is stable and chemically

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Fig. 6.36 Schematic illustration of (A) Josephson junction and (B) superconducting transistor. HREM images at the interfaces of (C) Tl2Ba2CuO6 and (D) TlBa2CaCu2O7. No permission required.

unreactive to the solution and the atmosphere. On the other hand, the surface structures of the fracture nearly perpendicular to the c-axis show complex structures formed by rearrangement of atoms near the edge. However, Hg-based copper oxides showed stable surface (T. Oku & Nakajima, 1998a, 1999). The characteristics of various surface structures should be taken into account, when the properties are sensitive to surface structures.

Fig. 6.37 HREM images of (A) TlBa2Ca3Cu4O11, (B) Tl0.5Pb0.5Sr2CuO5, (C) Pb2Sr2Y0.5Ca0.5Cu3O8, (D) PbBa0.7Sr1.3EuCeCu3O9, and (E) TlBa2CaCu2O7, taken along the a-axis together with projected structure models. No permission required.

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A description of the possibility of light element detection, in the form of 1 atom in the plane of projection, follows. An HREM image of Hg0.5Tl0.5Ba2CuO5 taken along the [010] direction (a-axis) is shown in Fig. 6.38A, together with a projected structure model (T. Oku & Nakajima, 1998b). The upper side of the image is a vacuum in the electron microscope. There is no contamination layer at the sample edge, and the surface is stable and chemically inert to both the solution (n-butanol) and the atmosphere

Fig. 6.38 (A) HREM image of Hg0.5Tl0.5Ba2CuO5 taken along the [010] direction. (B) Enlarged HREM image of Hg0.5Tl0.5Ba2CuOx. HT indicates the Hg and Tl atoms. (C) Image depth of 200–255 Gy scale. (D) Surface structure model of Hg0.5Tl0.5Ba2CuOx and (E) corresponding simulated HREM image. No permission required.

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(air). The HREM image clearly shows the metal atom arrangements in the crystal, and the lines indicate the unit cell. The darkness and size of the black spots corresponding to (Hg,Tl), Ba, and Cu positions can be identified as being nearly proportional to their atomic numbers. A sharp flat crystal edge consisting of {100} (a-planes) is observed at the surface in atomic scale. An enlarged HREM image of the surface of Fig. 6.38A is shown in Fig. 6.38B. A weak black contrast is observed outside the (Hg,Tl) layers, as indicated by arrows in Fig. 6.38A and B. Each metal atom has an original atomic position, and there is no rearrangement of metal atoms. There is also no stable Cu-(Hg,Tl)Cu layer at the surface. To observe the black contrast outside the (Hg,Tl) layers, the 200–255 Gy scale was extracted from the original image (0–255) of Fig. 6.38B, as shown in Fig. 6.38C. Black dots are clearly observed outside the (Hg,Tl) layers as indicated by arrows and are believed to be oxygen atoms (K. Hiraga et al., 1989). From the HREM observation, a surface structure model of Hg0.5Tl0.5Ba2CuOx was proposed, as shown in Fig. 6.38D. The atomic coordinates are based on the XRD analysis of HgBa2CuOx. The Hg, Tl, and oxygen occupancies of (Hg,Tl) layers were assumed to be 0.5. In Fig. 6.38D, oxygen atoms exist 0.34 nm outside the Hg layers, as determined from the observed HREM image of Fig. 6.38B. The simulated HREM image based on the proposed model agrees well with the experimental one, as shown in Fig. 6.38E. To apply superconducting thin films to electronic nanoscale devices, the {100} plane (a-plane) should be flat and stable in atomic scale because the carrier transport is parallel to the a-axis. Most highTC superconducting oxides have layered structures, with c-planes more stable than a-planes, as described in a previous work (T. Oku & Nakajima, 1998a). In the present investigation, stable a planes of Ba-O-Ba layers were observed outside the Hg layers in Hg0.5Tl0.5Ba2CuO5 (1201-type structure). It is believed that the Hg layers are unstable because of the oxygen vacancies and that the oxygen atoms are effective for the stabilization of the {100} surface of Hg layers. The present result indicates that direct atomic determination of light elements such as oxygen is possible under special conditions.

13

Summary

The microstructural information on the perovskite-type copper oxide superconductors is useful for the development of materials for superconducting devices such as highperformance superconducting wires, nuclear electromagnetic resonance analytical systems, linear motor car and the international thermonuclear experimental reactor (Foltyn et al., 2007; Kang et al., 2006; Larbalestier et al., 2001), and DEMO (Fietz et al., 2017; Fischer et al., 2020), and for the theoretical analysis of superconducting mechanism (Anzai et al., 2013; Dal Conte et al., 2012; Zeljkovic et al., 2012). Highresolution electron microscopy and electron diffraction are quite useful for the characterization of structures of high-TC superconducting oxides on the atomic scale. For obtaining valuable information on crystal structures from the observed highresolution images, severe conditions such as thinner crystals than 2 nm and definite defocus values of 35 to 45 nm are required. The structure images taken under

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the severe conditions include information not only on arrangements and accurate coordinates of cations but also on ordered arrangements of oxygen atoms and oxygen vacancies. Many examples of crystal structure analysis of the perovskite-type superconductors were presented from the observed structure images and electron diffractions. Defects, surface, interface, intergowth, and dislocations, which cannot be investigated by diffraction methods, were observed and examined. Modulated superstructures were also observed in various copper oxides, which showed satellite reflections with various modulation wave vectors. The appearance of superstructures is sensitive to the oxygen content, and the superlattices would be due to the ordering of oxygen atoms in the Cu-O planes. Thin films and bulk materials with high JC were also investigated, and many precipitates showing Moire fringes were observed in the image, which could act as the flux-pinning centers resulting in the high JC. Such nanostructural analysis will play an important role in the development of perovskite-type superconducting copper oxide materials in the future.

Acknowledgment The author acknowledge K. Hiraga, S. Nakajima, A. Tokiwa, M. Kikuchi, D. Shindo, M. Hirabayashi, Y. Syono, T. Kajitani, H. Yamane, K. Takagi, Y. Miyazaki, Y. Amamoto, T. Hirai, Y. Tokura, E. Bruneel, S. Hoste, K. Osamura, T. Kizu, K. Kosuge, N. Kobayashi, and S. Hosoya for excellent collaborative works, providing samples and fruitful discussion.

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Part 2 Oxide and halide perovskites: Functional and advanced applications

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Perovskite-structured ceramics in solid oxide fuel cell application

7

Nurul Akidah Baharuddina, Hamimah Abd Rahmanb, Abdullah Abdul Samatc, Nafisah Osmand, Nur Syafkeena Mohd Affandid, and Suhaida Dila Safiand a Fuel Cell Institute, Universiti Kebangsaan Malaysia, Bangi, Selangor, Malaysia, bFaculty of Mechanical and Manufacturing Engineering, Universiti Tun Hussein Onn Malaysia, Batu Pahat, Johor, Malaysia, cFaculty of Mechanical Engineering Technology, Universiti Malaysia Perlis (UniMAP), Arau, Perlis, Malaysia, dPhysics Department, Faculty of Applied Sciences, Universiti Teknologi MARA, Arau, Perlis, Malaysia

1

Introduction

The fuel cell is an electrochemical device that converts chemical energy into electrical energy via chemical reaction without any combustion process. The fuels consumed at electrodes (anode and cathode) during operation are hydrogen, H2, and oxygen, O2, or air. The only product from the pure H2 supplied fuel cell is water/steam, H2O. Fig. 7.1 shows the schematic of the general working principle for the fuel cell system. The power output of the fuel cell varies depending on the type of the fuel cell. The hydrogen gas, H2 (fuel), is supplied at the anode side, while the oxygen gas, O2 (oxidant) is fed at the cathode side. Upon reaching the triple-phase boundary (TPB) at the anodeelectrolyte interface, the H2 molecules are oxidized into a proton, H+. At the same time, oxygen molecules, O2, is reduced to the anion, O2
. The H+ will flow through the electrolyte component to react with O2
and form water, H2O, as a by-product. Meanwhile, the oxidation reaction of H2 produces electrons that flow from the anode to the cathode component via the outer circuit, yielding the desired electricity. There are several types of fuel cells which mainly differ by their electrolyte component and the operational temperature. For lower temperature fuel cells, the polymer exchange membrane fuel cell (PEMFC), alkaline fuel cell (AFC), and phosphoric acid fuel cell (PAFC) are among the regular candidates for this category (Abd Aziz et al., 2020; Tahir et al., 2022). These fuel cells are suitable for use in portable and mobile applications, e.g., in transportation. Besides these as-mentioned candidates, the lower temperature fuel cells also include direct alcohol fuel cells as well as a microbial fuel cell. Despite the difference in electrolyte materials usage, the working principle of each fuel cell is similar in terms of ion movement across electrolyte components. The name of the fuel cell commonly reflected its electrolyte materials. For example, the PEMFC consists of a membrane-based electrolyte of Nafion while the PAFC uses the acid-based liquid as the electrolyte component. Both molten carbonate fuel cell (MCFC) and solid oxide fuel cell (SOFC) are categorized as high-temperature fuel cells with operational temperatures ranging from 400°C to 1000°C (N.A. Perovskite Ceramics. https://doi.org/10.1016/B978-0-323-90586-2.00006-1 Copyright © 2023 Elsevier Ltd. All rights reserved.

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Fig. 7.1 General working principle of the fuel cell system.

Fig. 7.2 Types of fuel cells.

Baharuddin et al., 2017). The electrolyte for MCFC and SOFC is made up of molten carbonate and ceramic materials, respectively. The common types of fuel cells are shown in Fig. 7.2. In terms of ionic movement, the AFC, MCFC, and SOFC involve anions flowing through the electrolyte from cathode components OH
, CO32
, and O2
, respectively (N.A. Baharuddin et al., 2017; Tahir et al., 2022). Besides the extreme working environment for high-temperature fuel cells, both MCFC and SOFC show higher efficiency in producing electrical energy. Thus, both MCFC and SOFC are strong

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223

candidates for large-power stationary applications. For SOFC, a wider range for its operational temperatures (400–1000°C) will allow the utilization of this system in more applications. In addition, solid ceramics in SOFC is beneficial in terms of the system handling aspect compared to the MCFC that contains liquid molten carbonate.

1.1 Solid oxide fuel cells (SOFCs) Besides being known for their high efficiency, SOFCs are also known for their fuel flexibility, in which hydrocarbon fuel becomes an option. For the hydrocarbon-fueled SOFC, it is not entirely green technology as carbon is produced from the reaction. Nevertheless, it is still a low emission technology due to the high-operational temperature that helps reform the process and minimize the emission. The biogas produced from industrial and domestic waste is also a good fuel candidate for SOFC. In this context, the utilization of SOFC in biogas plants will substitute the Genset or diesel generator usage, which is known for its higher carbon footprint (Saadabadi et al., 2019). Tailoring the SOFC’s working principle, it is similar to the other types of fuel cells. The conventional SOFC allowed the anion O2
to pass through the electrolyte and form water at the anode side. Thus, the rate of oxygen reduction reaction (ORR) at the cathode site is significantly important (Huang & Goodenough, 2009; Singhal & Kendall, 2003). Thus, besides an excellent electrolyte component that allows higher ionic movements, selecting suitable cathode materials with a high electrocatalytic behavior toward ORR is crucial for the higher performance of SOFCs. On the other hand, the H2 is oxidized at the anode site, usually made up of metal-ceramic combinations. The ceramics part of the anode is usually the same as the electrolyte materials to maintain an excellent thermal and chemical compatibility between these two components. Fig. 7.3 shows the working principle for the conventional SOFCs. This type

Fig. 7.3 Working principle of the conventional O2
-SOFC system.

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Perovskite Ceramics

Fig. 7.4 Working principle of the H+-SOFC system.

of SOFCs is also commercially available in the market for stationary applications, for example, Bloom Energy (US), Ceres Power (UK), and Mitsubishi Power (Japan). Continuous progress in this field led to the introduction of a new type of SOFC called proton (H+)-conducting SOFC. The main difference for H+-SOFCs compared to the O2
-SOFC is that the mobile ions pass through the electrolyte component. By contrast with the conventional O2
-SOFCs, this H+-SOFC allows the proton H+ to flow from the anode side to the cathode . The H2O is then produced at the cathode side. Due to its working principle, newly developed material candidates are required for the electrolyte and cathode components, considering that the electrolyte will allow H+ movement rather than O2
, and the cathode still works well even after the steam is produced at its part. Fig. 7.4 shows the schematic of the working principle for the H+-SOFC system.

2

Perovskite-structured electrolyte for SOFCs

SOFCs have gained significant interest in scientific research in hydrogen energy and fuel cell technology due to their great potential as a relevant future power generation. In many types of fuel cells, they showed the same typical process and design, which are a positive electrode (cathode), a negative electrode (anode), and an electrolyted (Chroneos et al., 2011; Druce et al., 2014). The two main functions of an electrolyte are (i) as a separator located in between the anode and cathode parts and as (ii) a medium for the conduction of ions where an electrochemical reaction occurs. This electrolyte is the heart of an SOFC that induces ion flow by transferring ions from the fuel to the air (H+-SOFC) or from the air to the fuel (O2
-SOFC), as the SOFC comprises porous electrodes and a dense electrolyte (Fleischhauer et al., 2014). The flow of ions through the electrolyte is vital to complete a fuel cell operation principle. Thus, the efficiency of

Perovskite-structured ceramics in solid oxide fuel cell application

225

SOFCs is strongly influenced by the performance of electrolytes. For operating temperatures between 700°C and 800°C, the SOFC’s performance could be improved by up to 33% through active chemical processes induced by hydrogen usage, according to numerical simulations conducted by Park et al. (2018).

2.1 General properties of the SOFC electrolyte Generally, the electrolyte component must fulfill several requirements such as being dense and leak-tight to reduce the ohmic and thermal shock resistance, being stable in reducing and oxidizing environments, having good mechanical properties at operating temperature, showing high ionic conductivity (about 0.1 S cm
1) and low value of electronic transference number ( 800°C), proton conductivity drops because a lower concentration of charge carriers is caused by lower proton (water) solubility in the oxides. During this temperature, it is dominated by oxygen conduction through a vacancy mechanism. Therefore, perovskite is a suitable material in intermediate-to-low temperature operation due to an exothermic reaction of the adsorption of water. The exothermic reaction releases less energy than the total energy of the reactant, where only small activation energy is needed for proton conduction. Due to the thermal activation of the ions moving from one crystal lattice site to another, protons within the perovskite crystal lattice exhibit thermal migration in the direction of the electric field. In the literature, two types of proton mobility through the perovskite crystal lattice are commonly adapted: (i) vehicular mechanism and (ii) Grotthuss mechanism (Radenahmad et al., 2016; Zuo et al., 2012). The vehicular mechanism is the transportation of protons as passengers on mobile host species or diffusion of a proton into the lattice accompanied by the carrier (OH
, H3O+, or NH4+ group) and H+ ions as presented in Fig. 7.6. Most often, charge carriers of OH
are typically used to explain the proton conduction in perovskite oxide. As mentioned in Eq. (7.1), OH
can be obtained through the interaction of water with surface oxygen vacancies to form protonic defects in the form of hydroxyl complexes.

Fig. 7.6 Possible mechanism of a proton inside the electrolyte through vehicle mechanism. Adapted from Regalado Vera, C. Y., Ding, H., Peterson, D., Gibbons, W. T., Zhou, M., & Ding, D. (2021). A mini-review on proton conduction of BaZrO3-based perovskite electrolytes. Journal of Physics: Energy, 3(3). https://doi.org/10.1088/2515-7655/ac12ab.

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Perovskite Ceramics

Fig. 7.7 Sequence of three snapshots from ab initio MD simulations showing inter-octahedra proton hopping in orthorhombic CaZrO3 (the Ca ions are omitted for clarity) (Islam et al., 2001).

The Gr€ otthuss mechanism (Mather et al., 2021; Shi & Yun, 2020), on the other hand, describes the ‘free’ protons moving to the neighboring oxygen ions of which the H+ is located in the electron cloud of the oxygen ions. Hydroxide ions having a positive charge of OH with respect to the oxygen lattice are formed. While protons interact strongly in the created O-H bond, it is also affected by other surrounding (neighboring) oxygen anions. The protons jump from one oxygen ion to the neighboring oxygen ion and their hopping activation energy depends on the oxygen-oxygen (O-O) distance (less than 0.24 nm) (S´wierczek & Skubida, 2017). Such a phenomenon is also reported by Islam et al. (2001), as illustrated in Fig. 7.7, that explained that the movement of the proton occurs through a simple transfer of a proton from one oxygen ion to the neighboring ion for their octahedra structured CaZrO3 electrolyte (from (i) to (iii)). Based on the molecular dynamics (MD) simulations, they also found that the O-H group experienced rotational and stretching motion (displayed in (ii)). This allows the reorientation of the proton toward the neighboring oxygen ion before the transfer process. The ZrO6 tilting influences this diffusion path within the orthorhombic structure, which leads to shorter O-O separations between the vertices of the adjacent octahedra. However, the ideal perovskite structure is cubic. The cubic structure makes the positions of the oxygen ions energetically the same; thus, we deduce with equal probability that proton hopping is possible in any direction (Assirey, 2019).

3

Perovskite-structured electrodes for SOFCs

SOFCs present many exciting material challenges. Material properties of an electrode (anode and cathode) are important factors that determine the electrochemical performance of an overall SOFC system and technology. Knowledge and understanding of the general properties or requirements of SOFC electrodes are essential in selecting any materials to be used as SOFC electrodes. The perovskite-type ceramic oxide

Perovskite-structured ceramics in solid oxide fuel cell application

231

crystal structures are commonly employed as SOFC electrode materials because of their geometrical and electrostatic configuration stability. It is well known that many oxides of this type are stable even when defects are introduced into their structures. The selection of the SOFC electrodes must be based on the individual properties of each component (anode or cathode) and the compatibility between these two components with other SOFC components such as electrolyte and interconnect. This section discusses the general properties of SOFC electrodes and materials progress of perovskite-structured electrodes for O2
and H+-conducting SOFCs.

3.1 General properties of SOFC electrodes The primary function of the SOFC electrodes is to bring about a reaction between reactants (fuels from the anode side and oxidants from the cathode side) and electrolytes without themselves being consumed or corroded. At the anode, oxidation or burning of fuels such as hydrogen gas (H2), methane (CH4), and other natural gases takes place, while reduction of an oxidant such as oxygen gas (O2) and air takes place at the cathode. The general properties of SOFC electrodes, including the anode and cathode, are quite similar. An ideal electrode for the SOFC application should have the following properties (Abd Aziz et al., 2020; Y. Liu et al., 2021; A.A. Samat et al., 2021; Singh et al., 2021; Tahir et al., 2022): l

l

l

l

l

l

l

l

l

l

l

Mixed ionic and electronic conducting material. High electronic conductivity (>100 S cm
1) and low ionic conductivity (10
1 S cm
1). Low area-specific resistance (ASR) or polarization resistance, Rp value (800°C) SOFC. The ceramic interconnects are oxides, which are very stable in an oxidizing atmosphere, but their cost is high and they exhibit lower electrical conductivity at lower operating temperatures. However, the high operating temperature results in high system cost/complexity and poor long-term durability. Recently, decreasing the operating temperature of SOFCs to the low-temperature range (400–600°C) has attracted intensive attention,

Fig. 7.10 Schematic diagrams of electrochemical reaction mechanisms of different types of cathode materials for H+-conducting SOFC application: (A) PEC cathode, (B) MOEC cathode, (C) MPEC cathode, (D) MOEC-based composite cathode, and (E) THOEC cathode.

Perovskite-structured ceramics in solid oxide fuel cell application

241

Table 7.4 Types of cathode materials for H+-SOFC. Description

Type of cathode

Figure

Active reaction site is limited only at cathode-electrolyte interface or TPB. O2
moves along the cathode surface before reaching at TPB to react with H+ of electrolyte, forming H2O. A quite long pathway taken by the O2
to reach at the TPB results in high cathode Rp Active reaction site is also limited only at cathode-electrolyte interface or TPB. However, O2
can move long the cathode surface and in the bulk cathode to reach at TPB, forming H2O. Larger pathway of O2
movement through cathode surface and bulk cathode to the active reaction site reduces cathode Rp Active reaction site is also limited at cathode-electrolyte interface or TPB, but at a larger area. O2
moves along the cathode surface and H+ moves in the bulk cathode. H2O can form at TPB, on cathode surface and in bulk cathode. Larger active reaction results in lower cathode Rp than PEC Active reaction site is larger and not only limited at cathode-electrolyte interface or TPB. Pathway of O2
to the active reaction site is also larger. O2
can move along the cathode surface and in the bulk cathode to react with H+ to form H2O. H2O can form at TPB and in the bulk of composite cathode itself

Pure electronic conductor (PEC) Example: Platinum (Pt)

Fig. 7.14A

Mixed oxide ion (O2
) electronic conductor (MOEC) Examples: La0.6Sr0.4CoO
3 (Abdul Samat δ et al., 2019) Sm0.5Sr0.5CoO3
δ (Oda et al., 2014) La0.6Sr0.4Co0.2Fe0.8O3
δ (Ismail et al., 2020) BaFe0.6Co0.3Ce0.1O3
δ (Zhang & Li, 2016) Ba0.5Sr0.5Co0.8Fe0.1Ta0.1O3
δ (F. Wang, Xu, et al., 2021) Mixed protonic (H+) electronic conductor (MPEC) Examples: BaCe0.9Yb0.1O3
δ (E. Fabbri et al., 2011) BaZr0.7Pr0.1Y0.2O3
δ (E. Fabbri et al., 2009) BaCo0.4Fe0.4Zr0.2O3
δ (Zohourian et al., 2017) BaCexFe1
xO3
δ (Tao et al., 2009) BaCe0.5Fe0.3Bi0.2O3
δ (Shan et al., 2017) Composite cathode (MOEC + H+conducting electrolyte) Examples: La0.6Sr0.4CoO3
δBaZr0.1Ce0.7Y0.1Yb0.1O3
δ (M.S. Wang et al., 2015) Sm0.5Sr0.5CoO3
δBaCe0.8Zr0.1Y0.1O3
δ (Dailly et al., 2017) La0.6Sr0.4Co0.2Fe0.8O3
δBaZr0.1Ce0.7Y0.2O3
δ (K. Miyazaki, Ding, et al., 2020) La0.6Sr0.4Co0.2Fe0.8O3
δBaZr0.1Ce0.7Y0.1Yb0.1O3
δ (Shimada et al., 2021) Ba0.5Sr0.5Fe0.9Ni0.1O3
δSm0.2Ce0.8O1.9 (Y. Y. Ding et al., 2012)

Fig. 7.14B

Fig. 7.14C

Fig. 7.14D

Continued

242

Perovskite Ceramics

Table 7.4 Continued Description

Type of cathode

Figure

Overall cathode surface is the active reaction site. Therefore, H2O can form at any point on cathode surface and at cathode-electrolyte interface or TPB. Ideally, the cathode Rp is the lowest

Triple protonic-ionic-electronic (H+/O2
/e
) conductor (THOEC) Examples: SrSc0.175Nb0.025Co0.8O3
δ (Zhu et al., 2018) Ba0.5Sr0.5Co0.8Fe0.2O3
δ (An et al., 2018) La0.25Pr0.25Sr0.5FeO3
δ (Ma et al., 2020) La0.7Sr0.3Mn0.7Ni0.3O3
δ (N. Wang et al., 2020) PrBa0.5Sr0.5Co1.5Fe0.5O5+δ (S. Choi et al., 2018)

Fig. 7.14E

Fig. 7.11 Planar SOFC. From solid oxide fuel cells (SOFCs) (2021). Retrieved September 1, 2021, from https://www. doitpoms.ac.uk/tlplib/fuel-cells/high_temp_sofc.php under Creative Commons AttributionNonCommercial-ShareAlike 2.0 UK: England & Wales License.

especially in the usage of metal as an interconnect material. The metallic interconnect is cheaper than the ceramic ones. Metallic interconnecting is becoming common as it is composed of high electrical conductivity at low temperatures, for example, ferritic stainless steel (Brouzgou et al., 2017). Besides the high electrical conductivity of metallic interconnect, it should also be chemically and physically compatible with

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oxidation and reduction conditions at the cathode and anode. Hence, interconnectors must fulfill the following criteria (Brouzgou et al., 2017; Uehara et al., 2011; Yeh et al., 2015): (1) High tightness to avoid direct mixing of fuel and oxygen during operation. (2) Chemical stability and compatibility with other SOFC components in both anode oxidation and cathode reduction reaction and corrosion. (3) High electronic conduction with negligible ionic conductivity to enhance power output intensity. (4) The thermal expansion coefficient (TEC) should match with the fuel cell components (electrodes and electrolytes) and coating material. In general, the acceptable TEC values are in the range of 10
13  10
6 K
1. (5) Good mechanical strength and durability due to the high operating temperature to avoid damages. (6) High thermal conductivity to enable a uniform stack temperature. (7) Low cost and easier methods of fabrication and inexpensive raw materials.

Ferritic stainless steels are considered as the appropriate materials to be used as SOFC interconnects at low-temperature operations. The main alloying element in this stainless steel is chromium (Cr) in the range of 10–26 wt%. The metallic interconnects exhibit higher electronic conductivity compared with ceramics, but they are not stable in oxidizing atmospheres. Hence, at high temperatures, the main drawback is that Cr tends to evaporate and deposit on the cathode surface of SOFC (Brouzgou et al., 2017). This can cause damage to the stability of fuel cells in long-term conditions. Therefore, a protective coating is essential for metallic interconnects to overcome the problems mentioned above and prevent the out-diffusion of the chromia scale, which is badly hazardous to the overall cell performance.

4.1 General properties of perovskite coating for the SOFC metallic interconnect Several solutions have been approached during the last decade, predominating in the surface modification of the metallic interconnects via protective oxide layers (ceramic one) deposition on them. Application of protective ceramic coating is the practical approach to prevent Cr poisoning on metallic interconnects. Currently, ceramic conductive materials are widely used as thin protective layers deposited on metallic interconnects. The application of coating can increase the lifetime of metallic interconnects, especially under cathode conditions. Reactive element oxides, perovskites, spinels, and dual layers are the kinds of coatings that have also been developed (Brouzgou et al., 2017; K.H. Tan et al., 2020). Protective coatings for interconnects should have the listed properties (K.H. Tan et al., 2019): (1) (2) (3) (4)

High electrical conductivity with area-specific resistance (ASR) lower than 0.1 Ω cm2. Uniform and dense (high density) microstructure to minimize chromium diffusivity. High stability (chemical and physical) during operation of SOFC. Intimidate contact with the interconnect to avoid contamination which can affect the conductivity performance of the interconnect.

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Spinel and perovskite oxide coating are two primary protective coatings applied for planar SOFC interconnect (Mah et al., 2017). Spinel oxide coating with general formula AB2O4 is one of the currents focusing on protective coating materials. The electrical conductivity and thermal expansion can be controlled by adjusting the ratio of A and B cations (Petric & Ling, 2007). The application of spinel coatings has successfully proven its ability to suppress the volatilization of chromium oxide (Cr2O3) to the scale surface of the interconnect. However, spinel coating mainly focuses on hightemperature SOFC application where its potential and ability to perform at low temperature still need to be investigated and explored. Perovskite materials classified as mixed ionic and electronic conducting (MIEC) materials have been widely used for SOFC cathode but not as much as coating materials for the interconnect. Nowadays, rare-earth perovskite is applied as an interconnect coating material because it has a good coating adhesion on stainless steel and excellent electrical conductivity in intermediate to low temperature SOFCs (K.H. Tan et al., 2020). The perovskite structure has the general formula ABO3, where A and B are cations with a total charge of +6. The A cations are occupied by lower valence elements such as La, Sr, and Ca. Meanwhile, the B cations are occupied by a higher valence such as Ti, Co, Fe, and Ni. Fig. 7.12 shows the schematic structure of cubic perovskite. Most perovskite materials have a mixture of rare and alkaline earth (such as La and Sr, Ca, or Ba) at the A-site cation. Meanwhile, the B-site cation is a reducible transition metal such as Mn, Fe, Co, or Ni. Generally, ideal perovskite oxide has a cubic crystal structure at room temperature. However, due to the cationic substitutions at the A and B sites, differences in the ionic radii of the dopant and solvent host, and different charges in the A- and B-site cations, the structure can be distorted (usually octahedral) due to the oxygen/cationic vacancies and change in the angles between cations and oxygen (Kaur & Singh, 2020). The MIEC perovskite cathode has received much focus as this is a promising cathode material that can achieve high ionic conductivity at low temperatures (Da C^ orte et al., 2013). The MIEC cathode provides an increment of phase conduction at TPB and allows simultaneous transportation of electronic species that can increase the possible reaction, useful for the interconnect coating material.

Fig. 7.12 Perovskite structure of ABO3 (Liu, Tan, & Li, 2006).

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4.2 Materials progress of perovskite-structured coating for the SOFC metallic interconnect Perovskite coatings were discussed from their microstructure, oxidation behavior, and area-specific resistance. The compatibility of the perovskite coating with the ferritic stainless steel interconnect will also be observed through their thermal expansion difference. Perovskite oxide material as the protective coating has achieved lower areaspecific resistance, although perovskite layers do not prevent the formation of chromia on the surface of the oxidized coating/metal composites interconnect. This occurred due to the MIEC properties of the perovskite cathode material enabling the reduction of oxygen to oxygen ions which diffuse inward through the triple-phase boundary (TPB) and double-phase boundary (DPB) of the perovskite in the long-term operation (Fig. 7.13) (Ciucci et al., 2011). As this occurred, the oxygen ions would react with chromium to form the chromia scale on the surface of the coating/interconnect (de Larramendi et al., 2016). In order to avoid undesirable phenomena related to Cr species vaporization and diffusion into the cathode compartment that causes degradation of the SOFC electrochemical performance, which in turn shortens the service life, the stainless steel interconnect surface needs a modification. The modification is made by introducing protective coating layers on the stainless steel surface. Most of the materials used for the coating purpose is perovskite-structured cathode materials including lanthanum strontium cobalt ferrite (LSCF) ( J.P. Choi et al., 2021; Przybylski et al., 2014; Tsai et al., 2010), lanthanum strontium manganite (LSM) ( J.W. Chen et al., 2017; Hwang & Choi, 2009), manganese-cobalt-based spinel (Mn-Co) (Kruk et al., 2013; Zanchi et al., 2021), lanthanum manganite copper oxide (LaMnCuO) (Ananyev et al., 2018), and barium strontium cobalt ferrite (BSCF) (K.H. Tan et al., 2020). Those materials have shown characteristics as a good protective-conductive coating to achieve low electrical surface resistance, effectively prevent chromium migration from the chromium-containing steel substrate, matched thermal expansion coefficient, and high electrical conductivity at low-tointermediate SOFC operating temperatures (500–850°C).

Porous electrode

DPB

O2

Dense electrode

O2´ e´

DPB

Cathode TPB Electrolite Fig. 7.13 TPB and DPB of perovskite material (de Larramendi et al., 2016). Idoia Ruiz de Larramendi, Nagore Ortiz-Vitoriano, Isaen B. Dzul-Bautista and Teo´filo Rojo (2016). Originally published in https://www.intechopen.com/chapters/49510 under Creative Commons Attribution 3.0 License license. Available from: 10.5772/61304.

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Fig. 7.14 Micrographs of oxide scale/alloy interface of LSCF-coated Crofer22 APU at 800°C for 200 h by screen painting (Tsai et al., 2010).

Among various perovskite materials, LSCF is the most investigated perovskitetype material, used as cathodes in SOFC that exhibit ionic and electronic conductivities and high chemical stability with an electrolyte such as yttrium-stabilized zirconia, and has high electrochemical activity at intermediate to high temperature. Therefore, LSCF is a promising candidate that has been extensively studied as a perovskite protective layer for interconnect material. It was found that LSCF thick films (Przybylski et al., 2014) with a thickness of 20–30 mm developed by the screen printing technique on Crofer 22APU steel have a lower oxidation rate than uncoated steel. After exposure to isothermal oxidation in the air for 100 h, the area-specific resistance (ASR) values of 0.025–0.010 Ω cm2 were recorded at temperatures of 600–800°C. The LSCF coating can also maintain good adhesion to the Crofer 22APU steel substrate after annealing and cyclic oxidation. Furthermore, LSCF coated on Crofer 22APU steel via screen printing (Tsai et al., 2010), sintered at 1100°C, and then exposed to oxidation for 200 h (800°C) exhibited an ASR value of 0.054 Ω cm2. After 200 h, the ASR was further increased. The changes of ASR with oxidation time were attributed to the increment of oxide scale thickness. A thin layer oxide scale (1 μm) was observed after oxidation (Fig. 7.14). The other study showed similar behavior (Lee et al., 2010). Other than LSCF, LSM has been applied as a perovskite protective coating for interconnect. Previous studies identified a drawback of LSM coating where it tends to crack after undergoing heat treatment (calcination or annealing). The crack formation was due to the phase changes from amorphous to crystalline (Fig. 7.15), which caused volume shrinkage of the LSM coating film ( J.W. Chen et al., 2017; Chu et al., 2008; Rashtchi et al., 2013). The 5-mm LSM thin film ( J.W. Chen et al., 2017)

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Fig. 7.15 Micrographs of the Crofer22 APU (A) coated surfaces and (B) coated surfaces oxidized at 700°C for 1 h (Chu et al., 2008).

fabricated through the magnetron sputtering technique has managed to give the ASR value of 0.035 Ω cm2 after oxidation at 800°C for 1500 h. Additionally, the LSM protective coating (30–50 mm) on Crofer22 APU, SUS430, and SUS304 alloys ( J.W. Chen et al., 2017) was fabricated by the plasma spraying method and went through an oxidation process for 300 h at 800°C offering ASR values of 0.005–0.024 Ω cm2. The preoxidation process has been applied to solve the cracking problem of LSM coating; the use of nanostructured powders in the form of spherical particles is also recommended, as well as the application of plasma spray techniques that can form quite a dense layer without cracking (Brylewski et al., 2012; J.W. Chen et al., 2017). Lanthanum strontium cobalt (LSC) materials have also been investigated as a perovskite coating stainless steel interconnect for SOFC (Ren & Bao, 2014; Shong et al., 2011). The La0.8Sr0.2CrO3 coating layer (Ren & Bao, 2014) prepared by plasma spraying and modified by posttreatment through dip coating in LSC slurry has provided a denser coating layer with the ASR value of 0.015 Ω cm2 at 800°C after oxidation at 1000°C for 20 h compared to the unmodified LSC, which has the ASR value of 0.18 Ω cm2 at the same condition. The modification of the LSC coating layer appears to be the best solution to the interconnect’s high ASR increase rate when exposed to oxidising atmospheres for a longer duration. The abovementioned perovskite coating materials for the stainless steel interconnect mainly focus on the intermediate to high-temperature SOFC. The high sintering temperature of the perovskite materials is one of the main drawbacks even though they lead to a decrease in the interfacial contact resistance. The high sintering temperature can lead to their compaction followed by Cr diffusion, especially at the cathode compartment. Therefore, the application of low sintering temperature of perovskite coating materials and their ability to perform (low ASR and high electrical conductivity) at low operational temperature SOFC has become a significant concern. A recent study by K.H. Tan et al. (2020) focused on the application of composite barium strontium cobalt ferrite-samarium doped ceria carbonate (BSCF-SDCC) perovskite coating on ferritic stainless steel 430 (SUS430) aims for the low-temperature (400–600°C) SOFC operation. BSCF is preferably for interconnect protective material due to the least polarization resistance compared to other perovskite materials.

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The findings showed that a dense and compact microstructure of the BSCF-SDCC coating layer was obtained after sintering at 500–600°C. The existence of carbonate in the composite BSCF-SDCC has enabled it to be sintered at low temperatures (K.H. Tan et al., 2020). However, the increment of sintering temperature has decreased the carbonate bonds, and this phenomenon led to the formation of pores and cracks in the coatings as shown in Fig. 7.16E and F. The BSCF-SDCC coating sintered at

Fig. 7.16 SEM morphology of sintered BSCF-SDCC coating by EPD at 10 V and 10 min (K.H. Tan et al., 2020).

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550°C, 600°C, and 650°C fulfill the requirement of interconnect ASR value (below 0.1 Ω cm2) after 500 h oxidation. The ASR measurement revealed that the highest electrical conductivity was obtained for the BSCF-SDCC coating layer (80.7 mm) sintered at 600°C with 0.073 Ω cm2 for 500 h of oxidation at a low-temperature SOFC of 600°C (K.H. Tan et al., 2020). These findings showed the importance of controlling the sintering temperature to obtain good characteristics of perovskite coating for the SOFC interconnect. Hence, research into the development of perovskite coatings for low-temperature SOFC applications has a lot of potential.

5

Summary

This chapter discussed the brief knowledge of the SOFC technology and research outputs related to the perovskite-structured ceramics in this field. The utilization of perovskites-based materials in SOFC covers its significant components, including electrolyte, electrodes, and interconnect. The MIEC behavior of perovskites-based materials offers a larger TPB to allow higher ORR reaction, thus yielding more ions to react with the reduced fuel, H2, from the anode. The main mechanism for O2-diffusion in the perovskite structure is via the vacancy diffusion related to the mobile oxygen vacancies concentration. In addition, the flexibility of the ABOx structure toward dopants (ion species) results in the variation of material candidates for SOFC components. Controlling the perovskite’s A-site will help determine the optimum oxygen vacancies, then influence the ionic conductivity of the materials. On the other hand, the B-site is necessary for the electrical properties of the SOFC’s materials, specifically for electrode components. The efforts on perovskite-structured ceramics alteration are ongoing despite the numbers of materials reported due to the improvement potential at both the A- and B-site of the perovskites.

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Perovskite membranes for oxygen separation

8

Daniel Dornellas Athaydea,b, Julius Motuzasc, and Wander Vasconcelosd a School of Engineering and Architecture, FUMEC University, Belo Horizonte, Brazil, b Department of Chemical Engineering, Federal University of Minas Gerais, Belo Horizonte, Brazil, cFIM2 Lab—Functional Interfacial, Materials and Membrane Laboratory, School of Chemical Engineering, The University of Queensland, Brisbane, QLD, Australia, d Department of Metallurgical and Materials Engineering, Federal University of Minas Gerais, Belo Horizonte, Brazil

1

Introduction

Perovskite oxides have been used for membrane manufacture for over 30 years now, stemming from the pioneer study by Teraoka et al. in 1985 on SrCo0.8Fe0.2O3
δ (Teraoka et al., 1985). Since then, the research effort to develop and design inorganic perovskite membranes has been consistently growing, as one may notice from the increase in the number of publications shown in Fig. 8.1. These materials represent a potential economical, clean, and efficient alternative for oxygen manufacture, as oxygen (O2) is conventionally separated from atmospheric air by energy- and costintensive processes, such as cryogenic distillation and pressure swing adsorption. Moreover, other technological applications have been addressed for the use of perovskite membranes for pure oxygen production, such as oxy-fuel combustion and natural gas conversion to syngas. Perovskite membranes allow simultaneous transport of both ions and electrons and are called mixed ionic and electron conductors (MIEC). They are superior to other inorganic materials, such as fluorites, that conduct only ions and would require external electrodes for electron transport, which represents a higher complexity in the engineering design of membrane modules. Moreover, perovskite membranes are dense, allowing oxygen transport only as an ion (O2
) when submitted to high temperatures (typically within 700–1100°C). Meanwhile, the N2 present in atmospheric air is not able to diffuse through the dense perovskite membranes and thus has a theoretical infinite selectivity. In other words, the membrane permeate is composed of pure oxygen, a remarkable superiority when compared to porous ceramic membranes or polymeric membranes. The perovskite oxide structure is cubic with the general formula ABO3. The ideal ABO3 structure (without defects) is not capable of conducting ions. However, the presence of oxygen vacancies on the perovskite structure allows ionic conduction. High concentrations of oxygen defects can be obtained by partial substitution at the A-site cations to increase ionic conduction, whereas partial substitution at the B-site increases electronic conduction. Many compositions have been reported and Perovskite Ceramics. https://doi.org/10.1016/B978-0-323-90586-2.00010-3 Copyright © 2023 Elsevier Ltd. All rights reserved.

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Fig. 8.1 Number of published papers per year based on the keyword “perovskite membrane.” Extracted from Web of Science.

a series of reviews listed the large variety of tested perovskites (D.D. Athayde et al., 2015; P.L. Rachadel et al., 2013; ‘J. Sunarso et al., 2008; J. Sunarso et al., 2017). Among these compositions, the most promising perovskites to date for O2 separation from air are the Ba1
xSrxCo1
yFeyO3
δ (BSCF) and La1
xSrxCo1
y FeyO3
δ (LSCF) (x and y represent the partial substitution in A- and B-sites, respectively, and δ is the amount of oxygen vacancies). Both BSCF and LSCF were shown to achieve economically attractive performances for O2 production, although further improvements have already been reported. For instance, in the early 2000s, the O2 fluxes using BSCF perovskite membranes reached 1.4 mL cm
2 min
1 (Shao et al., 2000). After just a decade, in 2011, the BSCF perovskite had reached values 10 times higher, achieving O2 fluxes of 11. 4 mL cm
2 min
1 ( J. Sunarso et al., 2011). In 2017, a dual-phase copper oxide doped Ba1
xSrxCo1
yCuyO3
δ perovskite resulted in O2 fluxes of 27.5 mL cm
2 min
1 (Leo et al., 2017). All these upgrades were obtained within a couple of decades, showing that the research and development of perovskite membranes for O2 separation is still ongoing and is paramount for obtaining a technology that will allow widespread industrial use. The applications involving the inorganic perovskite oxides play an important role in the sustainable development of novel energy delivery technologies. For instance, the ability to transport O2
ions allows its use as electrodes in solid oxide fuel cells (SOFCs) ( J. Sunarso et al., 2017). As for perovskite membranes, they evidently can be used for pure oxygen production (Anderson et al., 2016). Another potential application of perovskite membranes is the integration of air separation units into coal gasification systems or in oxy-fuel combustion. In fact, the latter has been the main focus of researchers regarding the use of oxygen transport membranes (OTMs), with intense effort aimed at engineering and process design (Castillo, 2011; W. Chen et al., 2014; W. Chen et al., 2015; Engels et al., 2010; Stadler et al., 2011). The integration of OTMs in oxy-fuel combustion consists of the use of membrane modules to provide pure O2 for coal combustion in energy delivery systems. The absence of N2 during the combustion stage results in an exhaust gas without the presence of nitrogen compounds, such as NOX compounds. Moreover, the exhaust gas is composed mainly of

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CO2, with concentrations higher than 95%, allowing reutilization of the CO2 within the system or CO2 storage. This concept is an outstanding improvement when compared to conventional coal-based energy delivery systems, as it allows a near-zero emission combustion technology. Nevertheless, the commercial use of perovskite membranes as OTMs still has some challenges to be overcome. The high temperatures required for operation of these membranes lead to severe issues with membrane sealing, even resulting in system failure at higher temperatures (X. Tan et al., 2010). The stability at reducing atmospheres (i.e., CO2) is also a serious issue, especially when using cobalt containing perovskites, such as the BSCF compound. This has a significant impact on the design of membrane modules, because the integration of OTMs in oxy-fuel combustion is usually based on a configuration in which the CO2 is the sweep gas on the permeate side of the membrane. The LSCF perovskite is chemically more stable, though it reaches lower O2 fluxes. Moreover, the low mechanical stability of perovskite membranes, due to the higher brittleness typical of ceramic materials, has to be improved for the thin membranes usually required for higher oxygen fluxes (Song et al., 2019). Despite these challenges still to be overcome, perovskite membranes have already been successfully tested in pilot plants. For instance, the perovskite membranes used in a plant developed by Air Products in the United States delivered 16 tons per day of high purity (>99.9%) oxygen for long periods (15,000 h) (Anderson et al., 2016). The use of perovskite membranes is a very promising technology, with a high potential for the design of engineering systems that are more sustainable. There are challenges, and they are being frequently addressed and confronted by researchers. In this chapter, we will address these subjects and further structure, properties, performance, and challenges of the perovskite membranes will be developed. Firstly, the separation phenomena are explained, focusing on the perovskite structure and how it allows simultaneous ion and electron transport. Next, the synthesis and preparation (i.e., conformation) of perovskite membranes is shown in detail. This section will be followed by an analysis of the membrane’s performance under several conditions. Finally, future perspectives will be discussed, detailing further the challenges and the feasibility of this technology.

2

Perovskite structure and separation phenomena

2.1 Perovskite structure The ideal ABO3 perovskite displays a cubic structure (Fig. 8.2), with A being a large metal ion surrounded by 12 oxygen atoms and B, a smaller cation surrounded by six oxygen atoms. A classic perovskite example is CaTiO3, in which Ca is a cation of large ionic radius and Ti is a small and highly charged cation. Typically, A is commonly a rare earth, alkali or alkaline families (e.g., La, Na, Ca, Sr, Ba), while B is a transition metal cation (e.g. Fe, Co, Ni, Cu). The lattices are able to accommodate different cations, sizes, and charges. Usually, the compositional formula is A2+ B4+ O3; however, other formulas are also observed: A1+ B5+ O3 and A3+ B3+ O3.

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Fig. 8.2 Representative structure of perovskite compounds. Modified from Athayde, D. D., Souza, D. F., Silva, A. M. A., Vasconcelos, D., Nunes, E. H. M., Diniz da Costa, J. C. & Vasconcelos, W. L. (2015). Review of perovskite ceramic synthesis and membrane preparation methods. Ceramics International. https://doi. org/10.1016/j.ceramint.2016.01. 130.

The stability of the crystalline structure is also attained for perovskite oxides with more than 2 cations. It is frequent to find perovskites with 3 and 4 cations (such as SrCo0.8Fe0.2O3
δ and BSCF), although perovskites with 5 different cations were already synthesized [such as BSCFs doped with scandium (D.D. Athayde et al., 2017) or yttrium (P. Haworth et al., 2011)]. The cubic crystal shape is commonly preserved if the tolerance factor t, also called the Goldschmidt factor and detailed in Eq. (8.1), is within 1.0 and 0.75. r + rO t ¼ pffiffiffi A 2ð r B + r O Þ

(8.1)

where rA and rB represent the radii of the A-site and B-site cations, respectively, and rO is the ionic radius of oxygen. Tolerance factors outside of this range lead to the formation of different crystal structures. The capacity of perovskite oxides to conduct oxygen arises from the presence of defects in the perovskite structure, such as point defects, line defects, and plane defects. Examples of point defects that cause lattice diffusion of the ion entities are vacancies and interstitial atoms or ions. For oxygen transport, the focus is on perovskites with anionic (oxygen) vacancies with a formula that represents directly the amount of oxygen vacancies: ABO3
δ. These defects can be naturally found in the structure due to interactions with the environment or can be found due to the presence of impurities, the latter being a common strategy to obtain higher concentrations of oxygen vacancies. For example, the addition of a lower valence A-site cation (such as doping an A3+ B3+ O3 compound with a A2+ cation) requires a lower amount of oxygen anions (oxygen vacancies) in the structure to maintain the electrical neutrality. Nevertheless, some compounds predominantly form interstitial cations instead of oxygen vacancies by the addition of a lower valence A-site cation. Hence, comprehension of the structural characteristics of perovskite compounds is advised prior to the design of a novel material with a high concentration of oxygen vacancy. Whereas the A-site alkaline-earth or rare-earth cation is more commonly related to the oxygen vacancy concentration, the B-site cation is occupied by transition metals

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and it has a strong influence on electronic conductivity. The transition metal cations can often be found in different valence states. For instance, Fe ions that are usually found in the following states: Fe2+, Fe3+, and Fe4+. Therefore, the existence of B-site cations with different valence states—such as Ti, Fe, and Co—provides the conditions for electronic conduction. Improvement of ionic and electronic conductivity is frequently obtained by adding dopants to A- and B-sites, leading to compounds with 3 or 4 metal ions. The stability of the cubic perovskite structure also allows the existence of perovskite compounds with 5 metal ions, though these are less common. The general formula for perovskite with 4 cations is AxA0 x
1ByB0 y
1O3
δ. Excellent reading materials describing further structural details of perovskite oxides can be found elsewhere (Burggraaf & Cot, 1996; Roy, 1953; J. Sunarso et al., 2008).

2.2 Transport phenomena in perovskite membranes Perovskite membranes are dense, and consequently transport phenomena are related to the crystal structure. The transport mechanism of ions and electrons is different and independent. In fact, a perovskite oxide that shows high oxygen conductivity can either exhibit high or low electronic conductivity. It is important to notice that the oxygen flux is proportional to the concentration of oxygen vacancies. Moreover, the presence of ionic defects, electrons, and electron holes is a function of the conditions in which the material is submitted. In the case of perovskite oxides, the most predominant conditions are temperature and oxygen partial pressure. Depending on the separation conditions, the perovskite can behave as a prevalent oxygen conductor or prevalent electron conductor. Therefore, the range of pressures and temperatures has to be detailed when working with perovskites as mixed ionic electronic conductors. Oxygen transport takes place by oxygen vacancy diffusion through crystal lattices, while electronic conduction is due to electron hopping between metal ions with different valence states in the B-site. Both transports take place simultaneously and, in the case of an oxygen separation membrane, they are counterflow. Fig. 8.3 depicts the case of an oxygen transport membrane, where the feed side is an oxygen-rich gas with high O2 partial pressure (such as air). On the opposite membrane surface, also called the permeate side in membrane applications, is an oxygen-depleted gas with low O2 partial pressure. This can be obtained by applying vacuum or a sweep gas (such as CO2) on the permeate side. At high temperatures, typically higher than 800°C, the presence of oxygen vacancies in the perovskite oxide, as well as the availability of electrons on the surface, provide the thermodynamic conditions for a surface reaction (Eq. 2) where the oxygen reacts with the electron and fills the oxygen vacancy as an O2
ion: 1 O + 2 e
! O2
2 2

(8.2)

1 O + 2 e
2 2

(8.3)

O2
!

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Fig. 8.3 Transport mechanism in MIEC membranes. From Athayde, D. D., Souza, D. F., Silva, A. M. A., Vasconcelos, D., Nunes, E. H. M., Diniz da Costa, J. C. & Vasconcelos, W. L. (2015). Review of perovskite ceramic synthesis and membrane preparation methods. Ceramics International. https://doi.org/10.1016/j.ceramint. 2016.01.130.

Diffusion of the O2
ions occurs through lattice diffusion. This mechanism involves oxygen anions transport between lattices where there are oxygen vacancies and thus is also known as a vacancy mechanism. The oxygen flux can be understood as a bulkdiffusion phenomenon where some approaches consider the oxygen vacancy as the mobile carrier (Lin et al., 1994). Concentration of oxygen vacancies on nonstoichiometry oxides is highly dependent on the temperature, in which higher temperatures lead to higher concentrations of defects. That is the reason for using perovskite membranes at temperatures typically higher than 800°C. On the other hand, this is also one of the major challenges for the widespread commercial use of perovskite membranes for oxygen production (D.D. Athayde et al., 2015), since the high temperatures require a challenging and careful design of the sealing components and materials. On the permeate side, the oxygen anions undergo the surface reaction detailed in Eq. (8.3), forming two electrons and the O2 that will be delivered pure to the permeate stream (vacuum or sweep gas). The formed electrons are now available to be transported in the opposite direction in counterflow with the oxygen ions. This transport takes place by two main mechanisms. The existence of oxygen vacancy on perovskite oxides (similar to the metal cations excess in nonstoichiometry oxides) establishes the conditions for ionization of neutral levels near the conduction band and the electrons are free to be transported in the conduction band of the oxide. This is called an n-type conductor. Moreover, the existence of B-site cations with different valence states contributes to the overall electronic conduction. A small polaron mechanism allows the hopping of electrons between these cations and is known as a hopping conductivity contribution. The overall transport phenomena also involve more steps other than the bulk diffusion of oxygen and electrons. As one may notice in Fig. 8.4A, oxygen transport

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Fig. 8.4 (A) Oxygen partial pressure profile during oxygen transport through a perovskite membrane and (B) the membrane thickness on the overall oxygen flux. From Sunarso, J., Baumann, S., Serra, J. M., Meulenberg, W. A., Liu, S., Lin, Y. S. & Diniz da Costa, J. C. (2008). Mixed ionic-electronic conducting (MIEC) ceramic-based membranes for oxygen separation. Journal of Membrane Science, 320(1–2), 13–41. https://doi.org/10.1016/j. memsci.2008.03.074.

occurs in three sequential steps. In interface I, the surface reaction (Eq. 8.2) takes place, followed by the simultaneous bulk diffusion of charged species, while in interface II the surface reaction (Eq. 8.3) releases the oxygen to the permeate stream. Therefore, the overall oxygen flux rate can be controlled by any of these three steps, since the slowest rate controls the overall transport. This behavior is clearly noticed when analyzing oxygen transport with varying membrane thickness. Typically, the decreasing membrane thickness of dense ceramic materials lowers the resistance to bulk diffusion, as shown by the oxygen transport of a perovskite material in Fig. 8.4B. For this scenario, the overall oxygen transport is controlled by bulkdiffusion. As the membrane thickness is further reduced, the rate of surface electrochemical reactions outgrows the bulk-diffusion rate, hence becoming the slowest stage. In this scenario, the overall transport is controlled by the reaction kinetics and the membrane thickness no longer has an impact on the overall oxygen flux rates. The threshold between the diffusion-controlled and kinetics-controlled regions is also known as the critical thickness (Lc). This value is extremely important for membrane design, as the manufacture of thinner membranes is desired, although this is limited to the critical thickness. Beyond that value, it is not worthwhile to obtain thinner membranes. Values of the critical thickness depend on the rates and kinetics of the various steps for each perovskite compound, ranging typically within 20–3000 μm (P.L. Rachadel et al., 2013). For instance, analysis of oxygen permeation through La0.2Sr0.5Fe0.8Ga0.2O3
δ membranes revealed that the critical thickness is larger than 1.30 mm (Etchegoyen et al., 2006), while for silver coated Ba0.5Sr0.5 Co0.8Fe0.175Y0.025O3
δ membranes the critical thickness is 0.40 mm (P.F. Haworth et al., 2012). Clearly, the critical thickness is also dependent on operation temperature, as it has a strong influence on the surface reaction kinetics as well as on the bulkdiffusion rates.

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2.3 Estimation of oxygen flux The equations used to estimate the oxygen flux are derivations of Fick’s first law. The limiting step of the process defines which equations can be used to estimate the oxygen flux, and comprehensive reviews of the available equations are found elsewhere (Burggraaf & Cot, 1996; J. Sunarso et al., 2008). The main equations for perovskite materials were compiled in Table 8.1. The equations are derived from the Wagner equation [Eq. 8.4 (Wagner, 1975)] that describes the oxygen flux (JO2) in the case where the diffusion is the slowest step. Perovskite compounds exhibit electronic conductivity (σ e) considerably higher than the ionic conductivity (σ i). Eq. (8.5) ( J. Sunarso et al., 2008) and Eq. (8.6) (Qi et al., 2000) were developed in this manner, considering that the Wagner equation involves only the ionic conductivity, which can be estimated as a function of the oxygen partial pressure (PO2). Based on an Arrhenius approach, Tsai et al. (1997) developed Eq. (8.7) from the Wagner equation (Tsai et al., 1997), which was successfully used for membrane unit modeling (Engels et al., 2010; Stadler et al., 2011). For the case where the surface reactions are the rate Table 8.1 Main equations to estimate oxygen flux through perovskite membranes. Limiting step

Equationa

Bulkdiffusion limited

JO2 ¼ 42RT F2 L

Observations R P0 O 2

σi σe P0 0 O 2 ð σ i + σ e Þ d  0  P JO2 ¼ 42RT σ ln P0 0OO2 F2 L i 2

ln PO2

0
n i JO2 ¼ RTσ
P00 O2
n Þ 4F2 L ðP O2 O

J O2 ¼ Cmem Surface reaction limited

a

T L

e

KWagner 

JO2 ¼ k (Cs
Cg)

T

 ln

P 0 O2 P0 0 O 2



–

Eq. (8.4)b

Constant value for ionic conductivity (e.g., small PO2 gradient) Assuming that the ionic conductivity is a function of oxygen partial pressure   σ i ¼ σ i O P0 O2
n Cmem and are intrinsic KWagner material constants k is a first-order rate constant and Cs and Cg are, respectively, the concentrations of oxygen ions in the solid state and gas phase

Eq. (8.5)c

Eq. (8.6)d

Eq. (8.7)e

Eq. (8.8)f

R is the universal gas constant, T the membrane absolute temperature, F the Faraday constant, L membrane thickness, P0 O2 and P00 O2 the oxygen partial pressure on the rich side (feed) and the oxygen lean side (permeate). Wagner (1975). c J. Sunarso et al. (2008). d Qi et al. (2000). e Tsai et al. (1997). f Chater et al. (1992). b

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limiting step, the equations assume a different aspect, as one may notice in Eq. (8.8) (Chater et al., 1992). There were also some attempts to create generalized transport equations ( J. Sunarso et al., 2008).

2.4 Perovskite compounds Most works on developing perovskite materials for oxygen production focus on compounds based on these mixed oxides: SrCoO3
δ, SrFeO3
δ, LaCoO3
δ, and LaFeO3
δ. As already mentioned, the oxygen permeation through perovskites is mostly determined by the ionic conductivity for thick membranes, and therefore great research effort has been focused on increasing this ionic conductivity by increasing oxygen vacancy concentration. The strontium-based compounds usually deliver higher oxygen fluxes, whereas the lanthanum-based compounds have been proven to be more adequate in such conditions. Due to the high oxygen permeation fluxes of strontium cobaltite and strontium ferrate compounds, the main development was focused on these categories of materials by partial substitution of the A-site and B-site lattices. For instance, the SrCo0.8Fe0.2O3
δ perovskite revealed an interesting compound in the early stages of the development of perovskite OTM. However, the cubic perovskite structure is not stable at temperatures lower than 790°C and low oxygen partial pressures. These drawbacks caused material degradation and membrane cracking due to the differential decomposition of the membrane between the two sides, one in contact with the feed (rich in oxygen) and the other in contact with the permeate or sweep current (lean in oxygen) ( J. Sunarso et al., 2017). In the early 2000s, other perovskite structures were developed by partial substitution of the strontium in the A-site lattice. A remarkable example is the partial substitution with barium cations (Ba2+), resulting in the Ba0.5Sr0.5Co0.8Fe0.2O3
δ (BSCF) perovskite (Shao et al., 2000). The developed material with the incorporation of a cation with a larger ionic radius and the same valence hindered the oxidation of B-site cations when submitted to the reducing atmosphere at high temperatures. It also delivered high oxygen fluxes, with a superior stability in the presence of CO2 when compared to SrCo0.8Fe0.2O3
δ. Another notable development based on SrCo0.8Fe0.2O3
δ was obtained by partial substitution of the A-site strontium with lanthanum cations (La3+), leading to the La1
xSrxCo0.8Fe0.2O3
δ (LSCF) class of materials (Prado et al., 2004). Although delivering lower oxygen fluxes when compared to the BSCF perovskite, the LSCF structure maintained higher chemical and mechanical stability when in contact with pure nitrogen atmospheres (PO2 < 10
5 atm). This specific condition of a pure nitrogen atmosphere is frequently used to compare the stability of novel perovskite compounds (P. Haworth et al., 2011), though the condition of low oxygen partial pressures is considerably harsher than the actual conditions used in oxygen permeation tests. The higher stability of the LSCF is due to the preservation of high oxygen content of the structure when submitted to high pressures, suppressing the formation of different phases (brownmillerite) and maintaining a lower degree of oxygen vacancies. Concerns over the oxidation of B-site cobalt in cobalt-based perovskites when submitted to reducing atmospheres have directed researchers to alternative strategies for

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material development. This is the main drawback of practical applications, such as when oxygen depleted atmospheres are used as sweep gas in the permeate side of a four-point membrane module. The strategy of substituting cobalt for B-site cations with a fixed valence state (e.g. Ga3+, Ti4+, and Zr4+) is an alternative, as well as the production of cobalt-free structures. Many cobalt-free structures have been synthesized through the last two decades, such as BaCeO3, BaCe1
yFeyO3
δ, La(Sr)Ga (Mg)O3, La1
xSrxGa1
yFeyO3, confirming the high stability when under reducing atmospheres. An extensive list of perovskite compounds and the many structural properties and characteristics were reported by J. Sunarso et al. (2008) and in a more recent review focusing on perovskite applications by J. Sunarso et al. (2017). The strategy of using dual-phase membranes to increase the performance of the OTMs has also resulted in interesting outcomes ( J. Sunarso et al., 2017). This strategy conventionally involves production of a membrane with two compounds. One material is responsible for electronic conduction while the other is highly oxygen permeable (such as fluorites). In order to avoid any limitation of the surface reactions on the triple-phase boundary (between the electronic conductor, the ionic conductor, and the gas phase), MIEC perovskites are used as the electronic conductor rather than materials that can only conduct electrons. Interestingly, the use of MIEC perovskites for both phases has also been explored, with remarkable stability of the dual-phase membrane (Fang et al., 2015). More recently, a similar approach was done by incorporation of catalyst phases on a perovskite matrix, creating a composite material (Leo et al., 2017; Ma et al., 2019).

3

Membrane synthesis and preparation methods

The synthesis routes of perovskite compounds and methods used for membrane preparation and conformation have been discussed in detail elsewhere (D.D. Athayde et al., 2015; J. Sunarso et al., 2008). This topic aims at providing a concise list of techniques and methods, as well as recent advances regarding the subject.

3.1 Membrane synthesis There have been many techniques reported for perovskite oxide synthesis. The distinctions in the final compound vary mainly on purity and particle size and agglomeration, which has a strong impact on the membrane preparation and performance. Choosing the appropriate method is a challenging process and requires a certain notion of the main desired characteristics of the final perovskite compound. D.D. Athayde et al. (2015) have compiled the main methods and their details, as displayed in Table 8.2 (D.D. Athayde et al., 2015). The solid-state reaction is the most conventional method for perovskite synthesis. The metal precursor powders (oxides, carbonates, or salts) are mixed together in the adequate proportion. Proper mixing of the powders is vital in order to obtain a certain degree of homogeneity. The mixed precursors are then submitted to thermal treatment, typically at high temperatures (1100–1400°C). The reactions take place at

Table 8.2 Perovskite oxide main synthesis methods. Synthesis method

Particle size

Agglomeration

Purity

Precursors

Solid-state reaction

>1000 nm

Moderate

Poor

Coprecipitation

>10 nm

High

High

Sol-gel (Pechini method) Hydrothermal

>10 nm

Moderate

Excellent

Oxides, hydroxides, and inorganic salts Hydroxides, inorganic salts, organic salts, and alkoxides Nitrates

>100 nm

Low

Very high

Hydroxides and salts

Spray and freeze-drying Spray pyrolysis

>10 nm

Low

Excellent

>10 nm

Low

High

Organic and inorganic salts Inorganic salts

>100 nm

Low

Excellent

Microwaveassisted synthesis

Oxides or inorganic salts

Calcination temperature

Observations

1100–1400°C

Laborious mechanical milling and mixing of the precursors required

800°C

Washing steps may cause deficiency of specific cations Accurate control of the final material composition

800–1000°C

Does not require posterior calcination 900°C 700–1000°C

600–800°C

Requires high pressures (up to 15 MPa) inside the pressure vessel

High control of the particle morphology Final particles are typically spherical and agglomerate-free, with a high specific surface area A fast calcination of the material is carried out avoiding particle coarsening

Adapted from Athayde, D.D., Souza, D. F., Silva, A. M. A., Vasconcelos, D., Nunes, E. H. M., Diniz da Costa, J. C., & Vasconcelos, W. L. (2015). Review of perovskite ceramic synthesis and membrane preparation methods. Ceramics International https://doi.org/10.1016/j.ceramint.2016.01.130.

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the interface of the particles through diffusion of the ions from the bulk to the interface, requiring longer reaction periods. The typical result is a material with low purity due to the presence of unreacted compounds. The particles are usually large (>1000 nm) and agglomerated, requiring laborious mechanical milling after heat treatment in order to obtain a fine powder. Despite these drawbacks, the solid-state reaction is a fast and simple method for perovskite synthesis. A higher degree of mixing is achieved by the coprecipitation method, involving the preparation of a supersaturated solution containing the metal cations. This condition can be obtained, for instance, by the use of another solution called the precipitation agent. The precursors precipitate simultaneously and are easily separated from the solution, and the collected material is heat treated. Fine control of the solution parameters, such as mixing rate, pH, and concentration, leads to good homogeneity and purity of the perovskite compound. The main difficulty in this method is maintaining the adequate cation proportion when performing the washing steps after coprecipitation. The sol-gel is another interesting technique with a large variety of routes used for perovskite synthesis. The alkoxide route employs the metal alkoxides as precursors, delivering high purity compounds. For some cations (groups I and II) that present difficulty to obtain pure metal alkoxides, an alternative is using the alkoxide-salt route. Another sol-gel route that has gained intense focus is the ethylenediaminetetraacetic (EDTA)—citrate complexation route, also known as the Pechini method. Metal ions in solution form complexes with EDTA and citrate, followed by the addition of ethylene glycol to promote the polymerization. This process minimizes segregation of metal ions during decomposition and heat treatment, culminating in high purity and homogeneity. The fine control of the composition and the versatility of this route made it the most used process for preparing perovskite materials. Other methods have also been used for perovskite synthesis, although they are not widely used for perovskite materials. The hydrothermal method is a sol-gel-type method in which the synthesis process takes place in autoclaves at critical conditions (high pressures and temperatures). Spray drying involves fast vaporization of the solvent in small droplets, whereas freeze-drying explores the slow sublimation of the solvent, both involving minimal segregation and intimate mixing of the ions. As for spray pyrolysis, the precursor droplets come in contact with a high temperature flame, obtaining high drying rates at the same time as the precursors are pyrolyzed. Finally, microwave assisted synthesis takes advantage of the ability of the microwave radiation to penetrate into the sample bulk, reaching high heating rates. It is noteworthy that some of these methods are still under development and further details are found elsewhere (D.D. Athayde et al., 2015; J. Sunarso et al., 2008). Other attempts have been made for perovskite synthesis, such as combustion synthesis and physical vapor deposition (PVD) technique, though they are at very early stages of development.

3.2 Membrane preparation/conformation The perovskite powders are then processed to achieve the desired membrane geometry. Sintering of the shaped membrane is performed at high temperature, providing the conditions for densification and shrinkage of the particles. The particle size

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275

distribution plays an important role in this stage, as the driving force for the sintering process is the decrease of the surface area. Moreover, the presence of impurities can strongly affect membrane performance due to the presence of pores or secondary phase formation. Final properties of the membrane, such as the porosity, relative density, surface morphology, and grain size, are determined at this stage. For applications such as OTMs, the perovskite membrane is dense in order to guarantee complete separation of the O2 from air. Therefore, choosing the appropriate membrane preparation method is paramount for obtaining dense membranes with high selectivity. The simplest method involves pressing the perovskite powder into disks with thickness varying around 1–2 mm. In dry pressing, adjustments to the green body density can be easily done by controlling the applied pressure and application time. In some cases, vacuum can also be applied to remove any entrapped air. Conventionally, other methods can also be used, such as tape casting and extrusion. Table 8.3 compiles the

Table 8.3 Major advantages and disadvantages of the main membrane preparation method. Membrane preparation methods Dense disc membranes

Advantages

Drawbacks

–

–

Easiest method to prepare perovskite membranes

– – Tape cast

–

– – Extrusion

– –

Production of flat thin (0.150 mm) and dense films that are pressed together to form the membrane (1 mm); The membrane exhibit expressive microstructure homogeneity; Precise control of the film thickness Fine control of the tube dimensions; Does not require exact match of the thermal expansion of the membrane and the pressure vessel, since only one side is attached to the pressure vessel, allowing free expansion of the tube at high temperatures

– –

–

– –

Production of thick membranes (1–2 mm); Low oxygen fluxes are obtained; Increased sealing areas Lower membrane thicknesses result in brittle membranes; Difficulty to prepare gastight modules due to the high sealing area

Requires tailored rheological properties in order to produce a symmetric final shape; Two sintering steps are necessary; Production of tubes with thick walls (0.4–2.0 mm), reducing the oxygen fluxes

Continued

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Perovskite Ceramics

Table 8.3 Continued Membrane preparation methods

Advantages

Drawbacks

Hollow fibers

–

–

– –

– Asymmetric thin dense films on porous substrates

– – –

Freeze-cast

– – –

– –

Flat plate wafers

– –

Production of membranes with thin walls (0.200– 0.300 mm); Higher oxygen fluxes when compared to other tubular geometries; The thin external diameter allows high packing densities, leading to high oxygen production per unit volume; Highest oxygen fluxes reported to date Possibility to use very thin (0.020–0.065 mm) dense layers; The porous substrate enhances the mechanical stability of the membrane; The method allows surface modification with porous ceramic catalyst Production of mechanical stable supports with high porosity (above 60%); Formation of an ordered pore network; No significant resistance for gas diffusion through the support, thus, low-pressure drop through the porous substrate; Possibility to use very thin dense layers; The use of water as solvent makes this an environmentally friendly process The design allows an easy scale-up; Reduced sealing area per membrane area;

– – –

– –

– –

–

–

Reports indicate difficulties in obtaining a doping mixture with proper rheological conditions; May exhibit not interconnected pores; They are brittle and tend to break; The internal diameter is very small (99.9%) revealed the potential of these membranes for commercial utilization (Anderson et al., 2016). Another practical aspect of oxygen separation using perovskites is the sealing of the membrane contacts that can withstand high operating temperatures. The contact between the ceramic membrane and metal support is troublesome when working at high temperatures, due to the different thermal behavior. The use of tubular or hollow fibers decreases the sealing area and is more

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287

Fig. 8.11 Schematic representation and picture of a membrane module. From Nauels, N., Herzog, S., Modigell, M. & Broeckmann, C. (2019). Membrane module for pilot scale oxygen production. Journal of Membrane Science, 574(December 2018), 252–261. https://doi.org/10.1016/j.memsci.2018.12.061.

adequate to hinder any sealing leakage. Nevertheless, the membrane area that is able to take part in the oxygen separation process is decreased due to the area used for proper insulation (only 74% of the membrane is actually used (Nauels et al., 2019)).

4.2 Perovskite membranes for oxy-fuel combustion The main application of the membranes discussed in this chapter is oxygen production, as it has been accomplished by Air Products and many other pilot plants discussed in the text. Other applications have also been envisioned for the perovskite membranes, in which oxygen transport is desirable. The main application that has been analyzed (mainly by simulations) is oxy-fuel combustion. This concept aims to reduce CO2 emissions in thermal power plants, such as the ones based on coal combustion. In oxy-fuel combustion, the fuel combustion takes place with pure oxygen instead of using air. Consequently, there is no N2 or NOX in the flue gas, which consists mainly of H2O and CO2. The CO2 is more readily available, as the separation of these two gases is considerably simpler than when N2 or NOx is present. The use of cryogenic air separation to produce the pure oxygen stream involves high energy consumption, leading to a reduction of 8%–12% (in percentage points) plant efficiency. Meanwhile, the use of perovskite membranes has been discussed as the best alternative for this application, as it leads to lower efficiency loss. When comparing the threepoint and the four-point membrane modules (Fig. 8.12), the four-point system leads to higher net efficiency (Engels et al., 2010; Stadler et al., 2011). However, the hot flue gas is used as a sweeping gas in this concept, and hence the perovskite must be chemically stable when in contact with H2O and CO2, as well as other minor components. For instance, a pilot scale of the oxy-fuel combustion system has recently been trialed using an LSCF membrane and the thermochemical characterization of the perovskite material showed poor stability, as it reacted with components (mainly sulfur, but also with Cr and Zn) present on the flue gas (Portillo et al., 2020). Other applications of perovskite oxides are solid oxide fuel cells (SOFCs) and membrane reactors. These applications take advantage of the ability of perovskite

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Perovskite Ceramics

(A)

(B) combustion

flue gas

combustion 6

7 100 % O2 MIEC membrane N2

6

1 heat exchaner 2 heater/boiler flue gas 3 compressor 4 turbine 5 vacuum pump 6 recirculation fan 7 cooler N2

5

MIEC membrane 1 4

2

3

4

air 3 air

Fig. 8.12 Perovskite membrane integration in an oxy-fuel power plant, showing (A) the threepoint module and (B) the four-point module. From Engels, S., Beggel, F., Modigell, M. & Stadler, H. (2010). Simulation of a membrane unit for oxyfuel power plants under consideration of realistic BSCF membrane properties. Journal of Membrane Science, 359(1–2), 93–101. https://doi.org/10.1016/j.memsci.2010.01.048.

materials to conduct simultaneously electrons and anionic oxygen. A thorough discussion and literature review on these three applications has been published by J. Sunarso et al. (2017).

5

Future perspectives

The use of perovskite membranes for oxygen separation from air has been extensively studied since 1985, passing through the many stages of technology development (material development, lab-scale tests, membrane geometry, and scale-up), as reported by D.D. Athayde et al. (2015). Since then, further development has been done to overcome the three main challenges to the long-term commercial use of perovskite membranes: l

l

l

increase the oxygen flux; higher chemical stability; membrane reliability and mechanical stability.

These three aspects have been addressed throughout this text, detailing the strategies to overcome each one. Even with the numerous reports on this subject, the practical aspects of using these membranes still require further assessment for widespread commercial use. As listed by J. Sunarso et al. (2017), further improvements aimed at membrane material and fabrication and process design are necessary. In addition, some

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external factors involving CO2 penalties and more rigid energy legislation could stimulate advancements in this technology ( J. Sunarso et al., 2017).

6

Final remarks

The perovskite membrane technology has economical and environmental benefits when compared to conventional methods for oxygen production. Description of the perovskite oxide structure allowed an understanding of its properties and the transport phenomena involved during the oxygen transport through the membranes. Meanwhile, many perovskite compounds were presented for analysis of the influence of material properties on membrane transport. Based on these topics, a thorough compilation of the membrane synthesis and membrane preparation was displayed, as well as exploration of the benefits and drawbacks of each method. In addition, membrane performance is discussed, giving special attention to the impressive improvements obtained over the years since perovskite was first reported as an OTM for oxygen production in 1985. Finally, the challenges that are yet to be more deeply addressed by the researchers and engineers for widespread commercial use of this technology were explored. Providing efficient oxygen production for long periods with a reliable perovskite membrane is paramount for the next steps in technology deployment. A few examples of pilot and small-scale trials using perovskite membranes were shown in this chapter, with promising results for oxygen production. Nevertheless, further investments are still required to obtain an efficient, reliable, and reproducible technology. Moreover, further studies on the use of perovskite membranes for cleaner energy delivery would be a motivating approach for increasing the interest of the community in this technology.

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novel tubular freeze-cast substrates. Materials Research, 22(2). https://doi.org/10.1590/ 1980-5373-mr-2018-0639. Athayde, D. D., Motuzas, J., Diniz da Costa, J. C., & Vasconcelos, W. L. (2020). Novel two-step phase inversion and dry surface coated carbon membranes on alumina freeze-cast substrates for desalination. Desalination, 114862. https://doi.org/10.1016/j. desal.2020.114862. Athayde, D. D., Sousa, B. M., Dolabella, A. C. A., de Ribeiro, J. O. N., Vasconcelos, D. C. L., da Costa, J. C. D., et al. (2020). Manufacture of highly porous tubular alumina substrates with anisotropic pore structure by freeze-casting. Advanced Engineering Materials. https://doi. org/10.1002/adem.201901432. Athayde, D. D., Souza, D. F., Silva, A. M. A., Vasconcelos, D., Nunes, E. H. M., Diniz da Costa, J. C., et al. (2015). Review of perovskite ceramic synthesis and membrane preparation methods. Ceramics International. https://doi.org/10.1016/j.ceramint.2016.01.130. Athayde, D. D., Souza, D. F., Silva, A. M. A., Vasconcelos, D., Nunes, E. H. M., Motuzas, J., et al. (2017). Carbonation passivation layer of scandium loaded BSCF perovskite. Ceramics International. https://doi.org/10.1016/j.ceramint.2017.08.049. Baumann, S., Meulenberg, W. A., & Buchkremer, H. P. (2013). Manufacturing strategies for asymmetric ceramic membranes for efficient separation of oxygen from air. Journal of the European Ceramic Society, 33(7), 1251–1261. https://doi.org/10.1016/j. jeurceramsoc.2012.12.005. Baumann, S., Serra, J. M., Lobera, M. P., Escola´stico, S., Schulze-k€ uppers, F., & Meulenberg, W. A. (2011). Ultrahigh oxygen permeation flux through supported Ba0.5Sr0.5Co0.8Fe0.2O3
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Perovskite lead-free dielectric ceramics: Highly promising materials for energy storage applications

9

Mrinal Kanti Adaka,b and Debasis Dhaka a Department of Chemistry, Sidho-Kanho-Birsha University, Purulia, India, bDepartment of Chemistry, Indian Institute of Technology Delhi, New Delhi, India

1

Introduction

Global warming and exhaustion of fossil fuels proclaim loudly the issue of energy for the existence of life on Earth. As effective energy storage devices, fuel cells, lithium and sodium ion batteries, electrochemical supercapacitors, and electrostatic capacitors are explored worldwide to fill up the gap for renewable energy storage (Benyoussef et al., 2018). In modern electronic devices, battery and capacitors are vital components and are rapidly used in electronic power systems (Chandrasekhar et al., 2019). Owing to high-power density, excellent fatigue performance, intrinsic ultrafast chargedischarge capability (500°C). Also, other lead-free compositions with perovskite-like related structures, like tungsten-bronze, should be considered as candidates though few studies have been carried out on them (L. Pardo et al., 2018).

4.1 Alkaline niobates The ferroelectricity of the KNbO3-NaNbO3 (KNN) solid solution dates back from the 1950s (Shirane et al., 1954) exhibiting a complete solid solubility. The KNN phase diagram shows three phase boundaries at x ¼ 0.17, 0.35, and 0.5; however, most of the studies concentrate around the x ¼ 0.5 value since this composition, the MPB, exhibits a moderate dielectric constant and an optimum piezoelectric response (Garroni et al., 2018; Taghaddos et al., 2015; Z. Zhang et al., 2018; Y. Zhang & Li, 2019). Pure KNN ceramics are characterized by having a low densification and off-stoichiometry induced by ion volatilization during the densification process (sintering). Moreover, one of the major drawbacks of this composition is the limited thermal stability that can be achieved because of the onset of a phase transition between an orthorhombic and tetragonal polymorph. Soft-chemistry routes are succeeding processes for preparing KNN phases (Garroni et al., 2018; VillafuerteCastrejo´n et al., 2016). These yields nanosized powders (Lo´pez et al., 2010) and avoid potassium or sodium losses (Lo´pez-Jua´rez et al., 2014). Simultaneously, appropriate doping can extend the thermal stability range, produce highly mechanical quality factors, enhance the piezo-/ferroelectrical activity, or increase the defect dipole

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325

concentration (Y. Zhang & Li, 2019). Because of their important electrical characteristics (Carren˜o-Jim
enez et al., 2018; S. Zhang et al., 2009), this type of lead-free piezoceramics have been already implemented in actual electronic devices such as ultrasonic transducers for medical imaging (Taghaddos et al., 2015; Z. Zhang et al., 2018), wire bonding (Mathieson & De Angelis, 2016), and simple sensors (R. Castan˜eda-Guzma´n et al., 2017).

4.2 Barium titanate-based compositions Barium titanate, BaTiO3 (BT), remains from its early development (Hippel, 1950; Randall et al., 2004) as one of the most widely used and studied piezo-ferroelectric and dielectric materials (L. Pardo et al., 1999, 2018; Villafuerte-Castrejo´n et al., 2016). Concerning the compositions based on it, as a pseudobinary solid-solution system, (1
x)Ba(Zr0.2Ti0.8)O3-x(Ba0.7Ca0.3)TiO3 (BCTZ) features exceptional piezoelectric coefficients (Liu & Ren, 2009), specifically for compositions lying at the isopleth boundaries of BaTiO3-BaZrO3-CaTiO3. However, this composition is not amenable to commercial processing due to the low Curie temperature (93°C for x ¼ 0.50) and the high synthesis and sintering temperatures required for their production (>1300°C). Doping of the BCTZ solid solution has been primarily investigated to reduce the temperature of the thermal processes and improve the functional properties of the material, resulting in promising solid solutions with important electromechanical properties (Liang et al., 2015; Xue et al., 2011; Zhou et al., 2012). Alternative processing routes employing reactive precursors have been used successfully to produce high sensitivity, dense, fine-grained materials with relatively low sintering temperatures (A. Reyes-Montero et al., 2014). The composition with x ¼ 0.5 in the above-described pseudobinary BCTZ solid solution has been studied thoroughly (Acosta et al., 2014; Ehmke et al., 2012; Keeble et al., 2013; Liu & Ren, 2009). From the structural point of view, it commonly includes the presence of a rhombohedral (R, R3m) and a tetragonal (T, P4mm) phase, with an intermediate region that was initially proposed as MPB (with a tricritical point at 57°C for x ¼ 0.32). However, it has been found that the temperature dependence suggests a PPT-like nature (Keeble et al., 2013). In a vertical MPB behavior, the properties are only composition dependent and thus thermally stable, while for a PPT the structure is strongly determined by temperature. Later studies have revealed an intermediate region between R and T phases (Acosta et al., 2014; Bjørnetun Haugen et al., 2013; Keeble et al., 2013; L. Zhang et al., 2014), with a bridging orthorhombic phase (O, Amm2). The synergy of compositionally driven polymorphic and electric fielddriven structural phase transitions has given a new dimension to the poling process in lead-free materials (A. Reyes-Montero et al., 2021). The difficulty of fabricating large enough uniform lead-free piezoelectric materials with high piezoelectric coefficients has been preventing to date to make a true medical imaging ultrasonic array transducer with center frequency 0) for an adiabatic polarization/magnetization or negative (Δ TE,H < 0) for an adiabatic depolarization/demagnetization, respectively. This last makes this effect very attractive for solid-state refrigeration (M. M. Vopson et al., 2017). The refrigerant is a bilayer laminate composite (see Fig. 11.10B); the layered structure has a surface area A and consists of a piezo-ferromagnetic layer (M), a piezo-ferroelectric layer (E), and two metallic electrodes (C1 and C2). In order to maximize the multicaloric effect, the temperature of work T must occur near the Curie temperature of transition E of the ferroics phases (i.e., T  TH C  TC). However, for the ME coupling measurements different transition temperatures of the ferroics phase are required. Replacing E ∂ M/∂ T ¼ γ H and ∂ P/∂ T ¼ γ E and assuming TH C > TC, at the temperature of work E T close to one of the transition temperatures, i.e., T ¼ TH C ¼ TC; thus, at ∂ M/∂ H T ¼ γ ¼ 0. This result will modify Δ TE and Δ TH. The product ratio will lead to obtaining the magnetically induced ME coupling coefficient, αH ¼ ε0χ E. (Δ TH/Δ TE). E (Δ E/Δ H). Likewise, when TH C < TC and choosing the work temperature H E T ¼ TC ¼ TC will result in ∂ P/∂T ¼ γe ¼ 0, obtaining thus the electrically induced ME coupling coefficient, αE ¼ μ0χ H. (Δ TH/Δ TE). (Δ H/Δ E) (Fig. 11.10). Experimental measurement confirmation of ME coupling with a multicaloric effect was affirmed in single-phase materials and ferromagnetic and ferroelectric composite materials (Ursic et al., 2016).

2.3 Factors influencing in the ME coupling response The following section covers the fundamental aspect of composites with different connectivities and mechanisms that influence multiferroic behavior.

360

Perovskite Ceramics

2.3.1 Effects of interface The interface is the critical factor in determining the composites’ comprehensive properties. The application of composite types both in thin films and in bulk has been a promising approach for improving ME coupling performance to reduce the leakage current problem and maximize the interfaces between phases. Different strategies have been suggested to control the interfaces, interfacial interdiffusion, and undesired reactions in thin film and in bulk, giving rise to distinct interface architectures, as mentioned before. The types of interfaces in thin films are frequently characterized by the phases and their orientation with substrates. In bulk they are characterized by the phases and volume fraction of phases (Zhou et al., 2017). In this last, other factors are essential, such as difference in the effect particle size of the phases, sintering temperature, and intergrain effects. Fig. 11.11 shows a model of defects at the particulate composite interface during the sintering process. This proposed model is based on the evidence offered in several reports (F. D. Ma et al., 2014). We propose: (i) the singlephase particle sizes are uniformly distributed with approximate size, (ii) the particles during the sintering process have an electronic transfer process, (iii) at high temperatures, the ions present will produce an interdiffusion to the other particles along the grain boundary, and (iv) chemical reactions can trigger the formation of secondary phases in the multiferroics material. A strategy is filling the intergrain space with an excellent insulating material that may increase the bulk resistivity of the samples up to several orders of magnitude (Fig. 11.11). Badwal et al. (1993), through thermodynamic calculations, demonstrated that LayMnO3 can coexist with other phases when the occupation of site A, y, is less than 0.86. van Roosmalen et al. (1993) showed that the sintering temperature increases with Sr content in (La,Sr)MnO3; in addition, punctual defects (oxygen vacancies) decreased with increasing temperature. The control of this parameter is beneficial since it helps to understand the intergrain electric/magnetic interactions in the materials. Besides, it helps to enhance the ME coupling coefficient through structural modification, breaking of incommensurate spin structure, and intergrain magnetostatic Intergrain effects A-grain

B-grain XY+

ZT+

SQ+

EL+

Fig. 11.11 Interfacial diffusion between A and B phases in composites.

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interaction. In the case of single-phase multiferroics, determining the type and concentration of point defects in the surface and impurity in an intergranular space is an essential consideration in the ME coupling coefficient, for example, Bi0.9La0.1FeO3δ multiferroics (Pashchenko et al., 2019). Zhai et al. (2018) report a single-phase multiferroic hexaferrite Sr3Co2Fe24O41. This maximum ME coefficient appears around 100 Oe with a value of 15 mV cm1 Oe1, without defects in the surface and impurities in the grain boundary. To optimize single-phase materials in Bi10Fe6+yTi3 yO30+δ (Bi10-Fey, y ¼ 0, 0.2, 0.4) samples and ensure that the nonstoichiometry in the reaction system does not cause structural disorder in the obtained materials, fix the Fe/(Fe + Ti) ratio and slightly increase the Bi content in the Bi10Fe0 samples was slightly increased (G. Wang et al., 2018). Contrarily, a numerical estimation shows the origin of the giant linear ME effect in single-phase BiFeO3. The study is based on the symmetry approach of the Ginzburg-Landau type (Popkov et al., 2016). This approach considered that the polarization, antiferrodistortion, and antiferromagnetic momentum vectors are viewed as ordering parameters. For a particulate composites the ions produce interdiffusion to other particles along the grain boundary; in case of laminated composites, in addition to the grain boundary, interdiffusion is extended along with the entire interface. The representation of the interdiffusion process along the interface is shown in Fig. 11.12. Understanding the interdiffusion mechanism as a function of distance is essential in composites due to the weak value of the ME coupling coefficient. The interdiffusion kinetics, i.e., the speed of this interdiffusion process, is triggered by the temperature T according to Arrhenius’ law, D(T) ¼ D0e  Δ E/RT, where D is an intrinsic parameter of the material, R ¼ 8.314 J mol1 K1 is the gas constant, and Δ E is the apparent activation energy of the sintering process. In the initial phase, the interdiffusion of ions and defects between phases leads to product formation at the interface. Furthermore, the product formed at the interface has a different stoichiometry compared to the individual phases, consequently changing the structural parameters. This change in the lattice parameter will induce strain in the interface leading to fracture or change in the properties of the individual phases. After sintering (Fig. 11.12A) when the interdiffusion mechanism between the layers is insignificant, a sharp drop in the concentration of the A layer in the B layer is observed. We can infer that the interdiffusion will depend on the reaction at the interface, i.e., between the phases through the grain boundary. However, the chemical reactions will depend on the composition of each ferroics phase. In the presence of chemical defects, the interdiffusion process across the grain boundary is significant (see Fig. 11.12B). Percolation between layers in the initial process, in addition to chemical defects, facilitates interdiffusion between phases. This process leads to a slow drop in the concentration of the A layer in the B layer. When percolation is total, in addition to chemical defects, between phases the interdiffusion process can be significant along the interface (see Fig. 11.12). J. L. Clabel H. et al. (2016, 2014) considered the effect of grain size in La0.7Ba0.3MnO3 (L)-BaTiO3 (B) 2-2 type composite and claimed that grain boundary diffusion is activated at higher temperatures than surface diffusion (see Fig. 11.12D). When nanometer-sized precursor powders have been used, the contribution of surface

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Fig. 11.12 Mutual diffusion mechanism between the A and B phases (A) there is no diffusion, (B) partial diffusion, (C) total diffusion, (D) distribution of elements across the boundary of La0.7Ba0.3MnO3 (layers L)-BaTiO3 (layers B) 2-2 type ceramic composite, the vertical dotted line indicates the average boundary , (E) SEM analysis result of the representative 0.9LKNN/ 0.1CZFM laminated composite, and (F) element distribution across the boundary, where the upright dotted line indicates the average boundary. (A)–(D) From Clabel H., J. L., Zabotto, F. L., Nogueira, I. C., Schio, P., Garcia, D., De Lima, O. F., Leite, E. R., Moreira, F. M. A., & Cardoso, C. A. (2014). Magnetoelectric properties of laminated La0.7Ba0.3MnO3–BaTiO3 ceramic composites. Journal of Magnetism and Magnetic Materials, 364, 18–23. https://doi.org/10.1016/j.jmmm.2014.04.014 and Clabel H., J. L., Ferri, F. A., Zabotto, F. L., Rivera, V. A. G., Nogueira, I. C., Garcia, D., de Lima, O. F., Leite, E. R., Pereira-da-Silva, M. A., & Cardoso, C. A. (2016). Grain size and interfacial interdiffusion influence on the magnetic and dielectric properties of magnetoelectric La0.7Ba0.3MnO3– BaTiO3 composites. Journal of Magnetism and Magnetic Materials, 407, 160–166. https://doi. org/10.1016/j.jmmm.2016.01.082, with the permission of Elsevier. (E) and (F) From Yang, H., Zhang, J., Lin, Y., & Wang, T. (2017). High Curie temperature and enhanced magnetoelectric properties of the laminated Li0.058 (Na0.535K0.48)0.942NbO3/Co0.6Zn0.4Fe1.7Mn0.3O4 composites. Scientific Reports, 7. https://doi.org/10.1038/srep44855, with the permission of Springer Nature.

diffusion to the densification process cannot be ignored because it implies the presence of large surface areas that affect the development of the particle/pore structure during sintering. The grain growth takes place in the final stage of the sintering of micron-sized powders. However, in nanometer-sized powders, grain growth coincides with the densification from the beginning of the sintering process. Yang et al. (2017) fabricated laminated composites of Li0.058(Na0.535K0.48)0.942NbO3 (LKNN)/

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Co0.6Zn0.4Fe1.7Mn0.3O4 (CZFM) prepared by the conventional solid-state sintering method, showing from EDS analysis at the interface of the two phases. It can be found that a bit of interdiffusion occurs at the interface of the two phases (see Fig. 11.12E and F). In both laminated composites, it is shown that the solid-state reaction method presents low interdiffusion at the interface. Composites in the thin film have been considered a promising approach for improving the ME coupling coefficient by reducing the leakage current and the operating temperature. Different strategies have been suggested to control thin-film interfaces, including distinct interface architectures. Single-phase materials are interfaces with grain boundaries defined by the misorientation of two or more crystallographic planes from adjoining grains in polycrystalline solid thin films ( J. Ma et al., 2011). The grain boundary density can be significantly enhanced in the thin films as the nanograins can be obtained by transforming a 3D bulk into a 2D thin film. There are thick, randomly oriented polycrystalline grains in bulk materials that could not provide a sufficient pathway for the mobile (oxide) vacancies. The use of thin films remarkably allows tuning the grain size and orientation of the interfaces, besides complex sharp piezoelectric-magnetostrictive interfaces, i.e., avoiding interfacial diffusion.

2.3.2 Effects of the volume fraction It is well known that a higher volume fraction of particles of one of the ferroics phases causes more strain, and as a result, tensile strength improves while ductility decreases. During strain, these interfaces are excellent sites for the initiation and propagation of cracks. The variation of the ferroic volume fraction leads to stress in the single-layer forming two cracks that will interact. The crack interactions between the ferroics phases will determine the growth of the cracks. The cracks are magneto-electrically permeable, and applying an electric (or magnetic) field will influence the crack tip due to the stress field, affecting thus the ME coupling. At the end of the crack, the stress applied at the boundaries is focused in a small region, and when the stress is greater than the tensile strength, the crack will rapidly grow and quickly traverse the entire sample (Cheng et al., 2018). The interfacial mechanical coupling in laminated ceramics can be divided into materials with strong and weak interfaces. The first can increase mechanical performance by a compressive layer, which can absorb residual stress (residual stresses depend only on the volume ratio). The second can significantly improve the fracture energy dissipation by delaminating at the interface, breaking the crack propagation path (Cheng et al., 2018) (Fig. 11.13). Fig. 11.13 shows a theoretical model in the magneto-electro-elastic interfacial (MEE) regions. The two cracks are called crack I and crack II, respectively. Their half lengths are a0(I) and a0(II), and the horizontal distance between their centers is dc. The abscissas of the four crack tips are aI, bI, aII and bII, respectively. The thicknesses of the piezoelectric (PE), piezomagnetic (PM), and MEE layers are ht1, ht2, and ht3 (Zhou et al., 2017). Studies in laminate multiferroics (S. L. Liu et al., 2017) reveal the effects of the volume fraction on the fracture behavior and show that: (i) the effects of the piezoelectric volume fraction reflect the effects in the generated magnetostrictive strain and the stress in the piezoelectric material across the interfaces, and (ii) the

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tyz(0) ht1

(I)

FE layer

2a0 aI

(II)

2a0

MEE interfacial region

aII

bI

ht3

bII

FM layer

dc

ht2

tyz(0) (0) D(0) y or By

Fig. 11.13 Fracture model of a multiferroic composite containing a Magneto-electro-elastic interfacial region. From Zhou, K., Li, Y. D., & Liu, S. L. (2017). Effects of the volume fraction of piezoelectric particles in the magneto-electro-elastic interfacial region on the fracture behavior of a laminate multiferroic plate. Acta Mechanica, 228(4), 1229–1248. https://doi.org/10.1007/s00707-0161763-6, with the permission of Springer-Verlag.

interfacial fracture is due to the competition between the effect of the magnetostrictive strain and piezoelectric stress, and the thickness of the interface. The ME coupling via strain mediation in composite 0-3 type materials implies a highly nonlinear magnetoelectric coupling coefficient as a function of the magnetic field (Newacheck & Youssef, 2022). In these systems, the volume fraction strongly influences the ME coupling coefficient. A comparison between the predicted direct ME coefficient response using the computational framework presented by Newacheck and Youssef (2022) and the analytical results of C.W. Nan et al. (2001) as a function of the volume fraction of the spherical Terfenol-D particles is given (see Fig. 11.14). The computational results elucidate the nonlinear dependence of the direct ME coefficient on the content of the conductive magnetic particles (Fig. 11.14).

2.3.3 Effect of size The effect of particle size on microstructure, mechanical properties, and ME coupling response in different composites types plays a crucial role. Zeng et al. (2015) fabricated the composites of the 0-3 type of Terfenol-D/PZT by a hot pressing method. The Terfenol-D and PZT powders have a mean particle size of 20 and 5 μm, respectively, and the authors have investigated the effect of particle size on the ME coupling coefficient. The difference in particle size has tremendous importance in stress transmission, which contributes to enhancing the magnetoelectric effect. The maximum magnetoelectric coefficient of 82.58 mV cm1 Oe1 at an applied magnetic field of 200 Oe has been obtained. R. Gao et al. (2020) investigated the effect of particle size of CoFe2O4 (CFO) and BaTiO3 (BTO) composite fluids on the magnetoelectric coupling effect (see Fig. 11.15). The composite of CFO and BTO particles will form a core-shell with a bi-phase crystalline structure, where the average particle size is

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Fig. 11.14 (A) Comparison of the analytical and computation direct magnetoelectric coupling coefficient concerning the volume fraction of the Terfenol-D phase. (B–H) Contour plots of the magnetic flux, strain, and electric potential distributions observed from the top surface of the REV, showing the effect of the volume fraction on the distribution of these parameters and the resulting DME. The results exemplify an optimal volume fraction for this type of magnetoelectric composites. From Newacheck, S., & Youssef, G. (2022). Microscale magnetoelectricity: Effect of particles geometry, distribution, and volume fraction. Journal of Intelligent Material Systems and Structures, 33(10) 1338–1348. https://doi.org/10.1177/1045389X211053053., with the permission of SAGE Publications.

28 and 35 nm, respectively. This different particle size of CFO and BTO nanocomposites is resulted by the difference of CFO particle size of 12 and 18 nm. The maximal values of the coefficient are 10.2 and 6.1 V cm1 Oe (Fig. 11.15B–D). These values are several orders of magnitude larger than that obtained in conventional ceramics reported in the literature (X. Z. Chen, Hoop, et al., 2017; Viana et al., 2022). The high ME coupling coefficient is due to particle size since larger particles need a larger magnetic field to switch their magnetization direction because of the more significant drag coefficient (Fig. 11.15). Lisnevskaya et al. (2016) prepared magnetoelectric composites from the sol-gel and solid-state methods from PZTNB-1 and YIG powders. Composites with connectivities of 3-3, 3-0, and 0-3 were processed with particle size >3.0 and