Microstructure, Property and Processing of Functional Ceramics

Qingrui Yin Binghe Zhu Huarong Zeng Microstructure, Property and Processing of Functional Ceramics Qingrui Yin Bing...

Author: Qingrui Yin | Binghe Zhu | Huarong Zeng

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Qingrui Yin Binghe Zhu Huarong Zeng

Microstructure, Property and Processing of Functional Ceramics

Qingrui Yin

Binghe Zhu Huarong Zeng

Microstructure, Property and Processing of Functional Ceramics With 212 figures

AUTHORS Prof. Qingrui Yin Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai 200050, China E-mail: [email protected] Prof. Binghe Zhu Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai 200050, China E-mail: [email protected] Associate Prof. Huarong Zeng Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai 200050, China E-mail: [email protected] Based on an original Chinese edition:

ǉࡳ㛑䱊⫋ⱘᰒᖂ㒧ᵘǃᗻ㛑Ϣࠊ໛ᡔᴃǊ(Gongneng Taoci De Xianwei Jiegou, Xingneng Yu Zhibei Jishu), Metallurgical Industry Press, 2003

ISBN 978-7-5024-4571-3 ISBN 978-3-642-01693-6 e ISBN 978-3-642-01694-3

Metallurgical Industry Press, Beijing Springer Dordrecht Heidelberg London New York Springer Dordrecht Heidelberg London New York

Library of Congress Control Number:2009930028 This work is subject to copyright. All rights are reserved, whether the whole or part of the material is concerned, specifically the rights of translation, reprinting, reuse of illustrations, recitation, broadcasting, reproduction on microfilm or in any other way, and storage in data banks. Duplication of this publication or parts thereof is permitted only under the provisions of the German Copyright Law of September 9, 1965, in its current version, and permission for use must always be obtained from Springer-Verlag. Violations are liable to prosecution under the German Copyright Law.

ӄ 2009 Metallurgical Industry Press, Beijing and Springer-Verlag GmbH Berlin Heidelberg Co-published by Metallurgical Industry Press, Beijing and Springer-Verlag GmbH Berlin Heidelberg Springer is a part of Springer Science+Business Media Springer.com The use of general descriptive names, registered names, trademarks, etc. in this publication does not imply, even in the absence of a specific statement, that such names are exempt from the relevant protective laws and regulations and therefore free for general use. Cover design: Frido Steinen-Broo, Estudio Calamar, Spain Printed on acid-free paper

Preface

The functional ceramic materials (FCM) are potential for use in many electronic devices such as optical waveguides, non-volatile dynamic random access memories, micromotors, microactuators, thin film capacitors, and pyroelectric infrared detectors. FCM possesses unique properties like piezoelectricity, pyroelectricity, photoelectricity, photo-acoustic effect, photorefractive behavior, and non-linear optical activity that are closely depends closely on the common theme of composition-preparation-structure-property relationships in the solid state, especially microstructures (grain, grain boundary and domain structures, etc.) and their dynamic response to mechanical, electrical and optical loads at nanometer scale. Thus it is very important to understand the physical phenomenological behavior of ferroelectric structures and their dynamic evolution in nanoscale volumes. This is the context that motivated the publication of this book. The aim of this book is to present recent advances in the fabrication process of functional ceramic materials and their property study, particularly, in-depth observation/analysis of microstructures using the custom-built scanning electron acoustic microscopy (SEAM), acoustic and piezoresponse mode scanning probe microscopy based on atomic force microscopy. Along with the generally accepted concepts and experimental results there are numerous applications of functional ceramics and devices in industry. We hope that this book will make the readers aware of tremendous developments in the field of microstructure characterization and functional ceramic preparations. The first two chapters address fundamentals of microstructures in the functional ceramics. Chapter 1 presents the formation mechanism of microstructures including grains, grain boundaries, pores, domain structures, and their correlations with properties and processing for some typical ceramics like PLZT (lead lanthanum zirconate titanate) ceramics, PTC (positive temperature coefficient) ceramics, piezoelectric ceramics, ferroelectric ceramics, and so on. Chapter 2 discusses grain boundary phenomena such as grain boundary segregation and migration in the functional ceramics. The next two chapters focus on near-field microscopy characterizing ferroelectric domains of functional ceramics. Chapter 3 describes two custom-built acoustic microcopies used for ferroelectric domain imaging, scanning electron acoustic microcopy (SEAM) and scanning probe acoustic microscopy (SPAM), presenting

Ď Preface

their operation principle, analyzing domain contrast formation mechanism, showing their high-resolution observations of domain configurations non-destructively, especially their unique capability of imaging sub-surface structures, and their dynamic response to the electric field as well as their applications to a variety of materials including structure ceramics, metals, single crystals, composites and coatings, etc. Chapter 4 presents piezoresponse force microscopy (PFM) of nanoscale domains in functional ceramics, involving interpreting the electromechanical vs electrostatic contributions to PFM imaging contrast in the ambient and high vacuum environment, observing domain arrangements, investigating their evolution under the inhomogeneous tip fields and local domain switching behavior in ferroelectric thin film, lead-free piezoelectric ceramics and relaxor-type single crystals. Finally, the last two chapters cover the fabrication processing and the prospect of functional ceramics. Chapter 5 describes science and technology of fabrication processing for different kinds of functional ceramics including capacitor ceramics, ferrite ceramics, corundum ceramics, PTC ceramics, and superconductor ceramics,etc. Chapter 6 summarizes the future development of functional ceramics, starting with an overview of the ceramic evolution and emphasis of ceramic processing. This book is intended for advanced undergraduate and postgraduate students in the field of materials science and microscopy techniques. Scientists and industrial engineers working on the functional ceramics, microelectronics, optroelectronics, and sensors may also find this book useful.

4IBOHIBJ +BOVBSZ

Qingrui Yin, Binghe Zhu, Huarong Zeng

Ď Preface

Acknowledgements

We would like to thank our colleagues, Prof. Ji Zhou, Prof. Xiangming Chen, Prof. Haosu Luo, Prof. Yongxiang Li, and Prof. Guorong Li who contributed to this book. Also, we express gratitude to Yunshu Liu, Fukang Fan, Yingquan Zhu, Guangrui Ye, Taijun Deng, and Fengye Tao for their kind assistance. The authors are particularly indebted to Dr. Rubin Ye, Dr. Junkun Ma, Dr. Shiqun Xiao, Dr. Xiaohua Deng, Dr. Deji Fu, Dr. CV. Kannan and Dr. V. Vaithianathan for their appreciable support in proof-reading and revising the chapters. The authors greatly acknowledge National Science Foundation of China, Major State Basic Research Development Program of China, National High Technology Research and Development Program of China, and Shanghai Institute of Ceramics, Chinese Academy of Sciences for sponsoring various projects.

Contents

1

Microstructure and Properties of Functional Ceramics .............................1 1. 1 General Description ...............................................................................1 1. 2 Grain .....................................................................................................5 1. 2. 1 Grain category................................................................................5 1. 2. 2 Grain properties..............................................................................9 1. 3 Grain Boundary Structures ...................................................................13 1. 3. 1 Concepts of grain boundary structures ..........................................13 1. 3. 2 Properties of grain boundary structures .........................................14 1. 3. 3 Nano grain boundary structures ....................................................15 1. 4 Pore Phases..........................................................................................16 1. 5 Domain Structure.................................................................................18 1. 6 Mechanical Properties of Ferroelectric Ceramics ..................................27 1. 6. 1 General ........................................................................................27 1. 6. 2 Electric domain and internal stress................................................28 1. 6. 3 PLZT ceramics and internal stress.................................................33 1. 6. 4 PTC ceramics and internal stress...................................................39 1. 6. 5 Aging...........................................................................................40 1. 7 Capacitor Ceramics..............................................................................41 1. 7. 1 Ordinary dielectric materials for capacitor.....................................41 1. 7. 2 Relaxor ferroelectric materials ......................................................47 1. 7. 3 Microwave dielectric materials .....................................................48 1. 8 Piezoelectric Ceramics .........................................................................49 1. 8. 1 Microstructures of piezoelectric ceramics......................................50 1. 8. 2 Properties of piezoelectric ceramics ..............................................50 1. 9 Transparent Ferroelectric Ceramics ......................................................53 1. 9. 1 Microstructures of transparent ferroelectric ceramics.....................53 1. 9. 2 Experimental method and two phases of PLZT ceramics...............55 1. 9. 3 Domain switching properties of PLZT ceramics............................57 1. 9. 4 Grain boundaries in PLZT ceramics..............................................67 1. 9. 5 Summary......................................................................................77 1. 10 Thermistor Materials..........................................................................77

Ē

Contents

1. 10. 1 Microstructures and properties of PTC materials .........................78 1. 10. 2 NTC materials and segregation at grain boundaries .....................82 1. 11 Varistor Materials...............................................................................86 1. 12 Ceramics for Humidity Sensitive Resistor...........................................91 1. 13 Magnetic Ceramics ............................................................................92 1. 14 Biologically Functional Ceramics.......................................................94 1. 15 Functional Ceramic Films ..................................................................98 1. 16 Alumina Ceramics ...........................................................................104 1. 17 Summary .........................................................................................105 References ..................................................................................................106 2 Grain Boundary Phenomena of Functional Ceramics............................ 112 2. 1 Introduction ....................................................................................... 112 2. 2 Generalization of Grain Boundary...................................................... 115 2. 2. 1 Grain boundary structure ............................................................ 116 2. 2. 2 Grain boundary properties .......................................................... 118 2. 3 Grain Boundary Segregation .............................................................. 119 2. 3. 1 Generalization ............................................................................ 119 2. 3. 2 Boundary layer capacitors...........................................................122 2. 3. 3 PTC materials.............................................................................124 2. 3. 4 Magnetic ceramics......................................................................127 2. 3. 5 ZnO varistor materials................................................................128 2. 3. 6 Other examples of segregation....................................................129 2. 4 Grain Boundary Region .....................................................................134 2. 4. 1 General description about grain boundary region.........................134 2. 4. 2 Grain boundary region of BaTiO3 ceramics.................................134 2. 4 .3 Grain boundary region of PLZT ceramics ...................................135 2. 4. 4 Grain boundary region and stress ................................................139 2. 4. 5 “Core-shell” structure.................................................................141 2. 5 Grain Boundary Migration .................................................................143 2. 5. 1 Generalization ............................................................................143 2. 5. 2 Centripetal and acentric grain boundary migration ......................144 2. 5. 3 Liquid phase and abnormal grain growth during sintering ...........152 2. 6 Relation between Grain Boundary and Properties ...............................154 2. 6. 1 Influence on mechanical properties.............................................155 2. 6. 2 Influence on electric properties ...................................................162 2. 7 Summary ...........................................................................................166 References ..................................................................................................168 3 Near-field Acoustic Microscopy of Functional Ceramics........................176 3. 1

Introduction .......................................................................................176

ē

Contents Contents

3. 2 History and Development of Scanning Electron Acoustic Microscopy........................................................................................177 3. 3 Physical Principle of SEAM Imaging .................................................178 3. 4 Scanning Electron Acoustic Microscopy Image Processing System .....180 3. 5 Theory Studies of Electron-acoustic Imaging......................................182 3. 6 SEAM Imaging of Ferroic and Other Materials ..................................184 3. 6. 1 SEAM imaging features of ferroelectric domains ........................184 3. 6. 2 Electron-acoustic imaging of ferroelectric materials ....................185 3. 6. 3 Ferroelectric Bi4Ti3O12 single crystal ..........................................190 3. 6. 4 Ferroelasitc NdP5O6 single crystal ..............................................190 3. 7 Magnetic Domains in Austenitic Steel ................................................191 3. 8 Modulation Frequency Dependence of SEAM Imaging Domain Structures...........................................................................................193 3. 9 Electric Field Dependence of SEAM Imaging Domains......................195 3. 10 Temperature Dependence of Ferroelastic Domains in PMN- PT Single Crystals.................................................................................196 3. 11 SEAM imaging of Other Materials ...................................................200 3. 11. 1 Residual stress distribution in Ti3N4 coatings.............................200 3. 11. 2 Stress distribution in ferroelectric composites............................202 3. 11. 3 Stress distribution in Si3N4 and ZrSiO4 ceramics .......................203 3. 11. 4 Stress distribution of Al metal...................................................205 3. 11. 5 Surface structures and internal defects in lead-free piezoelectric ceramics ..............................................................206 3. 11. 6 Phase transitions in superconductor ceramics ............................208 3. 11. 7 SEAM imaging of MEMS devices ............................................209 3. 12 Scanning Probe Acoustic Microscopy...............................................209 3. 12. 1 Tip-vibration mode scanning probe acoustic microscope ...........210 3. 12. 2 Sample-vibration mode scanning probe acoustic microscopy..........214 3. 13 Comparisons of SEAM with SPAM..................................................225 References ..................................................................................................225 4 Piezoresponse Force Microscopy of Functional Ceramics......................229 4. 1 Introduction .......................................................................................229 4. 2 History and Development of Scanning Probe Microcopy ....................230 4. 3 Piezoresponse Force Microscopy........................................................231 4. 3. 1 Operation principle.....................................................................231 4. 3. 2 PFM imaging features ................................................................234 4. 4 PFM Imaging of Ferroelectric Domains..............................................236 4. 4. 1 Ferroelectric thin films ...............................................................236 4. 4. 2 Ferroelectric ceramics.................................................................241 4. 4. 3 Ferroelectric single crystals ........................................................247

Ē Contents Contents Ĕ 4. 5 Dynamic Behavior of Nanoscale Domain Structure ............................261 4. 5. 1 Domain writing ..........................................................................261 4. 5. 2 Domain nucleation and reversal ..................................................262 4. 6 PFM and SPAM Characterization of Ferroelectric Materials ...............273 4. 6. 1 Bi4Ti3O12 lead-free ceramics .......................................................273 4. 6. 2 PMN-PT single crystal ...............................................................275 4. 7 Summary ...........................................................................................279 References ..................................................................................................279 5 Fabrication Processes for Functional Ceramics......................................283 5. 1 Introduction .......................................................................................283 5. 1. 1 Capacitor ceramics .....................................................................287 5. 1. 2 Ferrite ceramics..........................................................................288 5. 1. 3 Corundum ceramics....................................................................289 5. 1. 4 Piezoelectric ceramics ................................................................290 5. 1. 5 PTC ceramics.............................................................................290 5. 1. 6 Varistor ceramics........................................................................291 5. 1. 7 Superconductor ceramics............................................................291 5. 2 Raw Material and Powder Preparation................................................292 5. 2. 1 Ball mill mixing and grinding .....................................................293 5. 2. 2 Powder preparation by oxide methods.........................................294 5. 2. 3 Powder preparation by co-precipitation.......................................298 5. 2. 4 Powder preparation by sol-gel method ........................................299 5. 2. 5 Powder preparation by hydrothermal method ..............................300 5. 2. 6 Powder preparation by spray pyrolysis........................................301 5. 3 Shaping and Forming of Functional Ceramics ....................................301 5. 3. 1 Processing of thin films ..............................................................302 5. 3. 2 Processing of thick films.............................................................305 5. 3. 3 Dry pressing...............................................................................307 5. 3. 4 Iso-static pressing....................................................................... 311 5. 3. 5 Hot injection moulding...............................................................312 5. 3. 6 Slip casting.................................................................................313 5. 4 Sintering............................................................................................314 5. 4. 1 Sintering mechanisms.................................................................314 5. 4. 2 Sintering process ........................................................................317 5. 4. 3 Grain growth ..............................................................................321 5. 4. 4 Abnormal grain growth...............................................................322 5. 4. 5 The effects of pressure and atmosphere on sintering....................323 5. 4. 6 Pressure sintering .......................................................................324 5. 4. 7 Micro-porosity sintering .............................................................325 5. 4. 8 Microwave sintering...................................................................326

Contents

5. 5 Mechanical Finishing.........................................................................327 5. 6 Electroding ........................................................................................329 5. 6. 1 Electroding from silver paste ......................................................330 5. 6. 2 Electroding from nickel plating...................................................332 5. 6. 3 Other electroding methods..........................................................334 References ..................................................................................................335 6 Review and Prospect of Functional Ceramics.........................................337 6. 1 Evolution of Ceramics........................................................................337 6. 2 Development of Functional Ceramics and Relation with Other Factors......................................................................................338 6. 3 Importance and Complexity of Understanding Functional Ceramic Effects and Mechanism ........................................................342 6. 4 Emphasis of Ceramic Processing........................................................344 6. 5 Future Development of Functional Ceramics ......................................345 6. 5. 1 Dielectric ceramics and devices ..................................................346 6. 5. 2 Chip type ceramic devices ..........................................................347 6. 5. 3 High performance, high temperature piezoelectric ceramics ........348 6. 5. 4 Lead-free piezoelectric ceramics.................................................349 6. 5. 5 Thermoelectric ceramics.............................................................350 6. 5. 6 Functional ceramic films ............................................................352 6. 5. 7 Functional crystals......................................................................356 6. 5. 8 Battery materials ........................................................................358 6. 5. 9 High temperature superconductive ceramics................................360 6. 5. 10 Fabrication of ceramic micro-components.................................360 References ..................................................................................................362 Index..............................................................................................................364 Appendix .......................................................................................................366

Microstructure and Properties of Functional Ceramics ჷ

This chapter mainly deals with the microstructure and property of functional ceramics. Following the description of grain, grain boundary, nano-grain boundary, pores, domain structure, and mechanical behavior, the emphasis is placed on the relationships among chemical, and phase composition, microstructure, and properties of various functional ceramics including capacitive ceramics, piezoelectric ceramics, transparent ceramics, thermistor ceramics, magnetic ceramics, bioceramics, thin film ceramics and corundum ceramics. About 90 photographs of representative microstructures are presented, which are useful for comparison and examination of the related structure characteristics, and are also taken as the criterion to judge the constitutive phase or defects and to evaluate the ceramic properties.

1.1

General Description

Modern advanced ceramics, have become key materials in the development of modern technology although ceramics are ancient materials. According to different applications, advanced ceramics could be roughly divided into structural ceramics and functional ceramics while functional ceramics account for a large ჷ

Pictures in Chapter 1 and 2 without specific reference are works of former fourth Lab in Shanghai Institute of Ceramics CAS. The authors are grateful to Zujun Gu, Ruifu Huang, Xiangyun Song, Jing Sun, Yujun Zhang for their great assistance in the composition of the book. The authors appreciate great contribution from Haikuan Ao, Yao Yao, Rongming Sun, Zhili Chen, Weiping Yin, Xinsen Zheng and other colleagues from former 3rd group of fourth Lab. Thanks also go to Zhiwen Yin, Pingchu Wang and Xiangting Li for the enlightenment from discussion with them.

1 Microstructure and Properties of Functional Ceramics

percentage of advanced ceramics. Generally, polycrystalline inorganic materials possessing electric, magnetic, elastic, biologic, superconductive, or other chemical functions are called functional ceramics while those ceramics possessing mechanical, thermal or some chemical functions are called structural ceramics. At present, the production value of functional ceramics and structural ceramics are 3 to 1. The annual production value of functional ceramics around the world has reached 7 billion USD, with 21% of capacitor, 18% of magnetic ceramics, 15%~16% of integrated package, 11% of piezoelectric ceramics, 5.6% of thermistor, 5.1 % of transducer, 2.4% of substrate, and 1.9% of varistor. All these functional ceramics have been widely used in the fields of computers, tele-communication, television, home appliances, space technology, automation, automobile and medical care, etc. Ceramic microstructures include various structural images obtained from different kinds of microscopes (Fig.1.1).

Single crystal

Ceramics

ü ü

ü ü ü

Polyphase

(poly-crystalline) Q quartz grain GB grain boundary M mullite G grain GP glass phase

ü ü

Pore (empty space) G grain P pore

ücrack

Crack

C

(a)

Fig.1.1

Microstructures of ceramics

(a)Illustrative diagram for the microstructure of ceramics; (b) Microstructure of PTC BaTiO3; (c) Microstructure of PLZT ceramics; (d) Quartzitic sandstone a natural hot pressed ceramics

ü

1.1

General Description

According to Petzow, all phase regions and flaws contained in structures would be reflected in microstructures, which determine many properties of materials. According to Pask(1984), microstructures should include sizes and distribution of grains and pores, phase composition and distribution, nature of grain boundary and its defects and flaws, composition homogeneity as well as domain structures. Ceramics are materials derived from powdery raw materials through various processing, and possess specific microstructures and properties. Thus microstructures comprehensively reflect previous processing, and bring specific properties to materials. Microstructural analysis is also important for determining phase diagrams, providing bases for property analysis, instructing modification on formulation, processing improvement, production rationalization, and failure analysis. The following are several examples which further explain the importance of microstructure analysis. Example 1: There was a newly built transformer substation in Shanghai. In a very hot summer the elevated temperature caused a dramatical rise of the oil pressure with a ceramic container, and gave a blast on it. Luckily, it happened during the trial run, otherwise it would probably have caused life threat and power shut down for a massive area. The microstructural analysis afterwards on that ceramic debris showed that the silica particle had sharp boundaries in the high-tension insulator ceramics, which provided evidences that silica particles did not fully melt and react with feldspar and other glass frits during sintering while the boundaries of silica particle of normal insulating ceramics are corroded with glass phases. The microstructure demonstrated that the ceramic body had not been fully sintered, thus it had low tensile strength and couldn’t survival under high oil pressure. Example 2: At a PTC heater manufacturer in Cixi city of Zhejiang province, the ceramic pieces were not broken after voltage test, but cracked in a large amount after packing and transportation, which caused a loss of hundreds of thousands of ceramic pieces (0.65 Yuan/piece at that time). In the analysis of microstructure of PTC ceramics, abnormally grown grains of large sizes were found. During the puncture testing, large grain would expand or contract along axis, which produced large residual stress and micro cracks. Thus the as sintered ceramic plates had normal strength, but became fragile and brittle after puncture testing because of micro cracks. After discovering the cause of problem, some additional additives were introduced to the composition to restrain the abnormal grain growth , and the problem was solved (Zhu, Yao, Zhao, et al, 2001). Example 3: Ferroelectric thin films have wide applications in the field of nonvolatile ferroelectric memory and micro electromechanical system. The unit dimension of high-density ferroelectric memory has been reduced to 30~100 nm, thus physical properties, including formation of ferroelectric domain, polarization reversion mechanism, polarization fatigue, retaining, degradation and aging of


polarization, need to be investigated in sub-micron scale. Scanning Probe Microscopy based on Atomic Force Microscopes could satisfy the requirements with resolution of 30 nm, and can be used to exhibit details of polarization reversion, which provides an observation foundation for improvement of anti-degradation. In addition, these kinds of microscopes are very useful for investigation on microstructures of all other nano ceramics too. Properties of materials can be divided into two categories. One kind is intrinsic or inherent properties, which mainly depend on internal characteristics of compounds and crystal structure and also depend on attributes of ferroelectrics, ferromagnetism, semi-conductivity and super-conductivity. the other is non-intrinsic properties, which often have relationship with microstructures. It is ceramic researchers’ task to investigate relationship among processing, structures and properties, in which microstructures play a significant role. During many global symposiums in last thirty years, it has been widely accepted that microstructure plays a significant role in material sciences with a special emphasis on interface and grain boundaries. When grain dimensions decrease to a nano level, proportion of grain boundaries abruptly increases, and materials turn into nano ceramics. Control of microstructures is an important approach to obtain materials with desirable properties. In early stage, optical microscopes were main tools for observing microstructures, while observation instruments currently including optical microscopy, Transmission Electron Microscopy (TEM), Scanning Electron Microscopy (SEM), Scanning Electron Acoustic Microscopy (SEAM), Scanning Nonlinear Dielectrics Microscopy (SNDM), Atomic Force Microscopy (SFM), Infra-red or Laser Microscopy, are now widely used, with resolution from 1μm to 30nm and even to 0.5 nm, so as to enable comprehensive and precise investigation on material properties, such as optics, electronics, mechanics and thermology etc., and to widen applications of material functions. At present, it has been evolved from simply understanding materials’ microstructures to purposely controlling and preparing materials with specific microstructures. Many properties of ceramics conform to the combination principle. For example, mass per unit of volume decreased as porosity increases, which decreases permittivity. Density or thermal capacity is generally the proportional combination of all phases in materials, and thermal conductivity conforms to similar rules only with some exceptions. However, some other properties of ceramics do not conform to the combination principle. For example, small quantity of second-phase additive can dramatically change electric conductivity of materials. Some properties of materials, especially those related with functional processes and energy transformation can be affected by interaction between different grains and phases, and also fall into the second category. Many attributes or phenomena belong to interactive properties, such as alignment of electric

1.2

Grain

domains in ferroelectric ceramic under electric field, the function of high-energy grain boundary zone during the process, and the effect of space charge and induced field in grain boundaries. Key factors of functional ceramic microstructures will be discussed as below.

1.2

Grain

Grain category and grain property are very important to understand functional ceramic microstructures and physical property, here we introduce their concepts in details.

1. 2. 1 Grain category Grains (or main crystal phase) are main body in microstructures, with three kinds of textures, i.e. idiomorphic, semi-idiomorphic and xenomorphic granular texture (as shown in Fig.1.2) (Shi, 1982).


Fig.1.2

Crystalline morphology of ceramics (Shi, 1982)

(a) Idiomorphic texture (mullite porcelain); (b) Semi-idiomorphic texture (95% alumina porcelain); (c) Xenomorphic granular texture (smelted quartz crystals in porcelain)

Grains grown in a sound sintering condition will exhibit the characteristic shape of crystal structure itself, which forms idiomorphic textures. Grains grown in an unsound or restraining condition will just partially exhibit its characteristic shape, which forms semi-idiomorphic textures. In addition, grains of anomalous shape with various sizes forming xenomorphic granular texture are quite common, especially in functional ceramics. According to different requirement on functions, different main crystal phase will be selected. For example, materials with large permittivity ε, (such as BaTiO3 and rutile) will be selected as main crystal phase. Materials with low ionic conductivity or covalent bond (such as α-Al2O3 or BeO) will be choosed. ferroelectric materials without symmetric center in crystal structure are excellent electro-mechanical materials due to its high spontaneous polarization Ps. Phase transformation induced under fields (electric or mechanic) will introduce changes of optical properties in some materials, which can be used as electric-optical functional (optical switch) or storage materials. Dimensions of some materials can be repeatedly and steadily changed under applied field, which can be used for electrostriction or actuator. Some transition metals (rare earth elements) retain their intrinsic spin angular distance after bonded into compound with magnetism (such as Mn-Zn ferrite), which can be used as magnetic materials. The properties of grains or main crystal phases have greatly determined properties of materials. For example, during the development of high voltage porcelain in past fifty years, the main crystal phases have changed from mullite, silica to corundum, with an increase of strength, from 30~70MPa (hard porcelain), 95~125MPa (silica porcelain) to 175~200MPa (alumina porcelain). Besides main crystal, grain sizes can also affect properties. Since grain have been formed from original powder particles after the processes of diffusion, gas elimination and grain boundary migration during sintering, grain sizes

1.2

Grain

depend on powdery raw materials, composition, additive phase and sintering process. The perfectness and idiomorphism of grain growth, as well as the interlocking grade would affect functional properties of materials. Generally, homogeneity of grains is preferred for microstructures, so dispersion or agglomeration, narrow or wide size distribution of powdery raw materials, appropriate second phase would affect microstructures. Narrow size distribution raw materials, appropriate second phase and forming processing are in favor of homogenous microstructures. In some cases, among fine grains emerge few coarse grains (abnormal grain growth), which possess greatly different thermal expansion or contraction along crystal axis and anisotropy. Stress concentration is generally formed along large coarse grain boundaries, which are electrically and mechanically weak and where micro cracks originate (as shown in Fig.1.3) (Burke, 1963).


Fig.1.3 Microcracks appearing in the grain and grain boundary of ceramic material

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(a) Cracks in the grain of PTCR microstructure TC=240 ; (b) Cracks in grain of KNaNb2O3 ceramic; (c) Large grain in high voltage porcelain (BaTiO3); (d) Cracks in grains of high tension porcelain insulator

Rationality of sintering processes can be judged from grain shape and crystalline morphology (such as over-firing and under-firing as shown in Fig.1.4). For materials with domain structure, grain size smaller than a critical dimension will affect development of electric or magnetic domain, which will finally affect its dielectric or ferroelectric properties.

1.2

Fig.1.4

Grain

Crystalline morphology of the quartz grains in porcelain

(a) Quartz grains with some smelted edges in porcelain (Chu, Weng, Wang, 1994); (b) Quartz grains with more smelted edges in porcelain

1. 2. 2 Grain properties Grains possess similar properties to single crystals of the same composition. Ideally all micro zones in a grain possess the same properties, however differentiation in composition and properties could be found between boundary zone (shell) near boundaries and central zone (core) (Rawal, 1981). Two kinds of crystal forms could be even found in one grain (for example, cubic near the boundary and tetragonal at the core, as shown in Fig.1.5).


Fig.1.5

Core and shell structure in PLZT ceramics (optical, crossed Nicol)

Micro zones with little differences in optic axial angle have been found in one grain of ferrite ceramics. Ununiformity in nano-scale has been observed within grains of solid solution ceramic of PbInNbO3-PbMgNbO3 due to defects of processing (Babakishi, Tai, Choy, 2002). Since grains are formed through diffusion and boundary migration (growth) during sintering, uniformity within grains can not be obtained if the process is out of equilibrium, so trace and residua of mixture of raw materials, boundary migration and diffusion will remain in non-uniform grains, which is quite common in many solid-solution functional ceramics such as PZT, BaTiO3, and ferrite. According to Laurent et al (2001), when preparing La-PZT from oxides, pores from chemical etching due to composition gradation of Zr, Ti within grains since the diffusion speed of Zr, Ti, and La are different; when preparing La-PZT from precursors, neither pores nor composition gradation was found in grains. In addition, grain sizes could also affect properties. For example, in piezoelectric ceramics, coercive field Ec increases with decrease of grain sizes. Besides, the diffuseness of phase-transition is more significant in finer grain materials. Tartaj (2001) has proved that the tetragonality c/a of PbTiO3 ceramics decreases with the decrease of grain sizes, and obviously the internal stress status will be different. As for microstructures of ferroelectric ceramics, the relationship between grain sizes and properties is quite complicated. BaTiO3 will be taken as an example to further discuss the question (as indicated in section 1.15, Huarong Zeng has obtained the critical dimension for ferroelectric thin films, which indicated that ferroelectric vanishes with grain size under the critical dimension). Ceramics capacitor is one of the electronic devices with large amount of market volume. BaTiO3 capacitor widely used dielectric coefficient of normally sintered BaTiO3 ceramics is around 1200~1500, while that of fine-grained BaTiO3 ceramics

ˈ

1.2

Grain

could reach 4000~5000, or even 6000. However in single crystalline BaTiO3, ε along c axis is 4000, and ε along a axis is only 170, i.e. the average ε is around 950~1200, on which there is still no convinced explanation after around 50 years of researches. At present, there are two explanations: W R Buessem et al (1966), L Metoseriu et al (1999) and other scientists think that high value of ε is caused by internal stress; while G Arlt, D Hennings (1985) and others thinks that high ε is induced by a large number of domain walls in fine grain materials. According to the first theory, when cooling down from high temperature to TC, phase transition from cubic to tetragonality takes place in BaTiO3 single crystal, and the deformation caused by phase transition is allowed since its surfaces are free of restriction (for example, surfaces perpendicular to c axis move outwards and surfaces parallel to c axis move inwards). However in BaTiO3 ceramics, surfaces of grains could not move freely since they are surrounded by other grains so as to form internal stress under TC. Many 90 domains (twining) would be induced to reduce internal stress in coarse-grain ceramics (5~50) μm, but very few 90 domains are seen in fine-grain ceramics (around 1μm), which brings large internal stress to fine-grain ceramics. ε could reach to 3000 with internal stress of 64 MPa, and even 6000 with internal stress of 80 MPa. Thus concerning the decrease of tetragonality with decrease of grain sizes, high value of ε in fine-grain BaTiO3 ceramics could be explained by the combination of large internal stress under TC and lack of 90 domains. According to the second theory, few 90 domains in fine-grain BaTiO3 ceramics as mentioned in the first theory is actually caused by defects of the chemical etching method. G Arlt indicated that 90 domains are visible after improving corrosion preparation. He also indicated that the width of ferroelectric domain d decreases with decrease of grain sizes α, following the relation of α d1/2 (in the grain size range of 1~10μm) (Arlt, Hennings, 1985). G Arlt (1985) has proved that a large number of domain walls would make more contribution to dielectric coefficient of ε, which may reach 5000 with average grain diameter(a) of 0.8~1.0μm below temperature TC,When α decrease further (for example, less than 0.7μm), ε would dramatically decrease, and the crystal form at room temperature would change from tetragonal to trigonal or orthorhombic system (Arlt, Hennings, 1985). G Arlt et al do not eliminate the contribution from internal stress for the increase of ε and attribute it to the combination of the two facts. But which fact is more significant needs further study. L Metoseriu et al (1999) have shown that transition temperature and tetragonality decreased when the grain size decreases, but the internal stress will be increase at this time. It is different from BaTiO3 ceramics, K Okazaki et al (1973) has proved that TC

e

e

e

e

e

ĝ


increases with decrease of grain sizes in piezoelectric and ferroelectric PZT and PLZT ceramics, which means that piezoelectric properties enhance and cohesive field declines with the decrease of grain sizes. Space charge theory could be utilized to explain the phenomena. It is reasonable to think that with the decrease of grain sizes, restriction enhances and it would restrain dielectric and polarized properties. Finer grains also mean more grain boundaries, which would affect transfer of some properties. As for the dependence of grain sizes on phase transition, the increase or decrease of transition temperature depends on expansion or contraction of grain during phase transition. With the decrease of grain sizes, space charge field effect would increased, which would restrict movement of domain walls so as to enhance the stability of resonant frequency in piezoelectric filter materials. PZT materials of fine grains often have higher fracture toughness than that of coarse grains, e.g. 40% higher. Grain sizes in superconductor also affect its toughness, while magnetic materials of fine grains have a higher saturation magnetization, a lower coercive field and lower resonant frequency. Generally optical axis of each grain in ceramics is randomized, so the properties of ceramics are average value of all grains. Nowadays it is practical to align grain along certain direction so as to provide ceramics with oriented properties, like that in single crystals. Following methods could be adopted to obtain this microstructure: 1. Hot forging or hot pressing sintering. 2. Eutectic solidification. 3. Templated grain growth sintering. Hot forging method could be used to produce considerable directional alignment. For example, piezoelectricity in some layered ferroelectric system containing Bi is not quite strong, only with a Kp of 0.2. However, its Kt could be enhanced to 0.35~0.42 by doping MnO, NiO, Cr2O3. Due to advantages including desirable temperature stability of resonance frequency ((0~20) 10 – 4%), low dielectric constant ε (100~200), high Curie Point TC (above 550○C), good resistance against degradation, these materials could be effectively utilized at high frequency. Dielectric constant ε perpendicular to hot forging direction is highly different from that parallel to hot forging direction as shown in Table 1.1.

h

Table 1.1

Dielectric constants of some typical materials by hot forging method

Materials

ε parallel to hot forging direction

ε

perpendicular to hot forging direction

Bi4Ti3O12

270

1170

Na0.5Bi4.5Ti4O15

280

3030

PbBi4Ti4O15

1350

5300

M M Seabaugh (1997) prepared highly oriented corundum ceramics by using templated grain growth sintering, with materials as below:

1.3 Grain Boundary Structures

1. Colloid gibbsite ( 15 mass% ) was selected as alumina precursor. 2. α-Al2O3 fine particles ( a. Fig.1.9 shows domain structure of tetragonal BaTiO3 (macroscopic model).

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Fig.1.8

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The distortion of domain structure in BaTiO3 single crystal (Liu, Xu, 1990) (a) The topography image of cubic BaTiO3 crystal; (b) The 180 twin domain structure in tetragonal BaTiO3 crystal (solid line); (c) The 90 twin domain structure in tetragonal BaTiO3 crystal (solid line)

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Fig.1.9 Domain structure of tetragonal BaTiO3 (macroscopic model) (M. D. Liu, Y. C. Xu, 1990)

Electric domains with polarization axis parallel to crystal surface are called a-domains; while domains with polarization axis perpendicular to crystal surface are called c-domains. The difference of refractive indices along a axis and c axis at room temperature is: na nc=0.055, thus domains could be observed under polarized light due to its characteristics of birefringence. With orthogonal polarized light perpendicular to the crystal surface, c domains (as dark area) and a domains (as bright area) could be observed. In BaTiO3 crystal of orthorhombic system, Ps is along 011 direction of formerly cubic system, thus there are also 60 and 120 domains besides 90 and 180 domains; In BaTiO3 crystal of trigonal system, Ps is along direction of formerly cubic system, thus there are also 60 and 109 domains besides 180 domains. Since positive sides of domain corrode faster than negative sides during chemical etching, domains could be observed from etched surface with “up and down” like morphology. Recently developed scanning probe acoustic microscope (SPAM) could be used to observe

ˉ

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domain structures without any special treatment. Actually domain structures of ferroelectric crystals or ceramics are quite complicated and randomly oriented, containing 90 , 180 domains and other interactive structures, as shown in Fig.1.10 (a), (b), (c), (d). An external field needs to be applied to compel randomly oriented domains to realign coincidently to the polarized direction, and to expand domain volume to the same direction as the applied field, so as to produce an overall Ps. Fig.1.11 (a) shows the schematic poling process in piezoelectric ceramics PZT, and (b) shows domain orientation after poling and chemical etching.

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1.5 Domain Structure


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Fig.1.10

Domain and domain configuration

(a) 90 and 180 domains in piezoelectric ceramics Z-4; (b) Domain configuration (PLZT, chemical etching, Phase contrast microscopy, direction of electric field E: shown as arrow); (c) Ribbed domain, replica, TEM; (d) Intersection of domains (PLZT, 8/67/33, after chemical etching, TEM)

Fig.1.12 shows domain configuration under different poling conditions. If Ps inverses for 180 , no additional stress would be introduced; if non-180 domain wall motion occurs, deformation would be induced (for example, elongation along poling direction) to produce mechanical stress. Ceramic plates may fracture if the internal stress is too high. For poled ceramics, internal stress would be gradually released in deposition for a period of time, which is called aging. T Ogawa (2002) has estimated switching and reorientation of electric domains by measuring ferroelectricity and piezoelectricity under applied field and their relationship. For example, switching of 180 domains could be evaluated from

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the poling fields for minimum Kp and K33; switching of 90 , 71 , 109 domains could be evaluated from the poling fields for εmax or minimum value of frequency constant.

Fig.1.11 Poling process in piezoelectric ceramics (sketch) (Xu, et al, 1978) (a) and domain orientation after poling piezoelectric ceramics (PZT, chemical etching) (b)

Electric domains often nucleate at locations where defects (such as grain boundaries) or stress concentrate. Nucleation of 90 domains would eliminate substantial internal stress. The possibilities of nucleation at grain boundaries are as follows: intersection points of four gains > intersection lines of three grains > intersection plane of two grains. Several twin-crystals often share a point along grain boundaries, i.e. nucleation of domain may cross grain boundaries: when a grain turns into tetragonal system from cubic system, domains show band structures and stress would occur along grain boundaries, which makes the domain movement in neighboring grains move easier along specific direction. During phase transition, the front side of band domains tends to cross grain boundaries to release stress, so as to minimize system energy, as shown in following section (Fig.1.43 (c)). Band width of 90 domains may vary with strength of internal stress, i.e. larger band width under smaller stress and vice versa. ε is often larger with more domain walls, and domain width is proportional to grain size:

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Fig.1.12

Domain configuration under different poling conditions

(Li0.11Na0.89NbO3,Ec=2400V/mm, crossed nicols) (a) 1900 V/mm, < Ec, 120 /20min; (b) 2500V/mm, around Ec, 120 (c) 7000 V/mm, >Ec, 120 / 20min

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ĝ

ć

ć/20min;

90 domain width (grain size)1/2. 90 domains play a very important role in adjusting and relaxing stress. There are a large amount of 90 domains within grains of ferroelectric ceramics. However, very few domain walls would be left if a grain is separated out since these domain walls are created from twining to release stress produced during phase transition. When a grain is separated from other grains, volume change from

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phase transition could be easily accommodated by free surface movement, so there’s no need for domains to adjust strain. Poly-crystals always have a high stress status relative to single crystals, thus there’s always a driving force to reduce residual stress, which causes aging. Aging of properties is often related to following processes: 1. Change of domain configuration with time. 2. Trend of domain configuration towards equilibrium. 3. Migration of impurities and vacancy towards domain walls or grain boundaries. 4. Regularization of impurities and vacancy along poling axis etc. For example, back rotation of 90 domains increases ε, while back rotation of 180 domains decreases ε. The internal friction of ferroelectric ceramics also lies on domain configuration and the interaction between domain walls and lattice defects. N G Zhang et al (2001) have investigated poling fatigue of PLZT (2/70/30) ceramics, and discovered that after certain number of switching cycles fatigue occurred, and piezoelectricity degraded with a decreased Pr. In addition, fatigue occurred earlier at lower frequency. For example: Degradation of ferroelectricity occurred after 103 cycles at 10Hz; Degradation of ferroelectricity occurred after 104 cycles at 50Hz; No degradation observed after 1010 cycles at 100Hz; Little degradation observed after 1010 cycles at 500Hz; M L Eng (1999) has investigated domain switching properties of nano BaTiO3 ceramics. It has been concluded that in a specific field, domain switching would occur only when critical switching time τc is satisfied, i.e. if frequency of applied field is higher than 1/τc, electric domains could not be switched and no aging or fatigue could occur even under applied field. It is commonly accepted that fatigue originate from (1) micro cracks caused by mechanical stress, and (2) migration of defects under electric field. A Levstik et al (1997) has performed test of cycled electric load (15 kV/cm, 50Hz) on PLZT (8/65/35) ceramic and measured its piezoelectric coefficient d33 and quality factor Q. It has been discovered that d33 started to decrease after 4 105 cycles, Q started to decrease after 106 cycles, and micro cracks occurred after 3 106 cycles. The earlier degradation of d33 may be caused by domain switching. Cross research group from Pennsylvania University of the US has conducted systematic research on electric fatigue of ferroelectric ceramics(Jiang, Cross, 1993). If materials degrade dramatically under cycled loading of alternative field, it would limit their applications in high-strain actuators and non-volatile memory devices. Fatigue is related to many factors including surface status, preparation of electrodes applied, microstructures, composition and working temperatures etc., in which pores in microstructures play an important role.

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The number of switching cycles to produce 70% decrease of remnant polarization has been taken as the criteria to estimate electric fatigue of materials with different density. It is well known that there’s space charge on surfaces of poled ceramic to neutralize polarization, and pore sizes in ceramics could reach 1μm or even dozens of micron. When high voltage field applied, space charge could enter grains, grain boundaries to interact with domain walls, which would stabilize domain configuration and make domain switching more difficult under applied fields, so as to decline the polarization. Under applied field, domain may extend from one surface to the opposite side in as-received BaTiO3 single crystal. However after several cycles of switching, domain would be confined within crystal body instead of extending from one side to the other. Space charge often accumulates around pores and thick grain boundaries which behave like pins on domain, and cause electric fatigue. In addition, the ceramic layer next to electrodes would undergo some electrochemical reactions under applied field and change its color obviously, which is un-restorable; while above mentioned fatigue due to domain pining could be recovered to as-received status by heating up to phase transition temperature Tεmax. There’s also space charge within materials to neutralize the polarization: positive and negative space charge are accumulated on positive and negative sides of domain respectively and establish a space charge field. Space charge also accumulates at grain boundaries and could affect domain motion. Analysis on domain configuration of ferroelectric ceramics is an important method to understand relationship between poling condition and domain orientation, aging property and switching behavior. A large amount of electric domains and domain walls are produced to minimize internal stress during phase transition, so as to reduce strain energy and overall system energy, same as magnetic domains. Recently, strong piezoelectricity has been found in a single crystal of Pb(Zn1/3Nb2/3Ti0.09)O3 solid solution: K31=80%, K33=95%, d33=2500pC/N (K31=40%, K33=70% for ordinary soft PZT ceramics). To explain the strong piezoelectricity, Ogawa et al ( 2002) has conducted research. They found that under proper poling conditions, a single domain could be produced cross over a single crystal, which brings along strong piezoelectricity. However, if the poling voltage is unfavorably increased, multiple domains will be formed and piezoelectric property will be degraded. Thus proper poling conditions are crucial to produce ideal domain configuration, so as to enhance piezoelectric properties. However, there are also other theoretical explanations to the phenomenon of strong piezoelectricity in these materials.

1.6

1.6

Mechanical Properties of Ferroelectric Ceramics


Ferroelectric ceramics are operated at high electric fields and at mechanical loads to use their full electromechanical potential.The intrinsic coupling of electrical and mechanical effects leads to the development of fracture mechanical experiments and concepts in ferroelectric ceramics.

1. 6. 1 General Electric properties of ferroelectric ceramics have been comprehensively studied while only a few researches have been focused on their mechanical properties and internal stress which play a very important role on materials functions. Here is an example to demonstrate the effects of internal stresses on working lifetime of ceramics. When preparing transparent ferroelectric ceramics through hot-press sintering, 99% alumina ceramics were used to make hot-pressed dies which had to sustain pressure of 20~30MPa under 1150 for dozens of hours including multiple heating and cooling. They should also be reusable without damage. Casting process was initially used to produce hot pressed dies with 65mm in diameter and 85mm in height, and the dies were sintered at 1780ć. However the obtained dies consistently fractured after being used only once. Isostatic pressing process was introduced then to produce the dies, and they were still intact after being used over 50 times. The dies produced with isostatic pressing had homogeneous structures and low internal stresses, and retained their high sustaining strength after repeated uses; while dies produced with casing process contained substantial internal stresses which significantly decreased the strength. Thus internal stresses play significant roles on working lifetime of ceramics. Similarly, once refilled with boiling water, vacuum flask made without annealing would fracture immediately. Internal stresses in ceramics play a significant role on material properties, and methods have been developed to measure and control internal stresses. Recently, more interests have been focused on mechanical properties than electric properties for piezoelectric and ferroelectric ceramics (Pferner, 1999), since these materials would be exposed to severe mechanical loading in the applications such as ultrasonic transducers, actuators and piezoelectric transformers. In some newly applications such as of piezoelectric powered pressure heads and fuel injection system of automotive technology, the expectation on mechanical properties for ceramics are even highly. The research on mechanical properties also helps to understand electric properties. For example:

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BaTiO3 ceramics with fine grains are frequently used as capacitors, and their dielectric coefficient is even higher than that of average value of single crystals, which was quite elusive for almost fifty years. However the researches on mechanical properties and internal stresses have uncovered this puzzle. Details are discussed as below. Ceramics or polycrystals are prepared through sintering and cooling. Following factors could bring various deformations to grains during cooling: 1. Differences in expansion and contraction among grains. 2. Anisotropy of non-cubic crystal system. 3. Variety in grain sizes. 4. Composition diversity among grains. 5. Ionic doping. Since grains in poly-crystals are interlocked and restricted by each other, they can not expand or contract freely, which produces substantial internal stresses, with both tensile and compressive stresses existing in different area. At high temperatures, diffusion and deformation are available to partly eliminate or relax internal stresses (diffusion through grain boundary is also a stress relaxation mechanism), after ceramics cool down to room temperature, there have always some residue stresses remaining. Jonker has investigate with X-ray diffraction, the difference of crystal axis ratio c/a between BaTiO3 ceramics (restricted grains) and powders (free grains) with the same grain sizes, and discovered that the c/a in BaTiO3 ceramics was 14% smaller than that of BaTiO3 powders, which showed stress status of grains in ceramics and restriction of grain expansion. Cross has indicated that internal stress within grains could reach up to 60MPa, while P L Janega (1986) has calculated that internal stress in ferroelectric phase could reach 60~90MPa. Y Sakabe (2002) has studied internal stress by analyzing peaks of XRD (X-ray diffraction spectrum), and concluded that peak’s dullness was related to internal stresses. XRD peaks of (113) and (311) of powders, as-sintered ceramics and etched ceramics have been compared. It showed that chemical etching could remove internal stresses. Kiyoshi Okazaki (1992) has listed crystal axis ratio c/a and corresponding internal stresses of some PLZT ceramics with various compositions. When La in composition (x/0/100 )increases, c/a, and TC as well as internal stresses decrease.

1. 6. 2 Electric domain and internal stress There are ferroelectric polarization, electric domains and phase transition in ferroelectric ceramics. Actually electric domains or domain walls result from substantial internal stress caused by phase transition so as to reduce strain energy. This is also true for magnetic domains in ferromagnetic materials. Under applied

1.6


e

field, electric domains could re-orientate or switch. 180 domain switching does not produce strain and does not revert after removal of applied field. However, non-180 domain switching, e.g. 90 domain switching, would produce strain and deformation and would partially revert to its original orientation. For example, based on investigation on PZT ceramics with composition close to phase boundary by measuring strain and lattice parameters, about 53% of 90 domains would realign along applied field after saturation field applied, and about 44% would maintain the orientation after removal of the field (Jaffe, Cook, 1971). However, among these reoriented domains, some may readjust their orientation again over a long time period to reduce internal stress, such as ageing during stocking. Under applied field E, deformation S would be produced in ferroelectric ceramics. The dependence of S on E is similar to dependence of P (polarization) on E, (Berlincourt, 1981). Different materials have different S-E relations with different slope of dS/dE. Large slop of S-E curve means relatively small field could cause large strain, which tend to form micro cracks in the process of polarization under applied field. Usually strain S parallel to applied field is positive, which means materials expand along applied field with tensile stress; while strain S perpendicular to applied field is often negative, which means materials are under compressive stress along this direction. Thus extreme poling conditions, including excessively high voltage, excessively high temperature, and excessively long time, should be avoided to prevent nucleation and slow growth of micro-cracks. It has been found that stresses are concentrated in grain boundaries, so micro cracks tend to nucleate in boundaries and expand into grains (Chung, 1989). Coarse-grain ceramics are more vulnerable to cracks. Kroupa (1989) has found that cracks tend to form in central plane (parallel to electrodes). When samples were cut and polished along central plane, particles would fall off, which means microcracks had been formed to decrease the strength. R A Pferner (1999) found that pores often degrade strength. Long-rod hexagonal PbO crystals of 5~8μm have even found in pores in PZT ceramics, while grain sizes of background materials are 2μm. The size disparity would facilitate crack nucleation to lower strength. Grains of ceramics would contract during cooling process, and materials would be under tensile stress along large contraction, and compressive stress along small contraction. Large temperature change would cause large internal stress σ (Kala, 1983): σ ∝ (α ) ⋅ ( E ) ⋅ (ΔT )

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Where α is lattice expansion, E is elastic modulus and ΔT is temperature change. When internal stress surpasses material strength, cracks will be formed. A typical example is the powdery disintegration of PbTiO3 ceramics. Since the tetragonal distortion is quite large (c/a =1.063) with strong anisotropy, tremendous internal stress would be left in sintered samples after cooling. During deposition,


PbTiO3 ceramics may fracture along grain boundaries and grains would disaggregate into separated particles. If there are only few cracks in ceramics, ceramics will not fracture. When reheated to a certain temperature, no expansion could be observed due to crack healing, and increasing temperature further could bring actual expansion. By using acoustic emission method, it could be found that the temperature for maximum acoustic emission correspond to the temperature for expansion startup (Srikanth, 1992). In fact, to judge occurrence of micro crack by acoustic emission is a very important method. At the cooling section of a tunnel kiln, if the temperature of coming ceramics is too high micro-cracking could be heard, which is just a kind of acoustic emission audible to human ears. Brittleness of ceramics has limited their application, so efforts need to be made to improve their toughness or fracture toughness. Fracture toughness is usually expressed with critical stress intensity factor KIC, which is a constant, related to materials and indicates materials’ ability to resist crack growth. Following mechanisms could be utilized to improve fracture toughness: 1. Toughening by phase transition or micro cracks. 2. Toughening by crack bridging. Qingchun Zhang and R A Pferner (1999) have investigate fracture toughness of piezoelectric and ferroelectric ceramics, and indicated that the minimum value of KIC occurred around phase transition temperature, as shown in Fig.1.13. Strength evolution has similar trend, as shown in Fig.1.14. Qingchun Zhang also proved that a large amount of domain walls would be produced in poled ceramics and could exert retardation on crack growth, so as to improve their mechanical properties. However the improvement also depends on poling direction during polarization and loading direction during mechanical tests. Lishui Zhang and Qingchun Zhang have also found in different crystal system that crack growth is related to lattice distortion. Large lattice distortion brings along large internal stress, which restrains crack growth. Poling also affects crack growth.

Fig.1.13 Temperature dependence of fracture toughness in La-PZT materials

1.6


Fig.1.14

Temperature dependence of flexure strength in La-PZT materials

Except for cubic system, all ferroelectric ceramics are anisotropic with lattice parameters varied according to temperature changes, as shown in Table 1.2. Table 1.2 Crystal system

Variation of lattice parameter for BaTiO3 at different temperature

ć

Lattice parameter

Temperature /

Cubic

130

a1=a2=a3=4.009

Tetragonal

130

a1=a2=4.003, c=4.002

Tetragonal

0

a1=a2=3.992, c=4.035

Orthorhombic

0

Orthorhombic

–90

e Ą a =a =4.013, c=3.976, β=89e51Ą

a1=a2=3.012, c=3.989, β=89 51.6 1

2

According to Table 1.2, lattice parameters of BaTiO3 ceramics may change abruptly or gradually when ceramics are cooled from 130 to –90 . In addition, there are anisotropy and grain sizes distribution in ceramic body. Deformation of large grains along specific crystal axis would be larger, which impose tensile or compressive stresses on grains, i.e. internal stresses. For tetragonal system, internal stress is related to c/a, i.e. when c/a approaches 1, internal stresses caused by anisotropy gets close to zero (Okazaki, 1983). Reference (Ueoka, 1974) shows temperature dependence of specific heat, heat conductivity and linear expansion in BaTiO3 ceramics and reference (Okazaki, 1983) shows temperature dependence of elastic modulus and Poisson’s ratio in BaTiO3 ceramics. Minimal values of both linear expansion and elastic modulus occur at around phase transition temperature. Above TC, α and E increase with increase of temperature, and internal stress is proportional to product of α and E. Thus tremendous internal stress would be produced within PTC ceramics under

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applied field above TC. Besides, tensile strength (only about 1/5~1/10 compressive strength)of ceramics is far lower than compressive strength, ceramics are vulnerable to fracture or cracks under tensile stresses. Internal stress also affects phase transition by enhancing diffusion. R C Pohanka et al (1976) found that flexure strength above TC may differ from that under TC. For example, for hot-pressed BaTiO3 ceramics, strength is 1999.4 MPa at 150 , and 1241.1MPa at 25 . The difference in strength change of above and below Curie point is 758.4 MPa, which is caused by phase transition from cubic phase to tetragonal phase. In another word, internal stress caused by phase transition correspond 758.4MPa. Pferner also found that strength of PZT ceramics is higher above TC and lower around TC, which is caused by internal stress from phase transition. In addition, strength is lower around phase boundaries, where two phases coexist and the heterogeneity lowers its strength. When materials are poling under applied field, KIC parallel to E is larger than KIC perpendicular to E, which is believed to be due to that coherence between grain-boundaries along E is increased while coherence perpendicular to E is lowered ( Pferner, 1999 ). At morphotropic phase boundary or cubic-tetragonal transition temperature, domain walls are more moveable and domain switching only causes little stress, so as to bring large coupling coefficient Kp, large mechanical loss, small mechanical quanlity factor Qm, and low strength, low mechanical and electric hardness ( represented by large S11, compliance and ε ), large specific thermal capacity, and small α. In one word, mechanical, electric and thermal properties of these materials all present their extremum due to Soft-Mode factors. R A Pferner ( 1999 ) has investigated relationship between strain and stress of several PZT materials, and proved that materials with composition adjacent to morphotropic phase boundaries have relatively softer properties, which could be referred in reference (Tian, Lin, Zhu, 1993).

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It is noticeable that materials of different crystal systems have different internal stress, which also affects their properties. For example, in PZT solid solution ceramics, there are tetragonal and rhombohedral systems on two sides of phase boundary. Although the two phases have the same spontaneous polarization Ps, the spontaneous deformation in tetragonal phase is 3 times larger than that of rhombohedral phase, which causes difference in internal stress, so as to affect thermal properties of materials including thermal coefficient of dielectric constant T T T T ε 33 , i.e. TCε 33 , and thermal coefficient of elastic compliance S11 , i.e. TCS11 .

The thermal coefficient of dielectric constant is negative for tetragonal phase and positive for rhombohedral phase ( Anon, 1991 ). Mixed-oxide method is often used to prepare ceramics, and the sintering time is

1.6


not very long, so mass transfer is not complete and there is composition differentiation in different regions, which cause micro zones of stress. It has been found that composition differentiation exists among grains and even one grain may consist of two crystal phases: for example cubic phase at the shell and tetragonal phase at the core, known as shell-core structure. In ferrite grains, it has been found that differentiation of optical axis among some micro zones of 5~10 nm may reach 0.1 ~ 4 , which has been verified by TEM. Normal optical or electronic microscopes take use of optical or electric properties of materials to obtain optic or electric images, while acoustic microscopes take use of mechanical properties (e.g. acoustic speed and elastic modulus) to observe stress and strain through acoustic images. Resolution of recently developed Scanning Probe Acoustic Microscope has reached 30 nm. PLZT ceramics are sensitive to applied field (mechanical or electric) ( Meitzler, 1971), and their optical properties could be altered by micro stress, which could be utilized to observe stress and strain zones at grain boundaries.

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1. 6. 3 PLZT ceramics and internal stress Fig.1.15 shows a “smooth region” around grain boundaries in PLZT and PZT ceramic slices after thermal etching (Zhu, 1990). Thermal etching force mass transferring and migrating from high energy region to low energy region, so originally polished surfaces would present unevenness related to energy, which depends on grain orientation because grain orientation determines array of atoms and atomic planes. Since energy of grain boundaries differs from that of grains, specific morphology would be formed after thermal etching, usually with grooves at grain boundaries and steps and stripes on grain surfaces. Different optic axis orientation of adjacent grains would produce different grain-boundary angle with different energy, and “smooth regions” with different width, which means the properties of these zones differ from those of grains. These strain regions are formed due to grain boundary stress. Fig.1.16 shows distortion region near grain boundary. Stress between two grains could be released through deformation and distortion of boundary region. The PLZT ceramic shown in Fig.1.16 was thermally etched at 1100 , and at this temperature deformation and distortion are allowed to release or relax stress. Fig.1.17 shows bending of domain wall near grain boundary observed by TEM, and Fig.1.18 shows discontinuity of bond contour in grain boundary region of PLZT ceramics by TEM. Both the bending of domain wall in Fig.1.17 and the discontinuity of band contour in Fig.1.18 demonstrate stress differentiation between boundary region and grain.

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Fig.1.15

Ceramic microstructure of PLZT (a) and PZT ceramics (b) after thermal etching (TEM, replica) (Zhu, 1990)

1.6

Fig.1.16


Smooth or distortion region (arrow) near grain boundary of PLZT ceramics (after thermal etching, TEM)


Fig.1.17 Bending of domain wall near grain boundary of (a) PTC, BaTiO3 ceramics and (b) PLZT ceramics (TEM)

Fig.1.18

Discontinuity of band contour in grain boundary region of PLZT ceramics (TEM) B

üDiscontinuity; GüGrain boundary

Phase transition in PLZT ceramics (from cubic phase into tetragonal or trigonal phase, as shown in Table 1.3) could be induced in PLZT ceramics by applied field, and leads to optical transformation along with following change:

1.6


Properties of α and β phase in PLZT ceramics (Keve,Bye, 1975)

Table 1.3

α phase

β phase

Crystal system

Pseudo cubic Non-ferroelectrics, no domains

Tetragonal or rhombohedral Ferroelectrics, with domains

Optical properties

Isotropy Birefringence Δn = 0

Anisotropy Birefringence Δn > 0

Under crossed polarized light

Lighttight off-state, darkness

Translucent on-state, brightness

Under crossed polarized light and gypsumplate

1st order pink

High order color

Fig.1.19 shows grain boundary morphology of PLZT ceramics slice under crossed polarized light after applying and then removing electric field.

Fig.1.19

Bright grain boundary region after removing applied electric filed in PLZT slice

As shown in Fig.1.19, mostly β phase within grains has transformed into α phase. Due to grain boundary stress, there is residual β phase along grain boundaries shown as bright regions, which also demonstrate the influence of strain field at grain boundary (Zhu, 1990). Residual β phase has been observed in grain boundary via TEM as shown in Fig.1.20 (Chang, 1982). With electrically induced phase transition and changes in unit cell dimensions during phase transition (changes along direction parallel to applied field differ from those perpendicular to applied field), gain boundary region with compressive and tensile stresses could be observed (Zhu, 1984). Quartzite


sandstone, a kind of naturally hot-pressed poly-crystalline ceramics, consists of quartzite grains with different optical axis direction. With colored microscopy images, Allen (1968) has successfully observed allochroic grain boundary region, which is caused by piezoelectric birefringence effects from grain boundary strain.

Fig.1.20

Stress induced β phase in grain boundary region of PLZT (TEM)

Ununiformity could also be introduced into the ceramics during preparation. For example, inhomogeneity could be introduced during dry pressing to cause internal stress. Electric field cross over two electrodes is thought to be uniform but actually is not. Thus the spontaneous polarization Ps within ceramics is actually unevenly distributed. With modulated laser intensity method, Qingrui Yin (1987) has observed macro un-uniformity in distribution of Ps. This un-uniformity may

1.6


cause micro cracks parallel to electrodes after poling.

1. 6. 4 PTC ceramics and internal stress Since resistance of grain boundaries is larger than that of grains in PTC semiconductor ceramics, grain boundaries would become hot spots under applied electric field, and the temperature differentiation between boundaries and grains could reach 50 when measured with infrared microscopy (Mager, Meixner, Kleinschmidt, et al, 1985).Under applied electric field, temperature of PTC ceramics could increase rapidly with a speed of 106 K/s. For certain reasons, including incompetent electrodes, pores, defects, regional resistance differentiation, defects from pressing forming, reaction with setter materials during sintering, abnormal growth of grains, inhomogeneity would be produced to form enormous and instantaneous differentiation in expansion and constriction, which will cause huge internal stress and therefore form micro cracks. For the similar reason, some PTC samples with normal strength have been observed to fracture easily after being loaded electric field. In addition, it has been found that by soaking at temperature 100~150 lower than sintering temperature during cooling, dielectric strength of PTC specimens could be enhanced, which is brought by the crystallization of the second phase as well as stress relaxation within ceramics during soaking. With infrared radiation thermometer and zoom of screen, thermal images of temperature distribution could be observed, which could be used to estimate quality and homogeneity of PTC or varistor ceramics. Hot regions and spots in a thermal image indicate that specimens tend to suffer from failure in future operation under field, and specimens with excessively hot spots tend suffer from breakdown and damage after times of electric loading. Similar method has also been widely used by hospitals to detect hot region for diagnosis. When exposed to high voltage, lightning arrester made of zinc oxide would present uniformity in temperature with infrared thermal imaging, which reveals problems during preparation including de-lamination in pressing and uniformity in density (Amizi, 1986). It has been found that the temperature distribution detected by Thermal Infrared Imager corresponds to resistance distribution, and low resistance region often correspond to high temperature region, i.e. hot region. Cracks often originate from hot region, thus uniformity of temperature distribution could be used to evaluate quality of ceramics. This method could also be used to examine piezoelectric devices of high-power to early locate quality problems. In production, PTC heaters previously qualified by dielectric tests could also fracture from stress (non-electric breakdown) in operation under electric field after being packed by dielectric paper, and the reject rate might be very high, which

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shows packing could change distribution of temperature and stress, so as to cause damage. Due to agglomeration in raw material powders and un-uniformity from additives, abnormal grain growth will be promoted during sintering. Abnormally large grains will present tremendous changes in dimension along crystal axis while temperature varies, possibly leading to cracks between abnormally big grains and ordinary grains. Thus ceramic pieces with abnormally large grains have deteriorated mechanical and electrical properties. For example, with abnormally grown grains, strength of corundum ceramics could decrease by 30%~50%; dielectric strength of semiconducting BaTiO3 ceramics would decrease from 300~400 V/mm to 100V/mm (Ueoka, 1974). In coarse grains there are many cracks which deteriorate both machining ability and strength. In previous Fig.1.3(a), cracks could be found in abnormally grown grains in KNaNb2O3 ceramics, so it is impossible to manufacture small sized delay-line devices with these materials. Coarse grains and cracks could also be found in PTC ceramic microstructure as shown in Fig.1.3(b). J K Lee (2001) has discovered that the second phase of Ba2TiO4 particles could be found in BaTiO3 ceramics with excessive Ba, and diffusion rings (strain field) could be observed at the boundaries of background phase and the second phase.

1. 6. 5 Aging Functional ceramics all suffer from the problem of aging, i.e. property changes over ling time period. In ceramics there are various kinds of defects, impurity ions, and out-of-order region (glass phase), which slowly transfer or change. It has been observed that domain configuration of ferroelectric ceramics, including domain sizes, orientation and domain wall numbers, will change after stocking for a period of time, e.g. domain walls move to more stable positions, and some defects accumulate at domain walls etc. Different concentration of oxygen vacancy in the same kind of material could cause different speed of aging, and high concentration of oxygen vacancy lead to fast aging since vacancy affects stress relaxation. In ferrite ceramics, magnetic permeability will be gradually reduced as time goes, which often results from slowly changing in ordering of some ions (Mn2+, Fe2+). When ceramics are heated up to a certain temperature and then cooled down to room temperature, the process of aging would start over again. In one word, aging is caused by internal stress which will change after heating. In conclusion, internal stress certainly exists in functional ceramics, and its magnitude and distribution would, to some extend, affect material properties. From above facts, internal stress could be divided into two categories: intrinsic

1.7

Capacitor Ceramics

and extrinsic. Internal stress caused by composition, phase transition and anisotropy of crystals is intrinsic, while internal stress caused by un-uniformity of additives, integrality, sizes and distribution of grains, abnormal growth of grains, influence of electric field and temperature, ceramic preparation and machining, or even packing, falls into extrinsic category. To improve internal stress status and mechanical properties of ceramics by modifying composition and ceramic preparation plays a significant role in functional application as well as yield improvement. Microstructural characteristics of several kinds of functional ceramics will be discussed in following sections.

1.7

Capacitor Ceramics

Due to their capabilities in DC blocking, charge stocking, wave filtering, bypass and decoupling and radio frequency oscillator, capacitors are widely used in personal computers, phones, televisions, communication device and other fields. Capacitor is one of the electronic ceramics that are used most widely, with annual production of $ 7~8 billion all around the world (Zhu, 2000). Dielectric constant ε of ceramic capacitor vary from 10~105; film thickness has been reduced from 20 μm in 1980 to 1~2 μm currently; To meet requirement for miniaturization the number of layer of multilayer capacitors has been increased from 40 in 1980 to 400 currently, or even 1100~1200 (Reynold , 2001); capacitor volume has been reduced to only 4% of that in 1980. According to the report in 2002 (Nick Dellow, 2003), Murata co. of Japan has manufactured capacitor with film thickness of 1.6 μm and capacitance of 10 μF, and was working on the capacitor with film thickness of 1 μm. According to the report in 2003, multilayer capacitors with film thickness of 1 μm, internal electrode of nickel, 800 layers and capacitance of 400 μF have been successfully prepared.

ċ

1. 7. 1 Ordinary dielectric materials for capacitor Capacitors could be roughly divided into following categories: high frequency, ferroelectric, grain boundary layer, microwave, monolithic ceramic and other capacitors. Capacitor materials mainly include solid solution of titanate, stannate, and zirconate. Following microstructural characteristics could be given: 1. Except semiconductor grain boundary layer capacitors requiring medium-sized grains, most capacitors require fine grains in materials (e.g. 0.7~1μm grains for BaTiO3 capacitors) to achieve large ε. Additives including NiO and bismuth stannate could be included to restrain grain growth. To obtain flat ε-T curve,


capacitor ceramics are often added with levelers, including MgSnO3, CaZrO3, CaSnO3, Bi2(SnO3)3, CaTiO3. 2. A small amount (2%~3%) of crystal seed with the same composition but slightly larger grain size could be added to achieve more homogeneous structures. The method is often utilized in high ε materials. 3. Chemical ununiformity could be purposely controlled with techniques including utilization of shell-core structures, second phase additive to restrain grain growth, hot-pressing with two-phase or multi-phase mixture (sometimes four phases) etc. 4. Excessive addition of Ba into BaTiO3 ceramics to form liquid second-phase to promote sintering. In SrTiO3 ceramics, the temperature for liquid phase is too high, thus sintering aid (such as SiO2 and Al2O3) or non-stoichiometry is required. 5. With various ε-T relationship, TKε (temperature coefficient of capacitance), and TKf (temperature coefficient of frequency), positive or negative, as well as principle of superposition, proportion of two phases in solid solution could be adjusted to obtain materials with flat ε-T relationship, specific TKε and TKf. One type of capacitor materials with stable capacitance cross over certain temperature range is required. For example, chemical modification could be taken to obtain following system: BaTiO3-Nb2O5-Bi2O3, BaTiO3-Nb2O5-Co3O4, BaTiO3(Ho,Gd,Dy,Er)2O3-MgO and other systems are often used as X7R dielectrics, whose capacity change is no more than 15% from –55 to + 125 . The compositional un-uniformity within grains, i.e. shell-core structure with shell of paraelectric phase and core of ferroelectric phase, could stabilize temperature dependence of capacitance. However, recent development requires thinner film thickness and finer grain sizes, which makes the X7R materials more difficult to obtain a shell-core structure, since ε and capacitance will be decreased with reduced gain sizes and core volume. Instead of making it with shell-core structures, Y Sakabe (2002) doped CaO with following composition: (Ba1̣xCax)mTiO3+1mol%MgO+2mol%SiO2, x=0~0.10, m=1.003~1.009 Tetragonality of BaTiO3 ceramics has been successfully controlled, and stable capacitor materials of fine grains (grain size of 0.2 μm) with high dielectric constant (ε=3000) has been obtained. S H Yoon (2003) has discovered two type of capacitors with difference microstructures and ε-T relationship for different application: One is the materials with excessive donor, in which liquid phase remains at triple junctions instead of entering grain boundaries. Since grain growth has been accomplished by grain boundary migration and the speed of grain boundary migration is quite low, there’s no abnormal grain growth, and

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1.7

Capacitor Ceramics

dielectric property of materials changes only slightly with changing in temperature (little capacitance change from 50 to +100ć), this type materials could be used as stable capacitor . The other is materials with excessive acceptor, in which liquid phase enters grain boundaries. Due to high activity of the liquid phase, abnormal grain growth will be promoted and both large grains of 100μm and small grains of 1μm are produced. Although the additive is macroscopically dispersed, Curie peak in its ε-T curve is quite sharp and the materials is not fit for stable capacitor materials with stable dielectric property. In preparation of high voltage capacitors, Shoutian Chen discovered that previous calcinations play a significant role on microstructure and dielectric property. SBT material was taken as an example: SiTiO3 · Bi2O3 · nTiO2 · MgTiO3, and materials made by two different processes were compared: 1. Method one: all raw materials including SiTiO3, MgTiO3, Bi2O3, and TiO2, were crashed and mixed, then pressed into pallets and finally sintered at 1330ć. 2. Method two: mixed raw materials were calcined at 1180ć, crushed and pressed into pellets and finally sintered at 1330ć. Fig.1.21 shows SEM micrograph of the specimens prepared with two different processes. As shown in Fig.1.21(a), besides main crystal phase, there are many melted globular particles of Bi2O3. Without complete synthesis in method one, Bi2O3 has not been reacted into solid solution and would volatilize at high temperature, which decreases material density. With the second method, Bi2O3 has completely synthesized into solid solution to form needle like layered compound Sr2Bi4Ti5O18, so as to effectively restrain volatilization of Bi, as shown in Fig.1.21(b). The properties of materials derived from the two methods are: for method one, ceramic density of 4.33g/cm3, dielectric strength of 5.84kV/mm; for method two, ceramic density of 5.22g/cm3, dielectric strength of 7.50kV/mm.


Fig.1.21 High tension condenser ceramics (SBT) without calcination (a) and with calcination (b)

In conclusion, calcination plays a significant role on microstructures and properties of materials. In addition, Fig.1.22 shows microstructures of high voltage capacitor SPBT (Sr 1 ̣xPbx TiO3· Bi2 O3 · nTiO2 ). The ceramics have been prepared with similar processes except that a small amount of Nb2 O5 has been added in the second method to restrain volatilization of Bi and Pb and to reduce grain sizes. Material from the second method was shown in Fig.1.22(b), a dielectric strength 25% higher than that from the first method, shown in Fig.1.22(a), which shows that additives have affected grain sizes and therefore affected properties greatly.

1.7

Capacitor Ceramics

Fig.1.22 Microstructures of high tension condenser ceramics SPBT (a) and SPBT + Nb 2O5 (b)

Fig.1.23 shows various high voltage ceramic capacitors, Fig.1.24 and Fig.1.25 show pictures of various multilayer ceramic capacitors, and Fig.1.26 shows AC ceramic capacitors.

Fig.1.23 Various high voltage ceramic capacitors Research center on electrical insulating materials, Xi’an Jiao-Tong University Yi-Sheng Electronic Co., Ltd.


Fig.1.24

Chengdu Hong-Min Electronic Co., Ltd.

(a) Various multiplayer ceramic capacitor; (b) Size are compared with the finger

1.7

Fig.1.25

Capacitor Ceramics

Various multilayer ceramic capacitors (Shanghai Kyocera Electronic Co., Ltd.)

Fig.1.26

AC ceramic capacitor (Shanghai Sanyue Electronic Co., Ltd.)

1. 7. 2 Relaxor ferroelectric materials

¦

Due to the properties of high and electrostriction constant as well as low temperature sintering ability (most containing PbO), relaxor ferroelectric materials are excellent dielectric and actuator materials. A number of relaxor ferroelectric materials


containing Pb have been developed, including PZN, PNN, PMN, PPN, PFW, PMW, and PST, which could satisfy various application and be made as multi-layer devices. However replacing materials of Pb free are required due to environmental concern. In materials of PbMg1/3Nb2/3O3 (known as ABO3 type), degree of order in B site could affect properties. Many small domains (acceptor-type island) have been observed at Nb rich sites (donor-type substrate). These small sized domains (10~20nm) are dynamically under high-frequency switching and difficult to be observed, and they can only be observed at temperature below 0 . The second phase in grain boundary would greatly affect ε. When there are impurities in raw materials, impurities concentrate at grain boundaries to increase ε, and vice versa.

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1. 7. 3 Microwave dielectric materials Recently, thanks to the rapid development in satellite communication, video phones, pagers, high resolution TV, visual telephone and other microwave communication, varieties and output of microwave dielectric materials have increased dramatically. The development tend to increase ε in order to reduce device sizes, increased Q to reduce losses and have smaller temperature coefficient of frequency TCf. Preparation processing affects Q, e.g. larger Q could be achieved via wet chemical precursors, iso-static pressing, proper heat treatment, increment of ordering of B site in ABO3. Zero TCf could be achieved by adjusting proportion of elements at A position (e.g. Ba, Sr, and Ca) in ABO3 structure. W S Kim et al (2000) have indicated that microwave ceramics with TCf (temperature coefficient of frequency) of zero could be obtained by solid solution of CaTiO3 and Li 0.5 Nb0.5 TiO3 ceramics since TCf of CaTiO3 is positive, while that of Li0.5 Nb0.5TiO3 is negative, and both have large dielectric constant. For example, mixed oxide method has been used to prepare material of (1−x) Ca0.4Sm0.4TiO3· x Li0.5Nd0.5TiO3 (x = 0.3), with ε of 98, Q f =5100 GHz, TCf =10 10−4%/ (10ppm/ ) at frequency of 5 GHz. Dongxiang Zhou et al has reduced sintering temperature by 50 with Q f = 6500GHz by preparation method of liquid mixture. Crystal defects might be introduced into MgTiO3 ceramics during sintering, so as to affect Q value. Grain boundary phase (where excessive Mg segregated) could reduce microwave scattering and effect losses. Sintering in oxygen instead of air could greatly improve microwave properties of ceramics. A Feteria et al (2003) has conducted sintering for BaTi0.92Ga0.08O2.96 ceramic in flowing oxygen, and enhanced its fired density and quality factor at high frequency, e.g. Q f of 7815 GHz at frequency of 5.5GHz. Jianjiang Bian (1998) has performed study on microwave dielectrics of BMT (BaMgTaO3) materials. It has been proved that phase purity, uniformity of microstructure, and high density are essential to achieve a high quality factor, while

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1.8

Piezoelectric Ceramics

grain boundary, defects, impurity and ununiformity all affect microwave losses. He has obtain materials with following properties: εr =25.4 and Q f =150000. In recent years, co-firing with silver is required to prepare multilayer microwave dielectric devices, i.e. sintering temperature has to be decreased to no higher than 900 . For example, C H Wei (2003) has introduced 20% 3ZnO·2B2O3 into (Ca1–xNd2x/3)TiO3 and reduced the temperature from original 1300 to 850~900 . In addition, since dielectric resonators and high-temperature superconducting microwave dielectrics (HTSC) require high quality factor as well as low temperature coefficient of resonant frequency, following materials have been developed: (BaPb)O-Nb2O5-TiO2, ZrSnTiO4, Ba(Zn1/3Ta2/3)O3, Ba(Mg1/3Ta2/3)O3, (MgCa)TiO3, with which minimum size and increased stability could achieved. Extremely high Q value is required for application of HTSC, and LaAlO3 is often used due to its high Q (105 at 77K). Actually LaAlO3 is a typical example of all LnAlO3 materials where Ln could be Dy, Er, Gd, La, Nd, Pr, Sm, or Y. It has been found that YAlO3 is more appropriate than LaAlO3 for high-temperature super-conducting microwave antenna due to its excellent lattice matching and thermal expansion compatibility with substrate. LaAlO3 could be synthesized at 1400 and sintered at 1650 , Guohua Huang has decreased synthesizing temperature to 800 and obtained pure LaAlO3 powders by soft chemical synthesis, which will likely be expanded to preparation for other LnAlO3 materials.

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Piezoelectric ceramics are important materials for electromechanical energy transformation, and are widely used in sonar (military application), medical treatment, television, communication, navigation and automation. Smart devices of piezoelectric actuator and supersonic motors have recently become significant development. In 2000, Business Communication Co. USA published a 174 pages report on research development and market survey on piezoelectric ceramics, which was described as importance materials with many significant applications and prosperous developing prospect. In the report, 44 newly applications have been listed, including SKIS micro-robot, actuator for light switch, data actuator, earthquake sensor, smart wing for aircraft, detector for piping, and piezoelectric fibers. Besides, their applications have been promoted by the revolutionary development of wired and wireless communication. In addition, to suffice development of mobile phones communication, monolithic multilayer piezoelectric devices have been recently developed as filter devices. The frequency covers from dozens to thousands of billion, with applications including low voltage actuator, vibration actuator with high speed and high


sensitivity, sounder, vibration sensor and step-up transformer.

1. 8. 1 Microstructures of piezoelectric ceramics Microstructures of piezoelectric ceramics consist of many domain structures, whose spontaneous polarization is randomized. After applied with field E larger than coercive field Ec, domains tend to be reoriented along applied field, and ceramics would present piezoelectricity. If grain sizes are too small, domains are difficult to be developed, thus there won’t have obvious piezoelectricity. Usually open pores are undesirable since they degrade dielectric properties, especially in high humidity; if there are abnormally grown grains of large sizes in microstructures, ceramics are apt to fracture during poling under applied field. In BaTiO3 ceramics, excessive Ti would cause coarse grains. Direction of spontaneous polarization Ps in piezoelectric ceramics could be deduced from 90 domain wall orientation. Materials with certain grain orientation could be prepared by using hot forging, so as to provide strong piezoelectricity along certain direction. In high power application, low heat generation under strong vibration is required in order to reduce thermal damage of devices. C. Sakaki et al (2001) has discovered that fine-grain ceramics are beneficial for high power applications, since fine grains provide large Q so as to reduce the temperature rise. For example, at the largest vibration level of 0.3m/s, temperature rise of fine-grain ceramics (0.9μm) is 40% lower than that of coarse-grain ceramics (3μm), i.e. device temperature could be decreased from 70 to 40 . Fine grains bring more grain boundaries, which impose pinning effect on electric domains, so as to enlarge Q and to reduce temperature rise. However, fine grains also bring down d31 and Ps.

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1. 8. 2 Properties of piezoelectric ceramics With high TC and small ε, pure PbTiO3 ceramics are often applied in high temperature and high frequency infrared sensor and actuator. However, the phase transition from cubic to tetragonal structure during cooling would produce large anisotropic expansion or contraction, which brings large stress to cause fracture. According to Tartaj (2001), since cracks usually originate along grain boundaries and smaller grain sizes means more grain boundaries and strong grain boundary binding, fracture in fine-grain ceramics could be reduced. Others have prepared pure fine-grain PbTiO3 ceramics with Spark plasma sintering, but Tartaj has also obtained pure fine-grain PbTiO3 ceramics with normal sintering at 1100 by adding 10% nano crystal seeds prepared by sol-gel method. In recent years, research on reducing sintering temperature by adding sintering aid in piezoelectric ceramics has been common. Wang (2001) has obtained densified PbZr0.53Ti0.47O3 ceramics sintered at 800 for 2 hours by adding LiBiO2 +CuO, i.e.

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1.8


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the sintering temperature has been reduced by over 350 , but the ε and Kp reached 1000~1300 and 0.55~0.60 respectively. Hayash (2001) has obtained PZT ceramics with Kp of 0.65 sintered at 750 by adding 1 mass % LiBiO2; Via microwave hot-press sintering in oxygen, Takahashi (2001) obtained PZT ceramics with better properties than normally sintered PZT ceramics, as shown in Table 1.4.

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Table 1.4 Sintering method Normally sintering Microwave hot-press sintering

Properties of PZT ceramics under different sintering conditions Coupling coefficient Kp / %

Density / kg·m–3

h10 (7.6~7.65)h10 7.45

3

0.65 3

Piezoelectric constant d 31/ m·V–1

h10 360h10 260

0.75

–12

–12

Table 1.4 fully demonstrate the detrimental effect of pores on properties. Recently developed piezoelectric fibers (φ 43μm (40~50)mm) have been used as high frequency (25~70 MHz) transducer for medical application (Zhang, 2003). PbZnNbTiO3 materials have been discovered with strong piezoelectricity recently, e.g. with a d33 of 2500 Pc / N. Ogawa et al (2002) has proved that single domain could be obtained under proper poling conditions, so as to achieve high piezoelectricity, as shown in Table 1.5.

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Table 1.5 Piezoelectric properties of PbZnNbTiO3 materials under different poling conditions Poling voltage/V

Domain status

ε33/ε0

d 31

d 33

K33/%

K31/%

1000~1500

Single domain

about 4300

1700

2700

95

78

Among ferroelectric and piezoelectric materials, a kind of transparent PLZT (PbLaZrTiO3) ferroelectric materials could be achieved with techniques including raw material dispersion, sintering in oxygen, and liquid phase sintering. Phase transition could be induced under applied field, and provides materials with ferroelectricity. The optical axis of obtained materials could be changed along with changing direction of applied field, thus the materials could be used for electro-optic shutter and memory device, with the later application often in form of films. PLZT ceramics are transparent polycrystalline materials, which are quite sensitive to applied field (electric or stress). Therefore these ceramics are very appropriate for investigation on domains and grain boundaries, as discussed in following section. Due to environmental protection concern in recent years, research has been focused on lead free piezoelectric ceramics, including BaTiO3, KNbO3-NaNbO3, Bi1/2Na1/2TiO3, (Bi1/2Na1/2)0.94Ba0.06TiO3 systems. However their properties have not been as comparable as PZT ceramics. It is recently reported that researchers from Tokyo Institute of Technology of Japan have obtained lead free piezoelectric ceramic BIT with comparative properties of PZT ceramics. Fig.1.27, Fig.1.28 and Fig.1.29 show pictures of piezoelectric ceramic resonators, filters and buzzers.


Fig.1.27 Pictures of various piezoelectric ceramic resonators and filters (Zhejiang Jiakang Electronic Co., Ltd.)

Fig.1.28 Pictures of various SMD piezoelectric ceramic resonators (Shanghai Nicera Electronic Component Co., Ltd.)

1.9

Fig.1.29

1.9

Transparent Ferroelectric Ceramics

Pictures of Piezo-Buzzer element (Betacera Inc.)


The transparent ferroelectric ceramics exhibit some important application in optoelectronic devices 1ke large-aperture modulators, linear light gate arrays, segmented displays, near IR modulators etc.Specific features (grain boundary phenomena and domain dynamics in response to the electric field) of PLZT ceramcis are discussed.

1. 9. 1 Microstructures of transparent ferroelectric ceramics Techniques including raw material full mixing, sintering in oxygen and liquid phase sintering could be utilized to prepare transparent ceramics (Chu, Sun, Yin, 1978). Although the mechanism behind transparent PLZT ceramics has not been fully understood, it is probably due to the introduction of La brings distortion of unit cell, reduces anisotropy of lattice, and promotes homogeneity and densification of grains, and thus forms pore free structures. The authors of this book have observed domain movement in PZLT transparent ceramics, and proved that domains often originate from grain boundary regions with high strain energy, and domains could orientate promptly and joint to cross over numerable grain boundaries under applied electric field to form domain configuration of high orientation. From morphology, it could be found that space charge accumulate along grain boundaries proving that there are grain boundary


regions with several nano-meters from grain boundary. The following phenomena, i.e. the bright grain boundary region, shell-core structure, smooth or distorted region near grain boundary, and bending of domain walls all relate to the strained or disturbed grain boundary region, whose width is from 500 nm to 1000 nm. High homogeneity of hot-pressed PLZT ceramics is quite favorable for generation and transfer of some functional processes. Details will be discussed as following. Functional ceramics have various functions. Microstructures of theses ceramics consist of randomized grains and grain boundaries. Grains present single crystalline structure with atoms regularly aligned and optical axis orientation, while grain boundaries present out-of-order structure often with impurities or second phase, as shown in Fig.1.30.

Single crystal

Ceramics (poly-crystalline)

ügrain boundary ügrain

GB G

Fig.1.30

Polyphase

üquartz grain Mümullite GPüglass phase Q

Pore (empty space)

Crack

ügrain Cücrack Püpore

G

Illustrative diagram for the microstructure of ceramics

To understand the mechanism, the materials properties are characterized normally in macro-scale, However it is more powerful to test and characterize their functions or properties in micro-scale, such as under microscope. Microscopes could be utilized to directly observe generation and development of functional processes. For ferroelectric ceramics, functional processes could be more distinctly reflected from domain characterization. Due to advantageous properties including single phase, pore free, high transparency, phase transition temperature close to room temperature, high sensitivity to applied field (electric or stress), transparent PLZT ceramics prepared with hot-pressing (Haertling, 1999) are quite appropriate for observation of generation and development of these processes. Since magnification of optical microscopes is limited (e.g. 400), ordinary ceramics (with grain sizes of 2~6 μm) are out of resolution limit of optical microscopes, thus coarse grain ceramics (with grain sizes of 15~40 μm) need to be prepared. In addition, to avoid optical interference, a thin sample with thickness equivalent to grain size is needed, thus refraction and reflection of multi-grains under polarized light could be eliminated so that optical changes

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under applied electric field could be easily observed in the experiment. Besides, grain boundary region could act as power resources during functional processes due to its high strain energy, it would be discussed here also.

1. 9. 2 Experimental method and two phases of PLZT ceramics Excessive PbO has been introduced into PLZT ceramics with composition of Pb0.82La0.08Zr0.65Ti0.35O3 for the purpose of liquid phase sintering. Mixed oxide method and long-time hot-pressing in oxygen have been used to obtain transparent PLZT ceramics with coarse grains (15~35 μm) (Zhu, Ao, Yao, 1975). After hot-pressing, excessive lead would be volatilized. The sintered samples were cut and polished into slice of 10 μm, and then deposited with Cr-Au electrodes with a narrow slot (about 0.4 mm) of blank in the middle for light pass and observation. With polarized microscope and stage of adjustable electric field and temperature, optical changes could be observed though the slot, as shown in Fig.1.31.

Fig. 1.31 1

PLZT sample and electrode

üDeposited Cr-Au electrode; 2üResin; 3üGlass slide; 4üPLZT slice

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Thermal etching (at 1150 / 10 mins in lead atmosphere) or chemical etching (hydrochloric acid with a few drops of hydrofluoric acid) could sometimes be performed to expose grain boundaries; ion thinned specimen and its replica are prepared for TEM observation; and crossed electrodes as shown in Fig.1.32 can be used to apply an instantaneous 90 switching of electric field to dynamically observe domain changes. PLZT ceramics could present two phases, i.e. α and β phases, whose properties have been listed in previous Table 1.7. When environmental temperature is over the phase transition temperature Ti, reversible phase transition occurs when applying or removing electric field (Keve, Bye, 1975), i.e. When T >Ti :

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Fig. 1.32

Cross electrode configuration (cross Nicol)( See Color Picture in Appendix)

ė

Ę

(a) electric field=0; (b) field direction ; (c) field direction ; (d) similar to (c) but with gypsum plate

When temperature is below the phase transition temperature Ti, reversible phase transition disappears, i.e.

1.9


When T<Ti :

PLZT ceramics are originally in α phase (un-poled state) with nearly zero birefringence as well as optical isotropy, which present off and lighttight state under crossed polarized light. With gypsumplate inserted, it presents first order pink. After applying electric field, the α phase would be electrically induced into β phases, the ceramic would present on-state and narrow gap turn to bright, as shown in Fig.1.33.

Fig.1.33

Switch on and off state in the region between two electrodes, E = 0 or 1000 V/mm

1. 9. 3 Domain switching properties of PLZT ceramics With gypsum plate inserted, the hue would present colors of blue, green or yellow. The light permitted through the slot would change according to the changes of applied electric field, as shown in Fig.1.34.(a) Non-ferroelectric phase with optical off state when no electric field applied; (b) Ferroelectric phase with optically on state and anisotropy when field applied; (c) Semi-on state when field partially applied.


Fig.1.34 The microstructure (a ~ c) and the corresponding light intensity (I) passing through the gap between the two electrodes (d ) ( See Color Picture in Appendix) (a) E=0, paraelectrics, switch off state, I=500~600; (b) E=E, ferroelectrics switch on state, I >3000; (c) E=1/2E, half switch on, I >1500

1.9


When electric field applied, abundant strip shape domains would emerge, with most domain walls perpendicular to (or at an angle of 75 ~83 with) the applied field, as shown in Fig.1.35.

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Fig.1.35 Domain orientation between the electrode, strip domain crossed several grains, PLZT., optical, crossed Nicol with gypsum plate ( See Color Picture in Appendix)

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These strip shape domains are actually alignment of a large amount of 90 domains. It could also be seen from Fig.1.35 that the domain orientation and continuity are quite in order in the slot region. Liquid phase sintering under hot-pressing could facilitate the rearrangement of grains, and dissolve Zr,Ti,La elements in certain degree, which reduce their lattice stability, and increase the rate of mass transport and reactivity, and benefit densification and homogeneity. However grains in normally sintered ceramics are not so well bonded, as shown in Fig.1.36 (Meng, 1984). Therefore, continuity and transfer of properties among grains in transparent PLZT ceramics are excellent, as shown in Fig.1.37 (colored pages).


Fig.1.36

Comparison of the microstructures of ceramics fabricated by hot pressing and normal sintering

(a) PLZT (hot-pressing, TEM); (b) PTC: Pb0.9Ba0.1TiO3, SEM (Meng, 1984); (c) PTC: BaTiO3 SEM; (d) PZT, chemical etching, TEM;

Fig.1.38 shows structures of grains, grain boundaries, and domains. Similar to Fig.1.37, strip shape domains in Fig.1.38 may cross several grains, demonstrating little orderlessness in grain boundaries. Electron Acoustic Microscopes can be used to observe electric domains crossing grain boundaries. Once electric field is applied, domains would immediately emerge and align orderly, and the functional process takes less than a micro second. As shown in Fig.1.43, wedge-shaped domains can easily cross grain boundaries in hot-pressed PLZT ceramics, but not in normally sintered ceramics (Meng, 1984; Jaffe, Cook, 1971) because pores, impurities, and ununiformity would loose grain bonding, which imposes unfavorable effects on property transfer. As shown in Fig.1.39, after chemical etching and film replica, observation with TEM shows that some domains could cross grain boundaries, but others could not. This may result from different grain boundary angle between two grains, and coherent grain boundary brings excellent continuity between two grains.

1.9


Fig.1.37 Optical microphotograph of domains in PLZT ceramics (domains across grain boundaries, crossed Nicol with gypsum plate) ( See Color Picture in Appendix)

Fig.1.38 Electron acoustic image of grain, grain boundary and domain structure in BaTiO3 ceramics (the sample without polish and etching) ( See Color Picture in Appendix)


Fig.1.39

Domain crossing grain boundary (narrow grain boundary) (a) and domain terminating at grain boundary (wide grain boundary) (b)

Width of strip shaped domains and numbers of domain walls depend on magnitude of stress and strain: high stress brings a large amount of domain walls, while low stress brings less domain walls and large width of strip-shaped domains, as shown in Fig.1.40. If crossed electrodes as shown in Fig.1.32 are applied, adjustment on polarities of the four electrodes could cause 90 switch on electric field in the intersection region, so as to observe domain response under electric field switch (Table 1.6).

e

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Fig.1.40 The effect of internal stress on the domain state in the ceramics ( See Color Picture in Appendix) (a) More domain walls and narrow domain under larger stress conditions; (b) Less domain walls and wider domain under small stress conditions

Table 1.6

The applied electric field direction for domain switching

No.

a

b

c

d

Electric field applied

0

1/2 E

E

E

̨

̨

̩

Semi-on

On state

On state

Field E direction Switch state

Off state

Domain state

No domain

̩

Domain wall

̨

Domain wall

As shown in Fig.1.41(a), with no electric field applied PLZT ceramics exhibit α phase with optical isotropy, and present off sate under crossed polarized light. With gypsum board inserted, it presents Grade-one pink. After applying electric


field, the α phase would be electrically induced into β phases, i.e. ferroelectric phase, and domains are developed with domain walls perpendicular to applied electric field, i.e. polarization within domains is basically along applied field, as shown in Fig.1.41(b). Fig.1.41(c) shows the 90 switch of domain walls under 90 switch of applied field. Fig.1.41(a), (b), (c) show the same region.

e

e

Fig.1.41

Domain orientation under different electric field ( See Color Picture in Appendix) (a) E = 0; (b) E: ; (c) (

ė

Ę

Fig.1.42 is similar to Fig.1.41, but including semi-on state. Fig.1.43 shows nucleation and growth of electric domains from grain boundary region. When

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electric field intensity E increases from 670V/mm to 1000V/mm, domains develop to fill the whole grain, with some wedge-shaped domains even crossing grain boundaries into adjacent grains (as shown in Fig.1.43(c)). With high strain energy, grain boundary regions facilitate domain nucleation, which will expand to the whole grain under larger field intensity. When temperature is below phase transition temperature Ti, ferroelectric phase and its domains could be frozen and preserved. After chemical etching and film replica, domain morphology could be observed with TEM, as shown in Fig.1.44.

Fig.1.42 Domain orientation when the direction of electric field is rotated ( See Color Picture in Appendix) (a) Switch off; (b) Half-open; (c) Open; (d) The direction of domain rotated by 90

e


Fig.1.43 “Wedge” domain nucleated from grain boundary ABC, and the domain tip crossed the grain boundary as E rising ( See Color Picture in Appendix) (a) Electric field E = 670V/mm; (b) E=1000V/mm (optical photograph, crossed Nicols); (c) Same as (a) but with gypsum plate

As shown in Fig.1.44, smooth region near right side of grain boundary is related to the accumulation of space charge along grain boundaries. Due to the difference in electric resistance between bulk and grain boundaries, charge accumulates at grain boundaries when migrating along electric field direction. The charge accumulation affects chemical etching so as to leave various morphologies. If polarity of electric field is reversed, smooth region would move from right side to left side of grain boundaries. This clearly demonstrates the functional process of charge migration within ceramics.

1.9


Fig.1.44 Grain boundary morphology of PLZT (8/67/33) ceramics, TEM image of replica prepared from a sample after chemical etching under electric field E (a) Domains and space charge region appeared at Te < Ti; (b) Space charge remained at grain boundary and remnant domains (as dedicated by arrow) at Te >Ti

1. 9. 4 Grain boundaries in PLZT ceramics Thermal etching has been utilized to investigate energy state of gain boundary region. Polished samples are annealed at temperature 50~100 below sintering temperature for several minutes. Driven by thermal agitation, mass migrate from higher energy region to lower energy region, often accompanied by vaporization and deposition, causes grain boundary morphology and growth steps formed on polished surfaces. Different morphologies after thermal etching correspond to various energy states, and grain boundary region has high strain energy gradient due to its out-of-order structure. As shown in Fig.1.45 and 1.46 (Zhu, 1984), there are many thermal etched pits in each of the grains, and their morphology varies from grain to

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grain due to the fact that the optical axis is related to crystal orientation and the grains are randomly oriented. The random orientation of the grains is also demonstrated by looking at the surface steps in crystal growth as shown in Fig.1.47. Fig.1.48 illustrates the relationship between etched pits and grain orientation.

Fig.1.45 Thermal etch pit in PLZT ceramics (replica, TEM), where EP is for “etched pit” and S for “smooth region along grain boundary”

1.9


Fig.1.46 Thermal etch pit in PLZT ceramics (replica, TEM), where EP is for “etched pit” and S for “smooth region along grain boundary”


Fig.1.47

Surface steps from crystal growth of individual grains of Niobate piezo-ceramics (a) PZT piezo-ceramics after thermal etching (b) (replica TEM)

Fig.1.48 Schematics of etch pits morphology corresponding to different grain orientation (optical axis as indicated by arrow)

Fig.1.45 and Fig.1.46 show smooth region without etch pits near grain boundaries. The grain boundary width is several nanometers, while the affected region of grain boundary may reach 500~ 1500nm. As pointed out by Mistler et al (Mistler, Coble, 1974):

1.9


Effective grain boundary width = mismatch region between grains + space charge region on both sides of grain boundary. The authors believe that the observed smooth region in Fig.1.45 is equivalent to the effective grain boundary width. It was also observed that the smooth region varies slightly from boundary to boundary of different mis-orientation. After thermal etching in lead atmosphere, precipitates of PbO was found in grain boundary region of PLZT ceramics, as shown in Fig.1.49, which also demonstrated that high energy of boundary facilitates nucleation. Similar phenomena are often observed in metals.

Fig.1.49

Precipitates at grain boundary and triple junctions

As mentioned above, at temperature above Ti, when electric field removed, ferroelectric phase turns back into non-ferroelectric phase, so ceramics presents off-state with no domains, and the whole area will be dark. However grain boundary region is still bright with residual domains, which once again demonstrates property differences between within grain and grain boundaries, as shown in Fig.1.50 and Fig.1.51. The author also discovered that ferroelectric domains only emerge within grains, and grain boundaries are non-ferroelectric phase, i.e. two phases in one grain, as shown in Fig.1.52. Since PLZT ceramics are extremely sensitive to stress, stress state in grain boundary region is bound to affect phase transition, which is not


likely to occur in other insensitive materials.

Fig.1.50 Optical micrographs with crossed Nicol showing the remnant domains in grain boundary regions after removing the applied filed E as indicated by : (a), (b), (c) are a few examples and (d) is an enlarged image from location D in (c); direction of field

Ė

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Fig.1.51 Optical micrographs with crossed Nicol showing the remnant domains in grain boundary regions after removing the applied field E as indicated by

Ė

Micron level distortion or smooth regions was observed near grain boundary, as shown in Fig.1.53. Domain wall bending (Fig.1.53(b)) and band contour discontinuity (Fig.1.54) near grain boundary have also been observed. Results showed that the critical field intensity to cause phase transition is related to grain size and clamping state. With electrodes deposited on a thin film sample (reserved a blank gap of 0.4 mm for light pass and observation), electric field is applied to induce phase transition. The gap is dark (off-state) before phase transition, and turns bright (on-state) after phase transition. Table 1.7 lists various switching field intensity for different grain sizes and clamping states.


Fig.1.52 Optical micrographs with crossed Nicol and gypsum plate showing “Core-shell” structure in PLZT ceramics, ferroelectric domains exist only in the center of the grain

1.9

Fig.1.53


Distortion and smooth regions in grain boundary region of PLZT ceramics (a) and bending of domain walls (b)


Fig.1.54

Grain boundary region have optical properties different from within grains

(a) Grain boundary; (b) Grain boundary region; (c) Band contour discontinuance

As shown in Table 1.7, smaller grains and larger clamping would cause higher switching electric field. Table 1.7

Switching field intensity for different clamping states and grain sizes of PLZT ceramics

Grain size / μm

Plate thickness /μm

Switching voltage / V

15~40

10

180~250

15~40

500

200~450

5~10

200

500~680

15~40

ķ

1

>> 1000

ķ Bonded on glass slide by epoxy resin. Zhiwen Yin et al (1988) from Shanghai Institute of Ceramics, Chinese Academy of Sciences has investigated PLZT ceramics of the composition 8/65/35 and 7.9/70/30 with high resolution microscope and found in cubic α phase matrix asymmetrical strip shaped micro zones, i.e. micro-domains with dimension of 5~30 nm. There are more micro domains in 7.9/70/30 material than in 8/65/35 material. According to investigation by electron diffraction, these micro domains have the same structure as polarized β phase (orthorhombic system) under applied field. They also found that the amorphous grain boundaries in these PLZT ceramics are very narrow, with width of around 0.6 nm, or even 0.2~0.3 nm, which are easy for domains to cross.

Grain Boundary Phenomena of Functional Ceramics

This chapter describe various grain boundaries phenomena related to materials properties including: 1. Grain-boundary segregation; 2. Functions of grain-boundary in mass transfer during sintering; 3. Continuity and coherence of grain boundary; 4. Grain boundary under tensile or compressive stress; 5. Grain boundary serves as “source” and “sink” for vacancies; 6. Grain boundary migration and abnormal grain growth in sintering; 7. Grain boundary acts as captive centers & space charge accumulation spots. All the descriptions and discussions are closely connected with concrete functional ceramics.

2. 1

Introduction

As polycrystalline materials, ceramics consist of closely packed grains with different optical axis orientation. Fig.2.1 shows thermally etched surface of piezoelectric ceramic, and Fig2.2 shows SEM micrograph of niobate piezoceramics. From the growth steps or cleavage planes after etching, randomized orientation of grains could be observed. Fig.2.3 shows microstructure of PZLT ceramics, and Fig.2.4 shows microstructure of ZnO varistor containing twin structure, spinel phase and Bi-containing layers. Thus most oxide ceramics have similar structures, but sometimes with twin boundary. Fig.2.5 shows twin boundary in PLZT ceramic, reference (Ogawa,1995) shows twin boundary in BaTiO3 ceramic for PTCR.

2. 1

Introduction

Fig.2.1 SEM micrograph of PZT piezoelectric ceramics after thermal etching

Fig.2.2 SEM micrograph of niobate piezoceramics

Fig.2.3 SEM micrograph of transparent PLZT ceramics

2 Grain Boundary Phenomena of Functional Ceramics

Fig.2.4 SEM micrograph of ZnO material for varistor

Fig.2.5 Twin boundary in PLZT ceramic (optical, crossed Nicol with gypsum plate)

As the most active part in microstructures, grain boundaries could affect many properties and processes in materials. For example, with segregation of Ca and Si on grain boundaries in ferrite materials, a 0.1 μm layer of high resistance will form and reduce loss by 10 times (Paulus, 1996). Addition of CaO and GeO2 into Mn and Zn ferrite materials has similar effects (Hirota, 1966). During cooling process after sintering of PTC materials, a barrier layer of high resistance could be formed at grain boundaries. At temperature lower than TC, because of spontaneous polarization or stress induced polarization, grain boundary barrier is offset to bring abrupt decrease in resistance of materials. Actually the difference between resistance below TC and that above TC could reach 103~108 so the materials could be applied as PTC or self regulating heating element. In capacitor materials of high dielectric constant, substantial ionic diffusion in grain boundaries at high temperatures would lead to chemical valance compensation of Cu, Bi, Mn and other elements and formation of a highly insulating layer (with thickness of

2. 2 Generalization of Grain Boundary

0.1~1μm) in these regions so that dielectric with high apparent dielectric constant ε (dozens of thousands) could be obtained although the intrinsic dielectric constant of the material is only around 1000. In ZnO varistor material, the special relationship between current vs. voltage of grain boundaries could be utilized for various applications. At sintering temperature, there’s a very high carrier concentration within material; during cooling process, mobility of these carriers is reduced so that carrier concentration in grain boundaries would be preserved. The characteristic of grain boundary region with space charge layer is the major mechanism for ZnO varistors. Diffusion of excessive Cu cations along grain boundaries in n-type CdS semiconductor would lead to p-type grain boundaries, which could be used for photoelectric units of solar cell. Grain boundaries could also affect some functional processes in ceramics. Corundum ceramics are normally opaque, but a small amount of MgO dopent could facilitate sintering and densification to obtain transparent corundum ceramics; addition of ThO2 into Y2O3 ceramics and CaO into ThO2 ceramics could slow down grain boundary migration during sintering to eliminate most pores and to obtain transparent ceramics for optical applications. Besides sintering process, grain boundary also plays an important role on diffusion, phase transition, domain formation, internal field, aging, fracture and other processes. Adjustment on grain boundaries could be utilized to control and improve material properties and to develop new materials. Thus scientists who study and develop functional ceramics all pay much attention to grain boundary. In this chapter, based on previous reports during past 40 years, phenomena and properties related to grain boundary will be discussed in detail to attract more attention for development and application of more and better ceramics. Since metallurgists have primarily performed some research related to grain boundary, some examples of metals would also be introduced to explain relevant facts.

2. 2

Generalization of Grain Boundary

Interfaces between substances could be categorized as following:

Hetero-phase interface Interfaces

Gas-liquid (liquid surface) Gas-solid (solid surface) Liquid-liquid (emulsion liquid) Liquid-solid Solid-solid (phase boundary)

Homo-phase interface – Solid-solid (grain boundary)


2. 2. 1 Grain boundary structure Solid-solid interfaces could be divided into two kinds. The interface between two solids of the same phase but with different crystal axis directions is called grain boundary while the interface between different solid phases is called phase boundary. Grain boundary angle is the angle between crystal axes of two grain, as shown in Fig.2.6. Two grain boundaries with low or high angle are called low angle or high angle grain boundary. If θ1 =θ2 , it’s a symmetrical grain boundary, in which twin boundary or superposition might exist. Grain boundary energy is a function of grain boundary angle, and grain boundary energy of twin boundary is often the lowest among all energy with the same grain boundary angle. Usually the thickness of lattice distortion layer along grain boundary is defined as grain boundary thickness, which is less than 5~10nm.

Fig.2.6 Schematic of boundary angle between two grains

Fig.2.7 shows a grain boundary with impurity segregation of B between two grains with composition of A. Grain boundary region usually consists of solid solution, with atoms stacking in a way as shown in Fig 1.44 of reference (Pampuch, 1976). Atoms in grains are in order, while atoms in grain boundary are in disorder where atoms have higher energy than those in

Fig.2.7 Schematic of grain boundary region between two grains

2. 2

Generalization of Grain Boundary

grains. Atoms within grains are bonded from various directions, while atoms in grain boundary and surface are not fully bonded. Actually atoms on surface have even higher energy than those in grain boundary. For example, the surface energy for NaCl is 0.3 J/m2 while interface energy for NaCl is 0.27 J/m 2 , which could be further lowered by absorption of impurities (Zhu, Zhao, 1996). Grain boundary thickness of metal only covers several layers of atoms due to small distortion disturbance, but grain boundary thickness in inorganic oxide is extended because of change of wave energy caused by larger disturbance. In addition, grain boundary thickness differs with different boundary angles. At high temperatures, impurity atoms often diffuse into grain boundaries due to high activity. There would be a strong elastic strain field surrounding impurity atoms within grains, so as to establish relatively high chemical potential. But areas around impurity atoms in grain boundary have low chemical potential because of the open structure and low strain field. The differentiation in chemical potential drives diffusion of impurity atoms into grain boundary to form segregation. Due to dislocations, loose and dense region of atoms will be respectively formed in grain boundary region. Dense region would attract impurity atoms of small radius, while loose region would attract impurity atoms of large radius to reduce stress distortion. In the case of larger differentiation in radius between solvent and solute atoms, grain boundary has an even stronger effect of impurity absorption. Grain boundary could also attract vacancies, and these vacancies might converge to form little cavities. When cooled below specific temperature, excessive cavities would migrate into grain boundaries since it is closer and takes less energy than the migration of cavities to surface. Due to the disorder and open structures, grain boundary is often related with excessive free volume(Meiser, et al, 1980) so that atomic diffusion would be inevitably affected. Activity of defects is obvious only under high disorderliness. At high temperature, atomic diffusion at grain boundary is several orders of magnitude higher than that within grains. For example, oxygen concentration at grain boundary is 10 times higher than that within grains, and the diffusion of oxygen along dislocation in grain boundary could be 106~108 times higher than that within grains. The function of grain boundary for mass transfer during sintering is as important as that of roads for city transportation. Yao has proved that, for PTC materials, the diffusion during sintering with pre-soaking at 1240 is much more satisfying than that with long time soaking at 1350 . It’s because much bigger grains in material sintered at 1350 lead to substantial decrease in amount of grain boundaries which act as pathways for diffusion, while in material sintered at 1240 , mass transfer is instead more effective due to more grain boundaries from much smaller grains(Yao, 1990). The excessive free volume at grain boundaries would lead to reduced atomic

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density with sometimes only 70% of that within grains. Ruhle measured the mean inner potential (MIP) by using electroscopic image. It has been proved that the 15% decrease of MIP at NiO grain boundary is caused by the 15% decrease of atomic density(Ruhle, 1982). High energy of grain boundary could be decreased and transformed into energy required for new phase generation. During re-crystallization or phase transition, nucleations of new phase or re-crystallization often originate from grain boundary. It has been observed in metals that a great amount of nucleation and growth of new phases along grain boundary(Chadwick, 1972).

2. 2. 2 Grain boundary properties As for mechanical properties, grains exhibit typical elasticity due to ideally periodic structure, while grain boundaries sometimes exhibit viscoelasticity due to disordered structures and large free volume. At certain temperatures, grain boundaries could accommodate substantial local plastic flow, which could scatter periodic wave, thus grain boundaries have lower electric and thermal conductivity. Thermal conductivities of single crystal materials decrease as temperature increase, while these of poly-crystal materials increase as temperature increase, which is mainly caused by grain boundary scattering. Grain boundary could act as “sink” and “source” for vacancies in sintering processes, as well as location for absorption and release of strain and stress. Oxide ceramics usually consist of grains with ionic bonds, so defects at grain boundary will lead to boundary charge, i.e. positive charge with excessive cations and negative charge with excessive anions, thus an electric field will be established. These grain boundary charges will be compensated by the opposite space charges near grain boundary, which means a space charge layer extending to certain distance (dozens to one hundred nm) will be induced with opposite charge against grain boundary charges (similar to the double layer in solution). K Lehevoc(1953) has calculated the thickness of space charge layer in the edge of grains of ionic crystals, e.g. 0.22 μm for NaCl crystal at 600K. Coble has put forward the concept of effective thickness of grain boundary region λ, which covers: area of misfit and space charge regions on both sides of grain boundaries. With the only area of misfit, λ of metals is relatively small, while λ of ionic crystals could reach relatively deeper into grains because of space charge regions. For example, λ of Al2O3 is 12.4 nm at 1650 , and λ of MgO is 2μm at 1400 or around 0.5 μm according to others(Peterson, 1983). Diffusion properties of grain boundary region are not applicable for grains. Grain boundary region often act as captive centers, and could attract a large amount of charges to form high capacitance or potential under specific conditions

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2. 3

Grain Boundary Segregation

(e.g. in barrier layer capacitor and PTC). The concept of nonmagnetic grain boundary region (NMGB) has been introduced into magnetic materials. Lattice misfit at grain boundary region would lead to plastic strain, which could block ordered magnetization through boundary thus form a non-magnetic grain boundary. Although thickness δ of nonmagnetic grain boundary is not very large, e.g. δ=1 nm, it plays a significant role on magnetic permeability, especially in grain boundaries with impurity segregation. Thus clean grain boundary is required for high magnetic permeability. Laval et al (1998) have proposed respectively high energy grain boundary with impurity segregation and resultant high barriers, as well as low energy boundary without impurity segregation and resultant low barriers.

2. 3


Strain energy,electrostatic potential and solid solution limit in functional ceramics usually give rise to grain boundary segregation.The possible factors and potential application related to this behavior are analyzed.

2. 3. 1 Generalization Grain boundary segregation is usually caused by following factors: 1. Areas around impurity ions within grains have a strong field of elastic strain, while areas around impurity ions at grain boundaries have open structure and low strain field. Therefore, impurity ions will diffuse from grains into grain boundaries to form segregation so as to reduce strain energy and to eliminate or release stress; 2. Since grain boundary charge increases as temperature decreases, segregation could also be formed in the cooling process. For example, in Al2O3 ceramics with saturated MgO, positive charges at grain boundary could lead to segregation of Mg2+ (whose valance is lower than that of Al3+) to decrease electrostatic potential; 3. At the limit of solid solution, as temperature decreases solubility of solute in solvent lattice would decline thus increase segregation. Normally, solid solution energy in oxide solid solution (energy required for solid solution formation) is quite large with very low solid solution limit, so as to facilitate the segregation of solute. The above mentioned three factors to cause segregation, i.e. strain energy, electrostatic potential and solid solution limit, could make a major contribution under different conditions respectively. Since both strain energy and electrostatic


potential increase as temperature decrease and solid solution limit decreases as temperature decrease, segregation at grain boundary is inevitable when ceramics are cooled down from sintering temperature to room temperature. More segregation would be formed under slower cooling, and less segregation under more rapid cooling, and sub-micro-structural segregation at grain boundary is quite universal. When several impurities coexist in matrix, the element with largest differentiation in ionic radius from that of matrix element will firstly be segregated. Since impurity into lattice will increase free energy of crystals, impurity ions will be excluded from grains during re-crystallization. Through multi-step crystallization, impurity concentration will be greatly reduced in crystals. During sintering of ceramics, accompanied by grain growth and re-crystallization, grains will be purified and impurities will be drawn into grain boundaries. Sometimes, impurity concentration within grains is only (50~100) 10-4%, while that at grain boundary could reach 5%, i.e. 500~1000 times higher, which means grain boundary has the effect of collecting impurities. However some impurities could greatly reduce free energy or even form unlimited solid solution after entering grains, which is not the case described above. Additives could be doped to form a liquid second phase at grain boundary so that some elements will concentrate at grain boundaries. For example, when manufacturing PTC materials with industrial raw materials, Fe, K and other impurities are deleterious to the semiconductorization of grains, so it is difficult to obtain PTC materials with desirable properties. However by doping additives including Si and Al that will form liquid phase, these harmful impurities will be concentrated at grain boundaries so that materials could be successfully semiconductorized(Matsuo, et al, 1968). An other example is that, Mn added into BaTiO3 ceramics will stay in grains, but with addition of Mn and SiO2 to form liquid phase Mn will enter and stay at grain boundaries. During researches on vapor doping, Jianquan Qi has found that boron element concentrate at grain boundaries due to very limited solid solubility of B2O3 in BaTiO3 ceramics, and boron elements are so dense at grain boundaries that the grain boundaries are raised under compression, as shown in Fig.2.8. Under different stress state (tensile or compressive), different segregation will occur at grain boundary. For example, for a Mg-Cr alloy in tensile stress at 500 , at the grain boundaries parallel to the normal direction of tensile stress, hydride denuded zone could be found, which is caused by the diffusion of Mg from grain boundary region under compressive stress to those under tensile stress(Karim, 1969). Vickers has also observed that no segregation at some grain boundaries while some segregation at other perpendicular grain boundaries(Martin, et al, 1976). Since many ceramics do not belong to cubic system, grain boundaries may be exposed to tensile or compressive stress, which could bring different segregation behavior. Concerning segregation at grain boundaries, there are two kinds of facts and

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2. 3


viewpoints:

Fig.2.8 The microstructure of B vapor doped PTC ceramics with different Mn(Qi, 2003) (a)BaSrTiO3+0.25%Y2O3+0.5%SiO2+0.04%Mn; (b)BaSrTiO3+0.25%Y2O3+ 0.5%SiO2+0.08%Mn; (c)Y- BaTiO3

1. There is a second phase continuously surrounding grains with some triple junctions. By selective chemical etching, grains could be removed to obtain grain boundary framework, which could be observe with SEM(Matsuoka, 1981; Mukae, et al, 1981; Stuijts, 1977). 2. The study on grain boundaries with AES, SIMS, SIMS and TEM shows that, instead of the second phase enveloping grains, there’s only a thin impurity layer of several nano meters that have the same structures as that of grains. Referring to the term in metallurgy, it is called as segregation instead of the second phase. Sintering atmosphere also plays a role on segregation at grain boundary. For example, in Nb doped SrTiO3 ceramics, no segregation of Nb has been found when sintered in air; but when sintered in hydrogen atmosphere, Nb segregation has been found in the space charge layer 4~6 nm away from grain boundary,


which is a result of electric compensation for the segregation of some acceptor impurities, e.g. aluminum(Chung, 2002). Some examples will be discussed in the next sections including boundary segregation of materials for BL capacitors, PTC, ferrite and varistors. Segregation has also been utilized to prepare materials for both capacitors and varistors at the same time(Kaino,et al, 1982), and grain boundary engineering could be used to manufacture TiO2 based capacitors with high dielectric constant as well as varistors for low voltage application.

2. 3. 2 Boundary layer capacitors Boundary layer capacitor is a successful instance of controlling structures and properties through grain boundary engineering for electric ceramics. Boundary layer capacitors have microstructural characteristics of semiconducting grains and highly insulating grain boundaries, with the structures equivalent to many capacitors in series and in parallel connection, thus the materials have very high apparent dielectric constant (from dozens to hundreds of thousands) due to the polarizations of space charges between semiconducting grains and insulating grain boundaries. These capacitors are widely used for TV, VTR and radio, and a Japanese factory could manufacture a billion pieces in a year. Manufacturing processes include: firstly doping Nb, Y, La and Dy into BaTiO3 or BaSrTiO3 to obtain n-type semiconducting grains after the initial sintering, then coating surfaces of ceramics with oxides (e.g. oxides of Pb, Bi, and B) that will form glass phases at high temperatures. After the second sintering, these oxides will diffuse into grain boundaries to form insulating layers, so as to obtain the so-called n-i-n structure. Another option is to impel oxygen absorption or metal vacancy (Sr vacancy) to diffuse from grain boundaries to grain surfaces to establish acceptor state of high concentration or interface state of compensation, so as to form n-c-n structure. Jipin Zhong has proved that, for potential barriers of grain boundaries obtained through oxides or diffusion after coating, more acceptor impurities will enter grain boundaries to form higher grain boundary barrier after heat treatment at higher temperatures and longer time, so as to sustand higher voltage. Here are some examples of microstructure analysis on these materials. R Werniche(1981)has proved that slightly excessive TiO2 plays an important role on the liquid phase diffusion during the second sintering. Oxides of Pb, Bi, and B coated on the surfaces will melt with TiO2 at grain boundary to form liquid phase that rapidly diffuse throughout the whole body of ceramics via grain boundaries. At temperature of 1000~1100 , it takes only several minutes to penetrate through all grain boundaries in a sample with a thickness of 10 mm. However the process of melting and diffusion will be relatively difficult without

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2. 3


excessive TiO2. Therefore, the selective melting process is the major reason for the rapid formation of the second phase at grain boundary layer, with a thickness of about 0.1~10 μm and major composition of Pb2Bi4Ti5O18 and some other glass phases. P F Bongers (1981) has prepared SrTiO3 ceramics with grain boundaries only doped with Bi2O3. It could be observed that the second crystalline phase of Pb2Bi4Ti5O18 has enclosed grains. The thickness of the second phase could be obtained with different methods: 0.5~1 μm (SEM), 10 nm (ESCA), 100 nm (TEM), thus it could be regarded that all grains of BL capacitors have been totally enclosed by the second phase at grain boundaries. P E C. Franken (1981) has observed with TEM the diffusion layer containing Bi (close to grain surfaces) and the second phase between two grains, with a total thickness of around 100nm. H D Park (1981) has investigated microstructures of several kinds of BL capacitor materials and electric properties of a single grain boundary. For BaTiO3 based ceramics, grains have the composition of BaSrTiSnO3, and grain boundaries consist of crystalline phase containing Ti and impurities of Si and Al; for SrTiO3 based ceramics, with grain sizes of 20~60 μm, Bi elements mainly stay at grain boundary region with width of 30 nm and extend around 10 nm into grains. The measurement on single grain boundary demonstrated that grain boundary has a resistance of 1011~1012 Ω·cm with non-ohmic I-V characteristic while grains have a resistance of 50 Ω·cm and ohmic I-V characteristic. Park has put forward the energy band structure for BL capacitor material, i.e. a continuously insulating grain boundary layer, enclosed n-type semiconducting grains, and compensation layer on grain surfaces c, so as to form the structure of n-c-i-c-n (grain-n, grain surface-c. insulating grain boundary layer-i,). R Wernicke (1981) has proposed a double-layer model for SrTiO3 grain boundary layer capacitors: the second phase layer in the middle with a different composition from that of grains, and two diffusion layers on the two sides with similar composition to that of grains.The effective dielectric constant ε depends on these layers, especially the second phase d2, which is related to the second sintering temperature and could be observed with SEM. For example, with the second sintering temperature of 900 and 1200 , d2 is 0.9 μm and 0.2 μm respectively (with grain sizes of 20~50 μm), with the second phase layers of 0.1~2.0 μm and diffusion layer of 0.1~0.5 μm in thickness. At higher temperatures, lower viscosity of glass phase will cause smaller d2. At temperatures above 1100 , the diffusion layer is predominant since its thickness d1 increases exponentially with the increase of temperature due to the diffusion mechanism. Grain boundary region in BL capacitor materials conforms to the double-layer model, and Ihrig has also observed the double-layer structure (Ihrig H, 1979). Since the ionic radiuses of Ta and Nb ions (0.07~0.073 nm) are similar to that

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of Ti ions (0.068 nm), M F Yan (1984) has doped Ta and Nb into TiO2 ceramics as donors, and 0.1% Nb could bring 1010 decrease in electric resistance. When Ba element (with an ionic radius of 0.133 nm, much larger than that of Ti ions) is added into TiO2 ceramics, Ba element tends to segregate at grain boundaries during cooling to relax high elastic energy. Auger analysis has proved that the Ba concentration is 150 times higher at grain boundary than that within grains, so that addition of tantalum is offset and it forms dielectric layer at grain boundary with a thickness of 40 nm. The effective dielectric constant ε of the material could reach 105, but problems including high frequency loss need to be solved before it’s put into practical applications.

2. 3. 3 PTC materials To prepare PTC materials, semiconducting doping needs to be introduced into BaTiO3 ceramics to obtain n-type grains, and grain boundaries need to be oxidized to produce Ba vacancy so that grain boundaries are in acceptor state to establish potential barriers with accumulated charges. At temperatures below phase transition temperature Tc, spontaneous polarization and stress induced polarization would compensate grain boundary charge to reduce resistance, and the electric resistance differentiation above and below Tc could reach 103~108, so as to act as materials for over-current protection and self-regulated heating elements. Scientists of electric ceramics from various countries all regard PTC materials as a complicated example of material engineering due to its combined effects of semiconductor, grain boundary and phase transition. In order to explain PTC effect, H Brauer(1967), P Gerthen(1973) and B Hoffmann(1973) believe that a second phase layer of 500~600 nm at grain boundary could form an n-p-n structure, which corresponds to grain-grain boundary-grain structure. Some researchers, from the aspects of defect chemistry and diffusion mechanism, believe that grain boundary barrier mainly originate from grain boundary layer with Ba vacancy of acceptor state, but the existence of the second phase is still uncertain. Investigating with TEM, H B Heanstra et al (1980) have not observed a second phase enclosing grains, but found amorphous second phase with excessive Ti at triple junctions (Bongers, et al, 1981). D R Clarke has done with TEM X-ray micro-structural analysis and discovered both crystal and non-crystal phases between grains. Electron diffraction results show non-crystal phase at triple junctions, with a composition of eutectic phase between BaTiO3 and BaTi3O7 also containing Pb, Al and Fe. Between a few grains, crystal phase layer of 2 nm could be found (Clarke, 1981). M Drofenik (1982) has proved that at grain boundaries of Sb-BaTiO3 ceramic,

2. 3


there are both amorphous glass phase and microcrystal segregation with liquid phase of Ba6Ti17O40 at triple junctions. According to G H Jonker(1981), the second phase is necessary for preparation of both densified PTC ceramics and potential barriers of high quality. With 1% excessive Ti included, 1.5 vol % glass phases will be formed, but mostly at triple junctions according to experiments. Thin glass phase layer (less than 10 nm) enclosing grains can not be eliminated. In PTC materials, with cathode luminescence microscopy, H Ihrig(Ihrig, 1979) has observed green grains and red second phase, as well as addictive gradation near internal boundary within grains. G. Koschek has also observed potential barrier micrograph of grain boundaries and found that the layer thickness varies under different sintering conditions(Koschek, 1985). N S Hari et al (1998) has proved the existence of the second phase enclosing grains as an epitaxial phase. S B Desu et al (1990) also believes that there is donor segregation to form insulating layer at grain boundary, so as to form n-i-n structure. H Hemoto(1981) has investigated resistance dependence on temperature of single grain boundary in PTC ceramics (with relatively large grain sizes of 100 μm), and believed that there’s dielectric grain boundary region with a thickness of 1 μm (observed with SEM) as well as continuous second phase of microcrystal. The acceptors and barriers in grain boundary region would decrease mobility of electrons, so that grain boundary resistance is predominant even at temperature below TC. Above mentioned factors including the second phase at grain boundary, grain boundary insulating layers, surface layer of grains, subsurface of grains, the second phase as grain epitaxy, and point defects frozen in grain boundary region could be used to describe various defects and un-uniformity in grain boundary region. Since different compositions (various content of excessive Ti, various doping of Si, Al and other additives), different sintering conditions (atmosphere, heating and cooling rates, etc.) would cause differentiation in grain boundary segregation, inconsistent results might be obtained by different authors. However, all these results are actually linked to space charge layers or potential barriers in grain boundary region, which have been directly observed. R D Roseman et al (1998) have proved that structures of grain boundaries and electric domains play a key role on PTC effect. Roseman performed high temperature poling on PTC ceramics, and observed obvious PTC effect perpendicular to the poling field but weak or no PTC effect parallel to the poling field. He explained the weak PTC of the latter situation as a result of relatively higher coherence of grain boundaries. K Hayashi (1996) measured the resistance-temperature (ρ-T) relation of single grain boundary in ceramics, and discovered that PTC effect for highly ordered


grain boundaries including small-angle or twining grain boundary is weak or absent; while PTC effect for highly randomized and disordered grain boundaries with many defects is strong. This is because disordered grain boundaries facilitate segregation of acceptorimpurity, diffusion and absorption of oxygen, and generation of compound defects, which are produced during cooling after sintering to form grain boundary barrier and PTC effect. Thus PTC effect is closely related to coherence and mismatch states of grain boundaries. As for acceptor state of grain boundary layer, some (Jonker, 1964) believe that acceptors exist at grain boundaries in the form of oxygen absorption to reduce Ti3+ concentration, so as to form depletion layer and potential barrier. Others suggest that genuine donors are intrinsic oxygen vacancy from oxygen volatilization, e.g. in SrTiO3 ceramics (Burn et al, 1982), or in the form of oxidized impurities (Kahn,1985). According to one viewpoint, these acceptors will provide deep trap; according to another viewpoint, potential wells lie in layers of various thickness under grain surfaces, and are generated from concentration gradation of Ba vacancy. Both potential wells would lead to potential barrier structures. By using cathode luminescence method with integral spectrum, G Koschek (1985) has observed that the light-colored grain boundary region in 0.2 mol% Y2 O3 doped PTC material, is also the region with high concentration of Ba vacancy (with a thickness of several microns), which is related to cooling rate in sintering. By adjusting sintering and cooling, oxygen diffusion and oxidation state in grain boundary region could be controlled and altered. With excessive oxygen, Ti3+ concentration will be reduced to form depletion layer, which will alter the resistance in the resistance vs. temperature curves. This could be actually regarded as segregation of oxygen in grain boundary region. Recently, thanks to the development of materials with high Curie Point, solid solution of PbTiO3 -BaTiO3 ceramics have been investigated. Z C Huang (1987) prepared LaBa0.55 Pb0.45TiO3 ceramics by doping Si, Al, Mn and 1 mol% excessive Ba, and found low La concentration, high Si, Al and Mn concentration, as well as high Ba concentration at grain boundary region with thickness of 1μm, which seems to disprove the viewpoint of “Ba vacancy at grain boundary region”. However, for high temperature PTC materials with 45 mol% PbTiO3, the result might differ from that for pure BaTiO3 materials. In addition, concentration of PbO in PbTiO3 solid solutions could reach several percentages and varies according to different sealing conditions during sintering, which make the situation more complicated. When discussing segregation at grain boundary in BL capacitors and PTC ceramics, solubility of TiO2 in BaTiO3 ceramic is quite limited. With concentration higher than 0.1 mol%, TiO2 will be segregate out, and the Ba/Ti ratio is significant. With Ba/Ti ratio of 1, no second phase will be produced; with Ba/Ti ratio higher or lower than 1, second phase of Ba2TiO4 or Ba6Ti17O40 will be produced respectively.

2. 3


This might partially explain various results reported by different authors. BL capacitor and PTC materials are similar concerning n-type semiconducting grains and p-type semiconducting grain boundaries. However BL capacitor materials have higher p-type doping concentration and thicker p-type grain boundary.

2. 3. 4 Magnetic ceramics Soft ferrite materials are magnetic materials with low coercive field, with major applications in communications and electronic toys. With application frequency covering audio frequency to millions of mega Hz, magnetic materials could be divided into two major categories: Mn-Zn and Ni-Zn systems. Generally, high magnetic permeability μ and low loss are desired, and resistance needs to be increased to reduce eddy loss. T Akashi (1966) found that addition of CaO and SiO2 could increase resistance by 2~4 orders of magnitude and Ca element are segregated at grain boundary according to observation with isotope analysis (Paulus, 1966). Intensification of oxidation during cooling after sintering could also increase electric resistance. For example, Ni0.4Zn0.6Fe2O4 ceramic sintered at 1300 and then rapidly cooled had a resistivity of 10 Ω·m and FeO concentration of 0.42%, while the resistivity could reach 1000 Ω·m and FeO concentration could be decreased to 0.07% after the sintered ceramic was re-fired and then slowly cooled down (Moulson, Herbert, 1990). It is because that oxygen would diffuse along grain boundaries during slow cooling to form highly insulating grain boundaries and semiconducting grains. Bando (1971) also thought that there was grain boundary phase of high resistivity enveloping grains. M F Yan (1984) explained from the aspect of ionic sizes: Ca ionic radius (0.099 nm) and Si ionic radius (0.041nm) are far away from those of Mn, Zn and Fe ions (0.06~0.08 nm) in matrix, and Si4+ ions are under large electrostatic potential, so these ions tend to segregate, while Sn ions (0.07 nm) do not likely to segregate due to their good compatibility with matrix. Auger analysis has proved that Ca and Si concentration at grain boundary is hundreds times higher than that within grains. The layer abundant in Ca and Si elements is less than 40 nm, and the non-crystal layer of Ca and Si is around 3~5 nm in thickness according to the observation of TEM (Tsunekawa, 1979). From the lattice fringe micrograph of MnZnFe2O4 ceramic containing Ca and Si (Mishra,1981), there is a amorphous phase at grain boundary (Mishra, 1981). Analyzing the intersection with AES, Franken found out that the thickness of Ca and Si layer is 2 nm without other second phases (Franken, 1980). There are some micro zones of second phase at triple junctions. If the micro zones are large, stresses produced during cooling could cause fracture. In stead of glass phase of high resistivity, they believed that grain boundaries of high resistivity were caused by oxidation of grain boundary region and segregation of CaO.

ć


To improve magnetic permeability vs. temperature (μ-T) performance, Ti could be introduced into materials (Stijntjes, 1970). Investigation with AES showed that gradation of Ti distribution could be established at wide grain boundary region and within grains, so as to improve temperature dependence of μ property(Franken, 1978). However addition of TiO2 could also generate a small quantity of liquid phase to promote abnormal growth of grains, which will increase the loss (Yan, 1978). Non-stoichiometric oxygen distribution at grain boundary region also plays a significant role on magnetic permeability and mechanical strength. Chiang (1983) proved the existence of oxygen concentration at grain boundaries is close to that of material surfaces. Addition of SiO2 into La0.6Sr0.4MnO3 materials could increase magnetic resistance property, which is also caused by the effects of grain boundary (Takemoto, Jamagiwa, 2002). In conclusion, grain boundaries could affect properties of ferrite materials through following mechanisms: 1. To form high resistivity barriers between grains; 2. To affect sintering and grain growth by providing impurity segregation or liquid phase; 3. To provide paths for oxygen diffusion to alter grain boundary composition and properties.

2. 3. 5 ZnO varistor materials Thanks to its unique voltage vs. current (I-V) characteristics and voltage dependant resistor (VDR), ZnO ceramics have been widely used in electric circuits and power system since 1968 for voltage stabilization and surge absorption. It had been believed that the non-ohmic properties of ZnO ceramics doped with Bi2O3 were caused by the second phase produced at grain boundaries. Morris (1973) obtained residual grain boundary phase with clear structures by using selective etching to remove grains, which provided basis for the existence of the second phase along grain boundaries. SEM micrographs also showed the continuous second phase enveloping semiconducting grains, and this seems to be doubtless. However, Clark proved with TEM observation that grain boundary phase containing Bi was not wettable for ZnO grains and only lay at triple junctions (Clark, 1978). Lattice fringe images also confirmed no second phase at most grain boundaries. W G Kingery (1976) performed X-ray microscopic analysis with STEM and found Bi segregation region (less than 10 nm) along grain boundaries, but without a second phase. Thus some scientists believe that VDR characteristics are caused by grain boundary segregation instead of medial

2. 3


layer separated from grains. According to P. F. Bongers (1981), there are second phases including spinel, pyrochlore and Bi2O3, and skeleton of second phases could be obtained by removing ZnO grains with etching, but the contribution of second phases to VDR effects is neglectable. It has been proved that ZnO ceramics without Bi2O3 addition but sintered under oxygen atmosphere also show VDR effect (Hausmann,1975;Einjiger,1978; Fujitsu,1987). Therefore, oxidation layer at grain boundary (chemical absorption and diffusion of oxygen into grain boundary) is the key factor for VDR effect, while the existence of Bi element only facilitates oxidation of grain boundaries. Yao Yao investigated reduction and oxidation of ZnO ceramics and found that varistor mechanism is related to the absorption and loss of oxygen only at grain boundary, which was also confirmed by experiments with O18. Although it is controversial concerning the effect of the second phase, it is doubtless that there is a space charge layer between two grains, and grain boundary barrier could be adjusted by heat treatment (Kim, 1986).

2. 3. 6 Other examples of segregation There are also many other examples of segregation in materials. For example, when 0.25% Fe is doped into BaTiO3 ceramic, Fe concentration within grains is 0.08%~0.15%, while that at grain boundaries with a thickness of 40 nm could reach 0.4% (Oppoljer, 1979). Grain boundaries in BaTiO3-NaNbO3 ceramic would turn bright due to segregation here. Yan doped Nb, Ta into TiO2 ceramic to obtain resistance of 10~100 Ω, which is due to insulating barrier caused by segregation at grain boundary during cooling process after sintering. By adjusting partial pressure p of oxygen , nonlinear constant α could be altered, i.e. α would be reduced with decreasing p. In magnetic ceramics, grain boundaries also play a significant role. For example, in Mn-Zn or Ni-Fe ferrite materials, properties mainly depend on magnetic domain structures, which are greatly affected by grain boundaries. Grain boundaries could affect magnetic domain width and domain wall mobility at equilibrium; furthermore magnetic domains nucleate at grain boundaries. Generally three attributes of grain boundary need to be considered, i.e. lattice orientation, disorder and stress states, and impurity segregation. Magnetic permeability μ of ferrite materials is proportional to grain sizes ( or grain boundary numbers), i.e. lower μ with smaller grains. Ca and Si could be added into Mn and Zn ferrite materials to increase grain boundary resistance (10~100 times higher) to reduce eddy current loss, and this method has been taken for 40 years to improve material performance, but its mechanism has not been revealed until recent year. Since addition of Ca element could cause grain boundary segregation and other effects, in order to improve electric resistance as well as magnetic permeability,


following factors need to be considered comprehensively: 1. Grain boundary segregation will inevitably affect grain boundary migration and grain sizes, so as to affect magnetic permeability; 2. Sintering atmosphere could affect Fe2+ concentration and consequently affect magnetic permeability. Meanwhile Fe2+ concentration would restrain Ca segregation since Ca ions could enter octahedron lattice positions, which are often engaged by Fe2+, thus Ca segregation decreases with increasing Fe2+ concentration. In addition, atmosphere could also affect pore structures, which could affect grain boundary migration so as to affect consequent magnetic permeability; 3. Plastic strain from lattice mismatch of neighboring grains could generate nonmagnetic grain boundary (NMGB) with a thickness of δ. Impurity segregation or composition change could produce low-μ region or NMGB, which could also be produced with reduced Zn concentration at grain boundary. Although the thickness of the nonmagnetic grain boundary is limited, e.g. δ≈1 nm, it will block ordered magnetization through grain boundary so as to affect μ; 4. μ depend on porosity and distance between magnetic domain walls. In one word, properties of magnetic ceramics are related to grain sizes, grain boundary numbers, nonmagnetic grain boundary and grain boundary stresses. J Y Laval et al (1998) has systematically investigated electric barrier, crystallographic and chemical composition of individual grain boundary in Mn, Zn ferrite materials, and proposed three kinds of grain boundaries: 1. Ordinary grain boundary without specific orientation between neighboring grains. These grain boundaries are highly disordered with high energy, where Ca could be segregated and Fe2+/Fe3+ ratio would be reduced to produce grain boundary of high resistance and high potential barrier; 2. Low energy grain boundary with orientation between neighboring grains, with neither Ca segregation nor Fe2+/Fe3+ ratio variation. These grain boundaries have large loss due to low potential barriers; 3. Grain boundaries with glass phase, and potential barriers are determined by its composition. For the samples measured by Laval, the three kinds of grain boundaries, i.e. 1., 2., and 3., accounted for 80%, 16% and 4% respectively. High diffusion of oxygen at grain boundaries in ABO3 ceramics could be utilized to introduce excessive oxygen or “oxygen segregation” at grain boundary and to finally eliminate it via diffusion. During hot-press sintering in oxygen, oxygen instead of nitrogen in pores in PLZT green body will be eliminated easily through diffusion of oxygen vacancy, while nitrogen in pores is difficult to remove. Pores are removed and transparency is increased when ABO3 ceramics are sintered in oxygen, which is lucrative to obtain materials with low porosity and high density, as shown in Fig.2.9 (Chu, Sun, Yin, 1978).The work of K H H噀rdtl s investigations have indicated that in Perovskite materials the rate of diffusion of oxygen at the sintering temperature is

2. 3


several orders of magnitude high then that of nitrogen. (Chu, Sun, Yin, 1978).

Fig.2.9 The effects of sintering atmosphere on transparency and microstructure of PLZT ceramics

ü

(a)Effect of atmosphere on transparency of PLZT ceramics;(1 Hot pressed in oxygen, optical micrograph, transmitted light; 2 Hot pressed in nitrogen, optical micrograph, transmitted light (black spots – pores)); (b)Effect of atmosphere on microstructure of PLZT ceramics; (1 Hot pressed in oxygen; 2 Hot pressed in air (P-pores))

ü

ü

ü

According to Reijnen, at high partial pressure of oxygen, high mobility of pores at grain boundary could generate enough cation vacancy to facilitate bulk diffusion of cations and gas migration of oxygen (Reijnen, 1968). A small quantity of oxide addition into Al2O3 ceramics could improve creep resistance. For example, 0.05 mol% Lu2O3 into Al2O3 ceramics would be segregated at grain boundary to obtain this effect (Ikuhara, 2001). To remove


pores and obtain transparent corundum ceramic, 0.05%~0.5% MgO could be added and separated at grain boundary to slow grain boundary migration and restrain abnormal growth of grains, so as to prevent the captivity of pores within grains (Jorgensen, 1964). Similar mechanism could be utilized to prepare various transparent ceramics, including ThO2 (with addition of 2 mol% CaO), ThO2 (with addition of 5 mol % Y2O3), ZrO2 (with addition of 6 mol % Y2O3), and Y2O3 (with addition of 10 mol % ThO2 and 1 mol % Nd2O3). It was known previously that addition of MgO to obtain transparent corundum ceramic was due to segregation of MgO at grain boundary, but it has been now discovered that no MgO was found at grain boundary even if the addition reached 0.3%. However, Ca segregation at grain boundary was observed: Ca concentration in grains is lower than 0.01%, and reaches 2% at grain boundary. According to Yan, since ionic radius of Ca is larger than that of Mg and does not match that of Al, elastic energy of Ca ions is estimated to be five times higher than that of Mg ions, so Ca element is easier to segregate at grain boundary. In addition, in samples only containing Mg without Ca, Mg will be segregated at grain boundary (Yan, 1984; Yan, 1977). In the investigation on segregation of Al2O3 twin crystal, width of Ca and Si segregation region could reach 60~80 nm, and grain boundary layer is 55nm in thickness (Anon, 1979). Li believes that misfit between ions is the major driving force for segregation (Li, 1985). Effects of MgO in Al2O3 ceramics could be found in reference (Dorre, et al, 1984). It has been discovered that amorphous phase rich in PbO with a thickness less than 10 nm could be found at grain boundary in sintered PZT ceramics with 3 mol% excessive PbO. During cooling, no segregation occurs in the amorphous layer, which could affect poling. During poling, almost half of the DC voltage is acted on grain boundaries, which also provide high capacitance (Goo, et al, 1980). Amorphous phase with a thickness of 10 nm at grain boundary has been observed with TEM in PZT ceramic (w(Zr)/w(Ti)=52/48), and Pb concentration at grain boundary is higher than that in grains(Goo, 1981). TEM observation and micro-structural analysis demonstrated that high Pb amorphous phase plays a significant role on poling in spite of a small amount (Mishra, 1981). Electric properties could also be affected if Mn and Cr elements are un-uniformly distributed in grains and grain boundaries (Mishra,1981). Characteristics of high density of dislocation and gradation of strain, local defects and high energy, make grain boundaries locations for nucleation and segregation (Matin, 1976). Nucleation possibility varies as following: intersection of four grains > triple junction > interface of two grains (Azaroff, 1960). Grain orientation at interface plays a crucial role on nucleation of precipitation. After thermal etching in lead atmosphere, precipitates at triple junctions of hexagonal grains could be observed in PLZT ceramics, as shown in Fig.2.10 (Chu, 1984). Y J Chang (1982) has observed β phase induced by strain field in grain boundary

2. 3


region during TEM investigation on PLZT ceramics, as shown in Fig.2.11. Similarly, precipitation region extending from grain boundary to both sides has observed in stainless steels, and is considered to be related to the strain field in grain boundary region (Jones, 1976). However, no precipitates could be found at some grain boundaries with low stress due to near same orientation of neighboring grains. Concerning metal materials, grain boundary precipitation or segregation could affect properties including strength, hardness, and corrosion resistance. After some stainless steels are processed, precipitates of abundant Cr would be generated at grain boundary to deteriorate corrosion resistance. However nitrogen element could be introduced to segregate at grain boundary to prevent Cr from segregation so as to improve corrosion resistance. This is a typical example of replacement of one type of segregation by another (Briant, 1987), similar to the segregation of Ca instead of Mg, and segregation of Mg when no Ca is added in alumina ceramics.

Fig.2.10 Precipitates deposited at the triple junctions(Chu, 1984)

Fig.2.11 Stress induced

£-phase in the grain boundary regions(Chang, 1982)


2. 4 Grain Boundary Region Grain boundary regions with high energy affect substantially various functional effect.Here we describe their features in BaTiO3 and PLZT ceramics,discuss the relationship between the grain boundary region and stresses,core-shell structures.

2. 4. 1 General description about grain boundary region As shown in Fig.2.3, ceramics contain many grain boundaries. When further zooming in, e.g. to magnification of 10000 , grain boundary could be observed, as shown in Fig.2.9. Grain boundaries are about 0.5~5 nm in width, equivalent the size of several or dozens of cells, but it could affect distance beyond 5 nm. Actually one atom vacancy could interfere with neighboring atoms and cause to shift 0.x% to 26% atomic distance (Hahn, 1981). Besides, space charge at grain boundary could extend rather deeply into grains, e.g. 0.2 μm, similar to double layers in solution-solid interface. The strain field caused by disordered structure of grain boundary could affect optical, electric and mechanical properties in a certain region, which is called grain boundary region. By testing relevant properties, thickness of grain boundary region could be obtained (generally from 5 nm to several microns, depending on many factors). Sometimes there are precipitates in grain boundary region, sometimes there’s no precipitate to form so called “clear region”. Optical variation of clear region could be observed by using TEM, and pores have been swept out in some grain boundary regions. Grain boundary region is usually under stress, and tensile or compressive stresses could sometimes be distinguished. Due to different energy or space charge, grain boundary region would present different morphology after thermal or chemical etching. Since grain boundary region has different elastic properties from grains, acoustic micrograph could be obtained by using acoustic microscope. Thermally conductive properties of grains and grain boundaries could be observed by using thermal wave microscopes. Grain boundary region has a width of 5~1000 nm, and accounts for 3 vol%~30 vol % of bulk ceramics. Because of having high energy, grain boundary regions, as an active part in microstructures, could affect or participate in many functional processes, thus understanding of these regions are necessary. Specific examples as following will be included to discuss grain boundary region.

h

2. 4. 2 Grain boundary region of BaTiO3 ceramics According to D R Callaby (1965), on the surface of BaTiO3 single crystal there’s a surface layer with lower ε, lower strain (proved by electron diffraction tests) and

2. 4 Grain Boundary Region

larger domain wall motion than that within crystal. Electric domains tend to nucleate in the interface between crystal inside and surface layer. Surface layer has a thickness of 0.2~0.5 μm, or 0.2~10 μm according to others. For example, BaTiO3 whiskers of 14 μm in diameter have a surface charge layer with a thickness of 1μm. With electric field applied, surface layer is bright under crossed polarized light, and needle like electric domains nucleate within the layer and develop into inside body of crystal (Callaty, 1965). It is reported that the space charge in BaTiO3 ceramic is mainly distributed at grain boundary. The optical properties of a film with a thickness of an individual grain demonstrated that grain boundary region is 2~5 μm in width (with grain sizes of around 30 μm) and its optical properties differ from those of grains. Since grain boundary space charge would impose shielding effects against applied electric field, effective electric field for grains would be decreased, which also explains aging of materials. When an electric field E is applied on BaTiO3 ceramic, grains and grain boundaries perpendicular or parallel to electric field would take on different response. Electric field first affects grain boundaries. After removing applied electric field, grain boundary region will not recover until reheated above phase transition temperature.

2. 4 .3 Grain boundary region of PLZT ceramics Under crossed polarized light, as-sintered and polarized PLZT slices would present different grain boundary stress states, as shown in Fig.2.12. Grain boundary stress field could be observed as well as calculated by using crossed polarized light (Zraichenko, 1973). TEM observation has found that region of 1μm in thickness near grain boundary might be strain field region, as shown in Fig.2.13. We have conducted thermal etching on PLZT ceramics to directly observe grain boundary region, as shown in Fig.2.14, and found etch pits with various configuration (spot, line or rod) corresponding to different optical axes of grains.

Fig.2.12 Different grain boundary state of PLZT ceramics (optical micrograph) (a) As-sintered or thermally depoled, transmitted light, crossed Nicol; (b) DC electric field removed, 1kV/mm, crossed Nicol; (c) Same location as (a) and (b) but parallel


Fig.2.13 Strain field region near the grain boundary of PLZT materials

Thermally etching drives mass transfer and migration from high energy region to low energy region, so originally polished surfaces would present unevenness related to energy. Smooth region near grain boundary demonstrates that grain boundary region has different energy from grains. Since neighboring grains have various orientation and grain boundary angles, different grain boundary width, defect density and atomic energy could be produced (Zraichenko, 1973), so as to present different clear region. H Hemoto (1981) has investigated relationship of resistance ρ vs. temperature T of an individual grain boundary of PTC ceramic, and found that different grain boundary has different ρ-T relationship, which is likely to be caused by similar reason.

Fig.2.14 Smooth region of PLZT in the immediate vicinity of grain boundary after thermal etching (replica, TEM)

Relationship between grain boundary thickness Z and impurity concentration as well as grain sizes D is listed in Table 2.1 (Carnigla, 1966).


Table 2.1 The effect of grain boundary impurity concentration and grain sizes on the thickness of grain boundary Grain boundary impurity concentration / % Grain size D/μm

0

0.1 G.B. thickness Z/nm

50

0.1

500

1

5

0.1

50

1

>1

>0.2

5

1

G.B. thickness Z/nm

G.B. thickness Z/nm

100

1

0.002

10

1

0.02

1

1

0.2

ü

Note: Vb grain boundary volume; Vt

1.0 Volume fraction Vb/Vt

Volume fraction Vb/Vt

ütotal volume of ceramic.

Volume fraction Vb/Vt

In grain boundary region of disordered structure with abundant captive centers, barriers tend to be formed here, and depletion layer will be formed for ferroelectric ceramics. Being able to affect various properties of materials, space charges will accumulate at interfaces. Being able to accommodate stress and strain, grain boundary region could act as resource for some irreversible processes, e.g. relaxation. The region first responds to the applied field, and would keep partial residual effect after applied field removed. The alignment of electric domains at grain boundary region in ferroelectric materials could compensate charge of potential barriers on surfaces of grains, as shown in Fig.2.15, which provides Fig.2.15 Grain charge compensation by Ps an important basis for PTC effect. being aligned in 180 domains below TC We have observed alignment of ordered domains near grain boundary region in PLZT ceramics, as shown in Fig.2.16. Ununiformity of electric field in grain boundary region could facilitate electriferous particles, e.g. TiO2 particles in silicon oil, to precipitate near grain boundaries so as to make grain boundaries visible (Carnigla, 1966). Grains and grain boundary regions will present different morphologies in chemical etching, as shown in Fig.2.17, and smooth regions appear on one side of grain boundary. When applied field is reversed, smooth regions will appear on the other side of grain boundary. In addition, due to differentiation in space charge, grain boundaries perpendicular and parallel to electric field will also present different morphologies. By using emission electron microscope, barrier layer at grain boundary in semiconducting BaTiO3 ceramic could be observed (Rehme, 1966). Shadow region, about 5 μm in width (with grain sizes of 40 μm), on one side of grain boundary will move to the other side when applied electric field is reversed, Comparing Fig.2.17 with reference (Rehme, 1966),

e


we think these smooth regions are related to space charge. H B Heanstra (1977) has observed grains and grain boundaries of different hue caused by different electron concentration in PTC materials, and estimated the direction of currents.

Fig.2.16 Grain boundary with electric domain, on removal of the electric field E (1000V/min, direction of E: ) crossed Nicol

Ė

Fig.2.17 Space charge in grain boundary region and corresponding schematic graph (a) Remnant domain and space charge in grain boundary region after removing electric field E (chemical etching, electron micrograph, E direction: ); (b) Schematic graph of space charge polarization

̩


J T C Kemenade (1978) has observed potential barriers at grain boundary in ZnO ceramic. Y M Chiang (1982) investigated ZnO ceramic under applied field E, and found asymmetric distribution of Bi element cross grain boundary perpendicular to applied field, and symmetric distribution of Bi element cross grain boundary parallel to applied field. It means that Bi element would migrate along applied electric field and accumulate on one side of grain boundary, which is similar to cases in Fig.2.29. Applying DC voltage on SrTiO3 BL capacitors, potential difference of grains could be observed (Okazaki, 1981). Strain field and deficiency of solute ions in grain boundary region would form non-precipitation region on both sides of grain boundary (Jones, 1976; Sellars, 1976). Burnett (1978) conducted rapid cooling on highly pure steel and observed similar situation. However, opposite situation also occurs, i.e. typical discontinuous precipitation is formed along grain boundary region (Mclean, 1957). According to Unwin, vacancy concentration is much lower than that in grains, so as to act as an ideal sink for vacancy. The factor of no precipitation could be explained from the diffusion of vacancy and solute (Unwin, 1969). Width of grain boundary used to be regarded as distance of several atoms, but some scientists believe there is a wide grain boundary region or effective diffusion distance. Due to rapid diffusion of vacancy, hardness near grain boundary is reduced, thus quenching defects could be eliminated or not generated at all (Westbrook, 1968).

2. 4. 4 Grain boundary region and stress Vacancy will migrate under effect of stress gradation from grain boundary of compressive stress to those of tensile stress, i.e. Nabarrohering mechanism of vacancy migration (Herring, 1950). Herring has found frequent deficiency of hydride in grain boundary region under tensile stress, which is caused by atom migration under stress to affect precipitation at grain boundary (Burke, 1969). For PLZT slices, when electric field applied, phase transition will be induced; when electric field removed, grain boundaries perpendicular and parallel to electric field would be in different stress states (tensile or compressive), as shown in Fig.2.18 (Chu, 1984). The differentiation in colors is resulted from grain boundary stress caused by expansion and contraction of crystal cells during phase transition. Most ceramics are prepared with sintering and cooling processes. During cooling, stresses will be established in non-cubic ceramics, and grain boundaries are divided into tensile stress grain boundary and compressive stress grain boundary. In the investigation on silicon segregation in aluminum, precipitate denuded zone has also been observed (Rosenbaum, 1959). The width of the zone varies


under different cooling conditions, e.g. 0.5 μm after cooling in water, 2.0 μm after cooling in ethanol, 10.0 μm after cooling in air, and 15 μm after cooling in sands. It means preparation processing could affect diffusion, precipitation at grain boundary, so as to generate precipitate denuded zones with different width. Reference (Rice, 1966) shows the ZrO2 rich region in MgO ceramic. From the fracture surface it shows that grain boundary is around 0.5μm in width and the fracture surface might present a two-phase region. Two-phase region could also be found in the surface of as-sintered PTC ceramic with excessive TiO2 doped (1.25 mol%), as shown in Fig.2.19.

Fig.2.18 Optical microphotograph shows the remnant domain at the grain boundary region after removing electric field E, the different color correspond to different stress condition (under tension or under compression) ( See Color Picture in Appendix)

Fig.2.19 SEM micrograph of as-fired surface of PTC ceramics with excessive Ti (Ba0.55Pb0.45) 0.997La0.003TiO3 +1.25 mol%TiO2


As above experiments demonstrate, although grain boundary is thin in width, it could affect a grain boundary region of larger distance, especially when there are impurities or second phase. Besides, grain boundaries could be categorized into those in tensile stress and those in compressive stress, and grain boundaries perpendicular or parallel to applied electric field are different.

2. 4. 5 “Core-shell” structure “Core-shell” structure of grains will be discussed as following. Ideally, grains in polycrystalline materials should have uniform structure similar to that of single crystals. However, researches in recent years have revealed that shell layer of grains near grain boundary has different composition and properties from those of core part of grains. Kahn (1969) proposed in 1969 the dual structure model within grains. He prepared BaTiO3 with mixed oxide method and then sintered after doping. It has been found that dopant could diffuse into surface layer of grains. In the case of restraining grain growth (grains grow from 0.6 μm to 1.5 μm), dopant diffuse into shell part of grains, core part without dopant accounts for 50 vol %, and capacitance vs. temperature relation presents bimodal curve; while in the case of no restraining grain growth (grains grow from 0.6 μm to 60 μm), dopant diffuse into almost whole grains, core part without dopant are extremely small, and capacitance vs. temperature relation does not present bimodal curve. For a ceramic named K-8500 (BaTiO3 ceramic doped with CaZrO3 and a little SiO2), grain size is 7 μm, with a core part of mainly BaTiO3 of 2.7 μm accounting for 6 % of entire volume, and a shell part containing Ba, Ti, Ca, Zr and Si elements. Rawal (1981) investigated BaTiO3 ceramic for capacitors (with 10%~12% Bi4Ti3O13). A brighter shell layer in microstructure is cubic BaTiO3 with higher concentration of bismuth and Nb, while darker core part is tetragonal BaTiO3 with higher concentration of barium and no impurity. At higher sintering temperature, bismuth could diffuse into core part. M McCarteny (1981) investigated material for multilayer capacitors (BaTiO3 + Bi + Zr). Domain structures were found within grain cores, and shell layer with no domains is 0.2 μm in thickness (with corresponding grain sizes of 1~4 μm) or could even extend to half grain size. Besides, Zr and Bi concentration in shell layer is higher, which could reach several percentages, even higher at the place near grain boundary. At triple junctions, besides Zr and Bi elements, Si and Al concentration is also increased to form crystallized second phase with size of 80~100 nm, or even to grow into a small grain with size of 0.2~0.5 μm. In the investigation on barrier layer capacitors, D L Johnson (1977) observed duplex structure after etching in the material of BaTiO3+1 mol%Al2O3+ 0.65 mol% CuO+0.2mol%La2O3 sintered in


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10% H2+90%N2 at 1425 for 4 hours and then processed at 1200 for 1 hour in air. Due to high corrosion resistance, the raised zones correspond to duplex region. Width of duplex region is related to the re-oxidation process, i.e. larger width at higher processing temperature, which is explained as segregation related to re-oxidation. In the optical observation on PLZT slices, we have also observed shell-core structure, sometimes with domains in core part and no domains in shell layer, which means that cubic and tetragonal phases coexist in one grain, as shown in Fig.2.20. As above mentioned, cathode luminescence microscopy could be utilized to observe concentration gradation of additives within grains.

Fig.2.20 Core and shell structures of PLZT ceramics (transmited light, cross Nicol with gypsum plate) ( See Color Picture in Appendix)

Concluding above factors, ceramics could be usually regarded as non-equilibrant system and grain boundary migration is predominant during sintering, so as to lead to differentiation between shell and core part in grains. During grain growth, grain boundary migration could also generate different optical properties, electric properties (ε-T relation) and corrosion resistance in the region it swept through. Actually the un-uniform core-shell structure of grains could be utilized to prepare materials of flat ε-T relation, which could be used as stabilizer capacitors. Nakata (1985) investigated ferrite materials with high resolution TEM observation, and found many micro zones of 5~10 nm with slightly differentiation in optical axis orientation (0.1 ~ 0.4 ), which means that raw materials and processing would lead to in-homogeneity within grains and core-shell structure is just an extreme example.

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2. 5 Grain Boundary Migration

2. 5 Grain Boundary Migration During ceramic fabrication,densification processes compete with coarsening processes to determine the path of microstructural evolution.Grain growth is a key coarsening process.Here the grain boundary migration mechanism in PLZT ceramics,and the role of liquid phases in grain growth and microstructural evolution are presented.

2. 5. 1 Generalization During sintering of ceramics, pores are gradually removed and grain boundaries are then formed, while grain boundaries in turn act as vacancy sink to further eliminate pores to finally reach full densification. To reduce energy of the whole system, grain boundaries tend to migrate towards their curvature centers to decrease overall interfaces. In intermediate stage and later period of sintering, a large amount of grain boundary migration occurs. Densification of hot-pressed ceramics is related to grain boundary migration (Bazarova, 1979), which is also very important in ceramics prepared by other methods including forging, casting and CVD. In order to control ceramic microstructure, grain boundary migration needs to be understood. Grains with more than six sides tend to have concave sides while grains with less than six sides tend to have convex sides, That is grain boundary tend to migrate toward centers of curvature (Burke, 1968). Grain boundary mobility is related to grain orientation. Due to low atom step density, grain boundaries parallel to small Miller Indices crystallographic planes have low mobility (Gleiter, 1969). In addition, impurity ions (species and concentrations), solute ions, pores, second and liquid phase, and free energy contained in matrix could all affect grain boundary migration. Since free enthalpy reduction caused by reaction energy produced in multi-composition materials is much larger than that caused by grain boundary migration, grain boundary migration and grain growth would be restrained in these systems (Kools, 1985). In one word, grain boundary mobility varies greatly, e.g. reaching 104, under different conditions (purity, density of materials, and existence of second phase). For example, by doping of 500 10-4 % Sn into metal lead, grain boundary migration speed will increase for 5000 times to bring rapid grain boundary migration in local region, so as to lead to abnormal growth. Yoon (2003) has proved that grain boundary migration varies greatly in BaTiO3 ceramics with or without liquid phase. In material with inhomogeneous liquid phase, abnormal grain growth would lead to microstructure with most small grains of 1 μm and a few large grains of 100 μm; while in material without liquid phase, homogeneous microstructure with grain size of about 0.5 μm could be obtained.

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The ε-T curve of former material presents a sharp Curie peak, while the ε-T curve of later material is quite flat. Thus two materials have totally different applications and later could be used for stabilizer capacitor. Abnormal grain growth would bring mechanical strength decrease of 30%~50% to Al2O3 ceramics (Dorre, et al, 1984). It would also bring electric strength decrease to PTC materials and make them more vulnerable to micro cracks and fracture, which might lead to substantial loss. Actually grain boundary migration was first observed in metals. Grain boundary migration at around 150 was observed in tin alloy in 1920 (Carpenter, 1920), and Burke (1949) reported growth and disappear of zinc grains through grain boundary migration.

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2. 5. 2 Centripetal and acentric grain boundary migration When materials are under plastic deformation, two kinds of grain growth would occur in annealing process: first and second re-crystallization. When first re-crystallization occurs in matrix under plastic deformation, crystal nucleus free of strain could form and grow, and matrix grains are consumed at the same time. The driving force for the process is the energy difference between phases with and without strain (2.1~4.186 J/g), which is about 100 times higher than interface energy. During first re-crystallization, grain boundaries migrate away from curvature centers to increase volume of low-energy nucleus. Due to little plastic deformation in ceramics during processing, first re-crystallization is quite rare in ceramics but universal for metals. This kind of grain boundary migration away from curvature centers has been observed in NaCl, MgO and PLZT ceramics. Fig.2.21 and Fig.2.22 show the grain boundary migration away from curvature centers during thermal etching.

Fig.2.21 Grain boundary migrations away from curvature centers during thermal etching 1

üOriginal location; 2üLocation after migration


Fig.2.22 Grain boundary migration during thermal etching

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(a) 1150 /10min; (b) 1260 /10min O Original location; M Location after thermal etching at 1150 F Location after thermal etching at 1260

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Second re-crystallization in polycrystalline materials is realized by growth of individual grains and consumption of fine grains with no strains, as shown in Fig.2.23. The driving force is the energy difference between fine and coarse grains to finally decrease overall interfaces. Grain boundary migration toward curvature centers could obtain larger single grains, with energy increase from 0.4186 J/g to 2.1 J/g if a grain grew from 1μm to 1cm. Second re-crystallization is quite common in ceramics, e.g. Al 2O3 , PZT and ferrite ceramics.


Fig.2.23 Abnormal grain growth of PLZT ceramics during sintering, G.B. migrate toward the centers of curvature

Fig.2.24 Different stress state of atoms located on different side of curved grain boundary

ü ü

(Atom A under compression; Atom B under tension.)

There’s a chemical potential difference on two sides of curved grain boundary due to different stress states, as shown in Fig.2.24. Atom A at the convex side of grain boundary is under compressive stress with high energy, while Atom B at the concave side of grain boundary is under tensile stress with low energy, so atom A tends to migrate toward B due to energy difference, i.e. grain boundary migrating toward the curvature center, which leads to grain growth. The migration speed is proportional to curvature of grain boundary.


At equilibrium, all grain boundary energies are the same with grain boundary angle of 120°. In planar drawing, only hexagonal grains have grain boundary angle of 120°, grains with less than six sides tend to diminish, and grains with more than six sides tend to grow. The driving force for grain growth as well as grain growth speed is inversely proportional to grain sizes, which will be further explained with some examples. In investigation on pure aluminum, P A Beck et al (1950) have demonstrated: 1. Grain boundary migration induced by surface energy is toward curvature centers to result in reduction of interface area in unit volume, i.e. grain growth, so as to decrease overall energy of the system, as shown in his figure (Beck, 1950); 2. Grain boundary migration of re-crystallization driven by energy restored in materials during cold working is away from curvature centers to form annealed grains without strains. M F Ashby (1967) introduced deformation energy into copper with weight bearing, and then annealed the material at a certain temperature. Interface migration away from curvature centers driven by stored deformation energy was observed. Originally, particles with average size of about 200 nm were uniformly distributed. When migrating, interface could drag small particles so that denuded zones could be formed after sweeping of interface migration. The interface has a width of several micron, but its migration speed could reach 1 μm/min at 650 , with a dragging force of 10-7 N, 109 times larger than the gravity of one particle. C S Tedmon (1970) observed a small amount of γ phase in stainless steel matrix of α phase. At certain temperature, α-γ grain boundary (phase boundary) would migrate to expand γ phase. But sometimes precipitates at grain boundary could retard or pin its migration. In the investigation on thermal etching on PLZT ceramics, we have observed grain boundary migration away from curvature centers, as shown in Fig.2.25~Fig.2.28.

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Fig.2.25 TEM micrograph of smooth region formed by migration of grain boundary in PLZT ceramic 1

üOriginal position; 2üNew position


Fig.2.26 Scratches on polished surface of PLZT ceramic are swept away by moving G.B.

ü

ü

(a) TEM, 1 Original position; 2 New position; (b), (c), (d) Metallographic micrograph

Regions swept by grain boundary migration became smooth region with original scratches on polished surface swept away. Fig.2.27 shows the ordinal migration of triple junction, original angle of 116 , 137 and 107 became 115 , 126 and 119 , closer to equilibrium position of 120 . Lozinskii observed similar phenomena in the investigation on grain boundary migration in metals (Lozinskii, 1961). Fig.2.28 and Fig.2.29 show grain boundary migration retarded or pinned by precipitates. Sometimes precipitates remain at original grain boundary while grain

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Fig.2.27 Grain boundary migration of triple junction in order 1

üOriginal position, 2üNew position

Fig.2.28 Pinning effect of precipitates on grain boundary migration (TEM) 1

üOriginal; 2üFinal; 3üPrecipitate

boundary has moved to a new position, as shown in Fig.2.30. Besides, distorted grain boundaries were observed near surface of sample after thermal etching, as shown in the right part of Fig.2.31, which means grain boundary migration for different segment differs. However no distorted grain boundaries were observed in the center of the same sample, as shown in the left part of Fig.2.31. Fig.2.32 shows undulating grain boundary, which might result from stress relaxation during thermal etching. Wersing held similar viewpoint concerning this phenomena (Wersing, 1980).


Fig.2.29 Pinning effect of precipitates on grain boundary migration (metallographic micrograph)


Fig.2.30 Precipitates remain at the original grain boundary (a) TEM micrograph; (b) Metallographic micrograph 1 Original position; 2 New position; 3 Precipitate

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ü

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Grain boundary migration in PLZT ceramic during thermal etching might result from PbO acquiring or losing from atmosphere, i.e. chemically induced migration, characterized with migration away from curvature centers, which is common in metals. Segregation of some solute ions at grain boundary could play a stabilizing role on new grain boundary.


Fig.2.31 Straight (inner, left) and distorted (outer, right) grain boundary after thermal etching (metallographic micrograph)

Fig.2.32 Different sections of grain boundary, moving in opposite directions, to produce undulating boundary (TEM)

In investigations on metals, it has been found that grain boundary provides a pathway for rapid diffusion of solute ions, thus grain boundary is sometimes called as short circuit path. In addition, hydrogen element tends to segregate at grain boundary and a high speed diffusion current of hydrogen often occur at grain boundary, which could induce grain boundary migration. These migrating boundaries could eliminate some dislocations so that there are very few etch pits in the area swept by grain boundary migration (Tseng, 1986).

2. 5. 3 Liquid phase and abnormal grain growth during sintering Abnormal grain growth is usually occurred during ceramic sintering. Abnormal grain growth is caused by liquid phase produced by impurities rich at grain boundary to promote grains dissolving and precipitating. Grains grown from liquid phase often possess similar properties to crystals crystallized from solution,


i.e. grain nucleation and the growth of crystalline planes . Liquid phase plays a role of mass transfer during processes of grain growth, which determine abnormal grain growth in stead of being controlled by otherwise grain boundary migration. In case of two-phase inclusion or pores, grain boundary migration will be retarded to remove pores, so that area swept by grain boundary migration would become denuded zone. The phenomenon has been observed in sintered Y3Fe5O12 ceramics (Kooy, 1962). According to J E Burke et al (1963), grain boundary moves toward curvature center and sweep away pores. In area near pores, lattice vacancy concentration is higher than the equilibrium concentration within grains, so these vacancies would enter grain boundary and be removed, so that pores near grain boundaries could be removed faster than those inside grains. At the same time, grain boundaries act as sink against vacancy diffusion. Few pores in grain boundary region have also been observed in BaTiO3 ceramic doped with Fe(Seuter, 1974). R W Balluffi observed clear regions free of pores symmetrically on both sides of grain boundaries and explained the phenomenon from vacancy diffusion (Balluffi, 1955). Pores (Anon, 1972) and insoluble second phase particles (Coble, Burke, 1963) could impose effects of dragging or pinning on grain boundary migration to slow down its migration speed. J E Burke (1980) estimated that grain boundary mobility in pure Al2 O3 ceramic is 100 times higher than that in Al2O3 ceramic doped with MgO, which means that the doping could effectively restrain grain growth. According to Yan (1984), additive could restrain grain boundary migration when there are large differences in atomic radius or valence between additive and matrix atoms. He also found that 7 mol% La into PbZr0.65Ti0.35O3 ceramic could decrease grain boundary mobility by 250 times, which might be the main reason for PLZT ceramics to be easily sintered as transparent ceramics (Yan, 1977). Some liquid phases could speed up grain dissolving and precipitating of mass near grain boundary to improve mass transfer, thus grain growth with liquid phases is much larger than that under purely solid sintering. During sintering of PTC materials, there’s a small quantity of active liquid phase playing a significant role in promoting processes of dissolving and precipitating. Accompanied by grain growth, grain boundary migration would bring semi-conducting elements included into grains and it is not the case that the doner elements, e.g. niobium, diffused into grains (Drofenik, 1982;Lubitz, 1982). Sometimes grain shell with niobium and grain core without niobium could be obtained. W R Bussem (1971) has proved that diffusion of niobium along grain boundary is much larger than that in grains. M Kahn (1985) has proved that niobium is not very soluble in BaTiO3 ceramic. D Kolar (1981) suggested that thermal cycles could promote motion and homogeneity of grain boundaries. M Drofenik et al (1982) has discovered in


BaTiO3 ceramic doped with antimony that there is liquid phase at triple junctions and second phase (glass phase) at grain boundaries. After oxidation, antimony element dissolves into liquid phase; after re-crystallization, oxides are reduced to obtain semi-conductively doped BaTiO3 ceramic, and release oxygen; finally grain boundary region would be re-oxidized to form potential barriers, then sintering of PTC ceramics is finished. Therefore, the sintering is accompanied by processes of grains dissolving, and grain growth, and release of oxygen, in which liquid phase play a very important role of medium. A small amount of liquid phase (0.x %) could bring tremendous impact on microstructure of PTC ceramic. Thus different additives to form different liquid phase, e.g. Al2O3-SiO2-TiO2 or SiO2-Al2O3-Li2O-Sb2O5, would lead to different microstructure of coarse or fine grains, so as to provide greatly different dielectric properties for PTC ceramics. SiO2 could be added into magnetic material SrFe12O19 to obtain fine grain structure, but the adding quantity needs to be properly controlled. Experiments show that when SiO2 quantity is increased, SiO2 concentration within grains remains low and that in grain boundary region would be increased accordingly. However, the second phase of SiO2 stay at triple junctions in stead of grain boundary. If the second phase is solid, impurities at triple junctions would decrease grain boundary mobility due to dragging effects, so as to restrain grain growth. Therefore, if too little SiO2 added in, no segregation of SiO2 would be formed to restrain grain growth; if too much SiO2 added in, the second phase becomes liquid phase to promote mass transfer and grain growth instead of restraining it (Lubitz, 1981).

2. 6 Relation between Grain Boundary and Properties Influences of grain boundary on material properties could be considered from three aspects: 1. Influence on misfit angle, originates from the change of lattice orientation cross grain boundary; 2. Influence on structures, originates from factors including disordered structures and high strain energy; 3. Segregation of impurities at grain boundary. For materials with domains, i.e. ferromagnetic and ferroelectric materials, properties mainly depend on domain structure which is greatly related to grain boundary. For example, grain boundary could affect the domain width at equilibrium, impose pinning effect on domain wall motion, accommodate domain nucleation, and facilitate domains to run through grains. Mechanical and electric properties will be discussed relevant to grain boundary as following.

2. 6 Relation between Grain Boundary and Properties

2. 6. 1 Influence on mechanical properties After cooling process following sintering of ceramics, grains are exposed to stresses or restrictions, which would result in some physical phenomena and processes. For example, gradual evolution of grain boundary stresses could lead to aging of some properties. Phase transition could be affected by strains, e.g. segregation could be form to reduce strain energy at certain temperature, stresses bring coexistence of some phases, coercive fields is often affected by micro stress at grain boundary in materials with domains, and stresses could also affect properties including domain structures and polarization. Quartzite sandstone, a kind of naturally hot-pressed poly-crystalline ceramic, consists of quartzite grains with different optical axis direction. With colored microscopy, A W Allen (1968) has successfully observed allochroic grain boundary region, which is caused by piezoelectric double refractive effects from grain boundary strain. PLZT ceramics are sensitive to compressive stress. Fig.2.33 shows morphology of PLZT film under crossed polarized light after applying and then removing electric field, with dark phase as optically isotropic grains (α phase) and bright phase as optically anisotropic grains (β phase), which is induced or related to stresses. Fig.2.33 explicitly shows the effect of grain boundary stress on phase transition in PLZT ceramics.

Fig.2.33 Stress induced β phase and remnant domains in grain boundary region (optical, crossed Nicol)

E W Hart (1972) and J J Burke (1969) have suggested that phase transition temperature Ti of grain boundary region could differ from that of grains, since tensile or compressive stresses could raise or lower Ti (Firtsakh, 1980;


Guntersdorfer, 1967). As shown in above Fig.2.18, there are domains in parts of grain boundary regions but no domains within grains, which also means that they have different phase transition temperatures, because grain boundary regions have different band energy and coordination number from those of grains (Hart, 1972). Among electric ceramics, there are two categorizes of materials with domains, i.e. ferromagnetic and ferroelectric materials, whose properties are directly related to domain structures. With polarized light microscopes, domains in PLZT ceramics could be directly observed. When electric field E applied, a large amount of oriented domains instantly appear with most of them originating from grain boundary and extending to all grains. There are similar processes in piezoelectric and ferromagnetic materials, but difficult to directly observe. Relation between domains and grain boundaries in a long range needs further investigation, but present examples of observation will be given out as following. High stress at grain boundary could promote formation of 90º domains in materials (Pohanka, 1976). Most 90º domains in BaTiO3 ceramic originate or nucleate near pores or grain boundary region with high stress. It has been observed 90º domains induced by grain boundary stress to adjust spontaneous deformation so as to reduce overall strains. E T Keve (1975) suggested that domains at grain boundary take on ribbed shapes or run through grain boundary. Between neighboring grains, stress concentration could lead to twin crystal, domain generation and serial domains. Stress produced by twin crystals within one grain is big enough to cross grain boundary and affect twin crystals in neighboring grains to move forward (Devries, 157; Neklyudova, 1964). R C Devires (1957) has show banded structure across grain boundary in BaTiO3 ceramic. H Ihrig (1984) has shown domain continuation across grain boundary or long range domain correlation among different grains. We have observed similar phenomena in PLZT ceramic. Fig.2.34, shows the

Fig 2.34 Domains across the grain boundary (optical crossed Nicol with gypsum plate) ( See Color Picture in Appendix)

ü edomain; G üGrain boundary

D 90


micrograph of PLZT polycrystalline slice (with a thickness of a single grain) under applied electric field. The figure demonstrates penetration of domains through grains, and makes a comparison with the acoustic micrograph in Fig.3.6(f) in chapter three. Due to different etching speed of domains with different orientation and polarization, sample surface will present unevenness after chemical etching. Fig.2.35 shows the micrograph of replica of these samples. It could be seen that some domains could cross grain boundaries but other could not, as shown in Fig.2.36. In TEM investigation on PLZT ceramic, Xiantong Chen (1988) has proved that narrow grain boundaries only have a thickness of 0.4~0.6 nm, i.e. size of a crystal cell. When electric field applied, domain nucleation, growth, extension to the whole grain, instant orientation in many grains of PLZT could be observed, as shown in Fig.2.37 and Fig.2.38 (in colored pages). K Neklyudova (1964) has show the situation of a common point at grain boundary of polycrystalline Fe twin crystal. Dislocations propagate mainly through high stress region, e.g. grain boundary and triple junction, and finally extend to the whole sample, as shown in Fig.2.39 (Hull, 1975). Wang (1983) has successfully


Fig.2.35 TEM morphology of domain walls near grain boundaries in the chemical etched PLZT ceramic

ü

ü

G Grain boundary; W Domain wall

observed the rapid propagation of phase transition though out the whole sample. By utilizing Kerr effect, Fidler investigated nucleation and growth of reversed domain in Co5Sm material, and discovered that precipitates near grain boundary could act as nucleation center, and finger like domains could cross grain boundary (Fidler, 1979). Magnetic domains in metal magnetic materials tend to cross grain boundary due to relatively clearer grain boundary since there is little impurities in metals. Therefore, grain sizes do not play a significant role on magnetic permeability μ in metals. However it’s another story for ceramic magnetic materials. Since grain boundaries in ceramics are relatively thicker, sometimes with pores, impurities and chemical inhomogeneity, they could retard magnetic domain motion, thus it’s difficult to obtain high μ for ceramics. A Goldman (1990) has indicated that μ in ferrite materials is higher with larger grain sizes and correspondingly less grain boundaries. In Mn-Zn ferrite material, magnetic permeability μ is directly proportional to grain size d, i.e. larger d brings higher μ. Due to pinning effect on domain motion, pores and other defects would decrease μ, especially pores within grains. He suggested that thickness and chemical composition of grain boundary was the key factor to determine magnetic properties of ceramics. Actually, factor including pores, micro cracks, precipitates at grain boundary, second phase, and strain could all decrease μ.


Fig.2.36 Domains across and within grain boundaries (a) Across grain boundary; (b) Confined within grain boundary

Fig.2.37 Domain nucleation and growth from grain boundary in PLZT ceramic

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(a) Electric field E =670V/mm; (b) E =1000V/mm (direction of E:


Fig.2.38 (a)Domain orientation between two eletrodes and (b) similar to (a) but with gysum plate (optical, cross Nicol PLZT) ( See Color Picture in Appendix)

Fig.2.39 Propagation of slips from one grain to an adjacent grain


With above examples of ferroelectric and ferromagnetic materials, influence of grain boundary on nucleation and growth of domains, as well as domain propagation passed or retarded by grain boundary have been demonstrated, but effect of grain boundary on domains has not been fully understood. If available, high speed microscopy (quicker than microsecond) would help understand the processes. Grain boundary region has an atomic density around 70% of that of grains. Average grain boundary energy of three oxide crystals is 8.36~12.54 kJ/mol higher than that of grains, and cleavage energy is 50%~70% of grain energy. In addition, grain boundary has a lower elastic modulus, higher elastic yield strength, lower melting point and higher anion diffusion than those in grains (Carnigla, 1966). For example, oxygen diffusion at grain boundary is higher than that in grains in MgO ceramic (Hashimoto, 1972), and oxygen diffusion at grain boundary is 100 times higher than that in grains in Al2O3 ceramic (Oishi, 1960). At high temperature, vacancies are uniformly distributed. When quenched to a low temperature, vacancies would remain in grains while vacancies near surfaces or grain boundaries could be eliminated by sink, which determines hardness in metal materials (Aust, 1966). Internal stresses are universal in ceramic materials. R C Pohanka measured internal stresses in BaTiO3 ceramic, discovering that internal stresses in fine grain materials are relatively low and uniformly distributed, but locally high stress often occur in coarse grain materials (Pohanka, 1978). For materials of asymmetric crystal system, such as Al2O3 and BeO, stresses would be established between adjacent grains due to different contraction along different crystal axis during cooling after sintering. The lower is the temperature; the higher are the stresses, which might cause fracture in coarse grain materials. Actually elastic strain energy is directly proportional to the cube of grain size d, causing much larger stress in materials with larger d. For example, in some mine processing factory or electric porcelain manufacturer, quartzite could be preliminarily smashed by quenching with water after heating due to abruptly increased stress among grains. Pores and coarse grain boundaries are mechanically and electrically weak points in ceramics (Yamashita, 1984). Pure PZT materials tend to fracture along grain boundary, while some additives could diffuse into grain boundary to change intergranular fracture into transgranular fracture, so as to increase bending strength, as shown in Table 2.2. Table 2.2 The effects of grain boundary impurities on strength of ceramic materials

Kinds of Materials

Grain size/ m

Pure PZT

22

Cr2O3

(added)

Cr2O3 (Implantation)

22 3

Bending strength/MPa

̚600 ̚800 1000̚1200 300

600


Due to large internal stress from great anisotropy of crystal cell, PbTiO3 ceramic would sometimes fractures along grain boundary into separated grains after placement of several days, i.e. so called “pulverization in placement”. Impurities could be added to increase strength of grain boundary so that pulverization could be avoided (Matsuo, 1966). During polarization of piezoelectric PZT ceramics, polarization under strong electric field for a long time does not ensure a high electromechanical coupling coefficient Kp, since grain boundary cleavage might occur during cooling. For corundum ceramic with some glass phases at grain boundary, proper heat treatment could be performed to crystallize glass phase so as to improve toughness of material as well as to adjust internal stress states at grain boundary (Zwaniewski,1986). With electric field applied, PLZT ceramic would be transparent under crossed polarized light (on state) due to generation of ferroelectric phase; when electric field removed, material should be lighttight and turn dark (off state), but grain boundaries still present bright part, which means that stresses could affect off state blackness. Junkun Ma (1988) has proved that at temperatures under phase transition temperature, resistance of PTC materials still depend on grain boundary resistance or barrier, which is related to internal stress. Grain boundary potential barriers could be increased by heat treatment due to reduced internal stress. Ma and Kuwabara (1988) proved that microstructure or interlocking of grains could affect PTC effects and aging property. In PTC materials, dielectric strength could be increased by reducing grain sizes since potential barriers lie at grain boundary whose fraction could be increased. Kim (2002) provided following results for samples of the same sizes: Grain sizes /μm

Dielectric strength / V

30

300

15

500

5

700

Since grains and grain boundaries have different elastic properties, elastic waves would be attenuate at grain boundary due to its lower elastic modulus and larger elastic yield strength, thus acoustic micrograph could be obtained with acoustic microscope or scanning acoustic microscope with a resolution of nano-meters (Anon, 1985). Similar to optical and electric microscopy, much work could be performed on acoustic microscopy to expose and investigate mechanical properties of materials, which will make material research more comprehensive.

2. 6. 2 Influence on electric properties K. Okazaki (1981) has found that there are internal bias fields in polarized ferroelectric ceramics. In ferroelectric phase range, i.e. at temperature below TC, grain boundaries bear almost all applied electric field, so grain boundary need to


be improved to increase dielectric strength for materials. He took PTC materials for experiments: at temperatures below TC, material with some second phase at grain boundary has a high dielectric strength, and material with no segregation at grain boundary has a low dielectric strength; but at temperatures above TC, these two materials have close dielectric strengths. Therefore, proper second phase segregation at grain boundary is lucrative for improving dielectric and mechanical strength. Grains and grain boundaries could have greatly different electric properties. In material of SrTiO3 BL capacitor, grains are ohmic resistance with little temperature dependence, but grain boundaries have non-ohmic resistance with resistance decreasing with temperature rise (Park, 1981). Resistances of grains and grain boundaries in PTC materials also differ, reference (Basu, 1987) lists resistance of grains, grain boundaries, and ceramic body of Sb doped BaTiO3 material under different sintering conditions. Grain boundary resistance would be increased by prolonging soaking time. With soaking time of 600 minutes, grains did not grow further but grain boundary phase grew thicker with less barium and more titanium. Interface polarization of grain boundary could enhance capacitivity for 100 times in SrTiO3 BL capacitor (Burn, 1982). High temperature loss of a kind of silicon nitride ceramic is related to viscosity of grain boundary phase, which could be utilized to control high temperature dielectric property by adjusting glass phase at grain boundary (Clarke, 1982). Grain boundary region in ZnO varistor material is electron deficient region (Clarke, 1981). A kind of Be doped SiC ceramic possesses both high thermal conductivity and high electric insulation, with insulation mainly from depletion layer at grain boundary. Yan has proved that most Be element is distributed at grain boundary, and grains and grain boundaries have totally different I-V characteristics. Properties of material often go through with aging, i.e. degradation or deterioration with loading time, temperature, or electric field, which is also related to microstructures. Stress is a significant reason for aging, especially for materials with domain structures. For example, PTC BaTiO3 ceramic consists of oxidized grain boundaries and reduced grains, whose difference in conductivity could reach several orders of magnitude. There are a large amount of anion vacancies within grains but few at grain boundary. Vacancy diffusion will be accelerated under electric field to decrease vacancy concentration gradation. In addition, electrostriction stress would impose much larger compressive stress on grains than that on boundaries, and stress gradation would drive vacancies to diffuse into grain boundary to increase electric conductivity so that local conductivity is increased and likely cause final break-down (Schulze, 1980). Aging under alternating current in PbSrTiO3-Bi2O3-3TiO2 ceramic for high


voltage capacitors is related to damage of grain boundary barriers. After a long time of operation under alternating current, obstructive effects of grain boundaries on carriers would be gradually deteriorated with a large amount of carriers produced in grains, leading to current uprush to cause break down (Jia, 1987). In the degradation tests on ZnO varistor, conductivity will gradually increase as time passes under similar conditions of operation with applied field at 200 . Ionic movement at grain boundaries could affect height of grain boundary barriers, thickness of depletion layer, defect concentration and distribution. In addition, degradation is related to migration of Zn atoms at grain boundary (Philip, 1982). Yao Yao has observed different morphologies of grain boundaries in ZnO ceramic before and after electric field is applied. Fig.2.40(a) shows polished surface of 0.5 mol% Bi2O3 doped ZnO ceramic with some pores, and image of electron beam induced current (EBIC) has been observed at grain boundary when applied with electric field, as shown in Fig.2.40(b). It’s suggested that the heights of grain boundary barriers are the same with no current passing under electric field before aging, but asymmetric distribution might occur to grain boundary barriers after aging and to cause current, with the dark and bright phases related to electric field direction. There is volatilization and diffusion of Pb element during sintering of PZT ceramics. When doped into PZT ceramic, chromium element will concentrate at grain boundary to change frequency-temperature characteristics of the material (Nezezchleb 1980). Disorder and visco-elasticity of grain boundary could help relax stresses among grains. Scattering of elastic waves by grain boundary could explain low Q property of some PZT materials. A major reason for PZT aging lies in space charge field, which is related to varieties and quantities of impurities contained in materials. Under applied electric field, grain boundaries will be under tensile stress while grains under compressive stress, so that vacancies will further diffuse. For some La doped PZT materials, with decreasing of grain sizes, average internal stress distribution will be increased with phase transition temperature decreased and planar ε peaks (Burggraaf, 1975). Yamashita (1982) applied bias voltage on BaTiO3 ceramics of various grain sizes and discovered that these ceramics have different ε, furthermore grains and grain boundaries also have different ε. The internal electric field Eint will alter proportionally with the applied electric field Eext according to following equation: Eint ∝ Eext (ε g / ε g.b )

ć

Where εg is dielectric constant of grains and εg.b is that of grain boundaries. It could be found from above equation that when εg ≠ εg.b, Eint ≠ Eext, and breakdown will occur if Eint exceed a certain value. In investigation on BaTiO3 ceramics of fine and coarse grains, H. Heydrich (1981) has indicated that, for the relationship between spontaneous polarization Ps


and temperature T, large variation of dPs/dT exists in coarse grain (10 μm) ceramic, but no such variation in fine grain (0.3 μm) ceramic; ε –T curves for coarse grain ceramic presents peaks, but no peaks on that of fine grain ceramic. If grain boundary is 20 nm in thickness, grain boundaries in coarse grain ceramics account for 1 vol% while grain boundaries in fine grain ceramics account for 30 vol%, which suggests that these differences might result from grain boundary effects.

Fig.2.40 Effects of electric field on grain boundary of ZnO ceramics (a) Polished surface of ZnO varistor; (b) Polished surface of ZnO varistor under applied field

It was found that ferroelectric domains could still be observed in PTC BaTiO3 material above phase transition temperature TC, and attribute it to stress (Heanstra, 1977). Fig.2.28 shows domains in grain boundary regions but no domains within


grains, together with observed grain boundary stresses. Miller has suggested that ferroelectric phase above TC is caused by internal electric field that stabilizes tetragonal phase at higher temperature. Murakami has observed spontaneous polarization and piezoelectricity above TC and related it to impurities at grain boundary. In one word, ferroelectric phase is likely to exist at the special region of grain boundary. Grain boundary could also affect thermal conductivity. For example, yttrium could be doped into AlN ceramic to form liquid phase at grain boundary during sintering so as to increase its thermal conductivity by several times. Thermal shock resistance of material is closely related to grain boundary. For example, due to higher resistivity of grain boundaries than that of grains in PTC materials, under electric field, temperature of grain boundaries would be increased rapidly with a speed of 106 K/s, higher than that of grains, so as to produce a temperature difference up to 50 . Thus every grain boundary becomes a heat point with every grain as a sink. The stresses are substantial at the moment of applying electric field because of the abrupt temperature rise. Especially when materials are under temperature around TC, thermal fracture might occur due to the negative thermal expansion. With rapid advancements of superconducting materials recently, grain boundary is found to have an explicit impact on its continuity and conductivity of structures (Bender, 1987). A.Tampieri et al has investigated hot-pressing preparation of superconducting ceramic Bi1.84Pb0.34Sr1.91Ca2.03Cu3.06Ox, also called 2223 material for short, and found that proper hot-pressing conditions could increase densification and grain orientation, so as to promote connection and coupling of properties among grains and to increase threshold transfer current Jct. Sometimes improper presynthesis temperature and sintering conditions would produce non-superconducting phase at grain boundary, e.g. 2212 phase, which envelops grains with barriers to weaken connection among grains and decrease Jct (Tampieri, 1999; Duran, Tartaj, 1991). In one word, clean grain boundary is desirable in preparation of these materials.

ć

2. 7 Summary Grain boundary phenomena are quite complicated and could be concluded as following: 1. Grain boundary segregation often exists in ceramics, and segregation concentration at grain boundary could be 100~1000 times higher than that in grains. By adjusting segregation, material properties could be improved or new materials could be developed; 2. Disordered and open structure as well as possible liquid phase makes the mass

2. 7 Summary

diffusion in grain boundary region more remarkable than that in grains. Functions of grain boundary in mass transfer during sintering are similar to those of roads in city transportation. Oxygen tends to be removed during diffusion in ABO3 materials, thus sintering in oxygen could be utilized to obtain ceramics with high relative density. Segregation and diffusion could be utilized to design sintering processing for transparent ceramics, e.g. corundum, yttria, and PLZT ceramics. 3. Some grain boundaries are different in continuity and coherence. Electric field could be applied to alter its coherence in order to change properties of ceramics, e.g. elimination of PTC effects. 4. Tendency of oxidation and potential barriers of grain boundaries, as well as interaction between potential barrier and ferroelectric spontaneous polarization would produce PTC effects, which have been utilized for self-regulated heating elements and over-current protection. 5. During cooling process after sintering, grain boundaries are often under tensile or compressive stresses in ceramics of asymmetric crystal system. Grain boundary stresses and property differentiation between grain boundaries and grains (e.g. dielectric constant, vacancy concentration, electric conductivity, and stress states) are often roots for aging of materials. 6. Disorder and free space provide grain boundary with viscoelasticity so that grain boundaries could accommodate strain and stress and serve as source and sink for vacancy in sintering processes. The property of high energy often makes grain boundaries origins for nucleation or domain formation, and affects many processes (e.g. diffusion, segregation, phase transition, domain formation, aging and fracture) related to energy. 7. Grain boundaries often act as captive centers, which accumulate space charge. When applied with electric field, grain boundaries perpendicular or parallel to applied field often has different properties, or two sides of grain boundary could have different compositions and properties. 8. To reduce energy of the whole system, grain boundaries tend to migrate towards their curvature centers to decrease overall interfaces, while acentric grain boundary migration often occurs to increase volume of low-energy phases. Grain boundary migrations often leave areas with low porosity or no pores. Grain boundary mobility varies greatly, with difference up to 104, under different conditions including impurity and pores,. Therefore doping could be utilized to alter grain boundary mobility to prepare transparent ceramics, e.g. La doped PZT ceramics. Sintering and abnormal grain growth is also directly related to grain boundary migration. 9. Most ceramics have not reached equilibrium after sintering, thus some grains are not ideally single-phase. Shell part near grain boundary and core part of grains have different properties, i.e. shell-core structure. Grain boundaries often

Ā ā

Ā

ā


have a thickness covering several atoms, while grain boundary region could reach 0.x ~ x μm with different properties from grain inside. Proportion of shell and core parts as well as volume fraction of grain boundary region could also affect overall properties of ceramics. In conclusion, in-depth understanding of grain boundary could help control material properties for better development and application of materials.

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Yan M F, Rhodes W W (1982) Preparation and properties of TiO2 varistors. Appl Phys Letters 40(6):536–537 Yoon S H, Lee J H, Kim D Y, et al (2003) Effect of the liquid-phase characteristic on the microstructures and dielectric properties of donor-(niobium) and acceptor-(magnesium) doped barium titanate. J Amer Ceram Soc 86(1):88–92 Zhu B H, Zhao M Y (1996) Recent Trends of Function Ceramics. Physics in Chinese 25(12):718–724 Znidarsic A, Drofenik (1999) High-resistivity grain boundaries in CaO-doped MnZn ferrites for high-frequency power application. J Amer Ceram Soc 82(2):359–365

Near-field Acoustic Microscopy of Functional Ceramics

Two types of near-field acoustic microscopies, scanning electron acoustic microscopy (SEAM) and scanning probe acoustic microscopy (SPAM) , are presented in terms of their operation principle, imaging contrast mechanism and their applications to characterize microstructures of functional ceramics as well as other materials including structure ceramics, metal, single crystals, composites and coatings, etc. Due to their unique features, SEAM and SPAM provide a powerful tool to visualize the buried structures in a variety of materials.

3. 1

Introduction

Scanning electron microscope (SEM) has been an important tool for visualizing material microstructures since it was commercialized in 1830’s. However, SEM could only be used to observe the specimen surface due to its imaging mechanism. In 1970s, scanning acoustic microscope (SAM) was firstly developed by R. Lemons and C. F. Quate (1974) and employed to visualize interior microstructures of the sample non-destructively. But SAM imaging resolution is less than one micrometer scale due to its limitations to the acoustic wavelength. Atomic force microscope (AFM), an ultrahigh resolution imaging technique for a variety of materials, is still limited to characterizing the sample surface morphology. Hence it is necessary to develop a high resolution acoustic microscopy for in-situ exploring subsurface microstructure of material and devices. In this chapter, two acoustic microcopies including scanning electron acoustic microcopy (SEAM) and scanning probe acoustic microscopy (SPAM), are introduced and discussed in details. SEAM combines SEM’s high resolution with sub-surface imaging capability of SAM, while AFM-based SPAM could be

3.10 Temperature Dependence of Ferroelastic Domains in PMN-PT Single Crystals

utilized to image ferroelectric domains and the sample’s interior defects. The two acoustic techniques are also referred to as near-field acoustic microscope since their spatial resolution is not correlated with their excited acoustic wavelength. Here the operation principle of the two acoustic microscopes, their imaging mechanism and their applications to functional materials are presented in details, especially for observing domain configurations of polar functional ceramics. In addition, their applications to a variety of materials including structure ceramics, metal, and composites are also described.

3.2 History and Development of Scanning Electron Acoustic Microscopy As a novel microscopy for subsurface imaging technique, scanning electron acoustic microscopy provides us a practical, convenient and reliable tool for effectively observing and investigating material microstructures and properties. SEAM was firstly proposed by G S Cargill (1980), and E Brandis and A Rosencwaig (1980), respectively. Similar to the thermal wave principle of the optical-acoustic microscopy (OAM), SEAM detects microstructure change in thermal elasticity with a coupling piezoelectric sensor while the specimen is subjected to the modulated energy beam. Therefore SEAM and OAM were generalized as thermal wave imaging technique. However, other theories except thermal elasticity were also proposed to interpret electron-acoustic imaging contrast mechanism (Balk, 1980). Combining the electron optics, weak signal detection, high-sensitivity piezoelectric sensor, image processing with computer technology, SEAM has unique features in non-destructive imaging of surface and sub-surface structures, which make SEAM look increasingly attractive in material fields of many countries including United States (Brandis, Rosencwaig, 1980; Kirkendall, Remmel, 1984), United Kingdom (Cargill, 1980; Davies, Howie, Staveley, 1982), Germany (Balk, Kultscher, 1983; Balk, Kultscher, 1984), Russian (Averkiev, 1984), France (Bresse, Papadopoulo, 1987), Japan (Ikoma, Murayama, Mizuno, 1984; Takenoshita, Managaki, Mizuno, 1984), Spain (Urchulutegui, Piqueras, Liopis, 1989; Urohulutegui, 1993), etc., and image successfully a variety of materials such as metals, semiconductors, and structural ceramics, inorganic functional materials (Yin, Jiang, Hui , 1994; Zhang, Jiang, Yin, 1995; Zhang, Jiang, Hui, et al, 1996), biomaterials (Balk, 1994) and superconductor (Piqueras, 1994) as well. Especially in China, the scientists from Shanghai Institute of Ceramics, Chinese Academy of Sciences (SICCAS), Tongji University, Nanjing University, Suzhou University, etc. contributed greatly to the imaging techniques, theory and applications for SEAM. Prof. Qingrui Yin, a renowned scientist from SICCAS, conducted an extensive study in SEAM ranging from instrument development,

3 Near-field Acoustic Microscopy of Functional Ceramics

applications and electron-acoustic imaging theory since late 1980s (Yin, 1985; Yin, Wang, Qiang, 1991). The first SEAM in China was successfully developed in 1989 (Yin, Tang, Zhang, et al, 1990), since then electron acoustic imaging of various materials and the corresponding theory study were performed (Jiang, Yin, 1991; Zhang, Jiang, Yang, et al, 1996; Zhang, Jiang, Yin, 1996). Prof. Menglu Qian, a scientist from Tongji University in Shanghai, proposed three-dimension thermal wave coupling theory for SEAM imaging (Qian, 1989; Qian 1995). These outstanding results have gained international recognition, and laid a solid foundation for SEAM research in China.

3. 3

Physical Principle of SEAM Imaging

Electron bombardment of a sample is unique to microprobe analysis and produces a large number of effects from the target material (Fig.3.1). The incident electrons interact with specimen atoms are significantly scattered by them. Most of the energy of an electron beam will eventually end up heating the sample. For SEAM, it detects the acoustic waves transmitted from the bottom of the specimens. According to the thermal wave coupling mechanism, SEAM imaging process could be described as below. When a modulated electron beam incidents a specimen surface, electrons interact with the specimen within a certain distance away from the sample surface; while about 40% ~ 80% of electron energy is transformed into heat. Theoretical analysis showed that the volume of the electrons lose energy is nearly a sphere (Murata, 1971). Since electron energy has been modulated with certain frequency, the spherical interaction volume is equivalent to an instantaneous hot spot. Its temperature varies with the same frequency as the modulated electron energy, thus like a thermal wave. For an indefinitely large solid body, the equation for the thermal wave propagation could be expressed as (Carslaw, Jaeger, 1973): T ( x, t ) ∝ Ae −kx + Be kx (3.1) k=

dt =

1+ i dt

2K λ = th ωρ C 2π

(3.2) (3.3)

Where A and B are constants, ω (=2πf ) is the angular frequency, ρ is the specimen density, C is the specific thermal capacity, dt is the thermal diffusion length, and λth is the thermal wavelength. From Equation 3.1, thermal wave in the bulk sample is a type of highly-damping wave, and its propagation distance (about several micrometers) depends on materials. The thermal wave intensity is related to the material thermal property in


the volume of the hot spot (see Equation 3.2 and Equation 3.3). As a result, the propagating thermal waves could be reflected or scattered due to the thermal inhomogeneity like defects and micro-cracks in the sample.

Fig.3.1

Various information generated from electron interaction on specimen

Clearly, periodical expansion and contraction, due to the periodically thermal waves, would result in periodically elastic deformation with the same modulated frequency, which generates acoustic waves carrying the sample’s thermal properties. The acoustic wave usually could propagate to a long distance because of its wavelength up to several millimeters. Thus it was incorporated into a piezoelectric detector coupling well with the acoustic waves inside the specimen to obtain the acoustic image reflecting the sample’s thermal-elastic property. Fig.3.2 shows a schematic diagram of the physical principle of SEAM. It is worthwhile to note that the finally resulting image is the acoustic signal, not the general acoustic image. Due to the long wavelength, acoustic waves only act as medium carrying thermal wave information without interacting with microstructures or defects in the specimen. So that thermal wave variations from thermal inhomogeneity in the specimen lead to the acoustic imaging contrasts. However, situation would become more complicated if defect sizes are comparable to the acoustic wavelength. In fact, such case is quite rare experimentally.


Fig.3.2

Schematic diagram of the physical principle of SEAM

3. 4 Scanning Electron Acoustic Microscopy Image Processing System A SEAM system shown in Fig.3.3 consists of six parts: 1. Electron beam gun. 2. Electron beam blanking. 3. Electron acoustic information processor. 4. Specimen and detector module. 5. Monitor. 6. Computer controller.

Fig.3.3

Computer-controlled SEAM experimental setup


In SEAM, a pair of deflection plates is set in front of the first electromagnetic lens of normal SEM. The deflection plates are controlled by impulse power amplifier so that the periodically modulated signal at the deflecting plates would bring periodic energy variation to electron beam. A PZT piezoelectric detector underneath the specimen is used to receive electron-acoustic signals. The weak signals are processed with pre-amplifier and with lock-in amplifier, and then transfered to the monitors or computers for electron-acoustic images. Combined with general functions of SEM, SEAM could be used for displaying both electronacoustic images and secondary electron images. Square waves produced by signal generator will be firstly modulated to 50 kV by a power supply and then loaded on the deflection plates to obtain periodic deflection voltages. Square wave used here is to avoid smearing of beam spot from rise and fall of sine waves, which would interfere with the useful signal and affect the electron-acoustic imaging resolution. Thus the square-wave could improve the signal to noise ratio and the imaging quality of SEAM as well. In SEAM, a type of piezoelectric PZT material with high-sensitivity, wide response frequency range and high sensitive to various excited modes was selected as the electron-acoustic detector. Extra efforts were made to shield stray electrons in the set-up of detector and signal cable connection in order to reduce noise interference. To shield stray electric and magnetic field (i.e. to avoid the interference of high-frequency power supply on high-frequency modulated electron acoustic imaging), the whole system is separately grounded in a completely shielded laboratory. In SEAM, a number of chemical agents including silicon oil, vacuum grease, graphite conductive adhesive and silver conductive adhesive etc. were used for the coupling between the specimen and the detector. Our studies demonstrated that the silver conductive adhesive provides an excellent acoustic-coupling between the specimens and transducers for strong electron-acoustic signals and high-quality electron-acoustic images. Since electron-acoustic signals directly from detectors are quite weak along with noise, weak-signal processing techniques have to be needed. In our experiments, firstly, signals from detectors are put into a pre-amplifier, and then to a lock-in analyzer to remove noise for a high signal-to-noise ratio. A pre-amplifier with low noise (EG&G PARC MODEL 5316A manufactured by American Princeton Co.) was used in SEAM. Since electron beams have been modulated with periodically continual square waves, a lock-in amplifier for detecting periodically continual signals (EG&G PARC MODEL 5302) was used instead of BOXCAR being for periodically impulse signals. The system for controlling and collecting digital electron-acoustic image was also established in our SEAM. A digital scanning generator was utilized to substitute an analogue scanning generator. Thus electron beam controlled by a


computer scans on a sample, and dynamic magnitude control and signal-to-noise improvement were used in the collection of electron-acoustic signals. In addition, a computer-aided imaging system was used to replace CRT monitor of SEM, which not only improves the quality of electron-acoustic images but also facilitate off-line processing of SEAM images. Obscure images in gray scale could be converted into color images for better distinction. Digital storage and output have also avoided a series of trivial works such as film developing and photo preservation. Computer controlled digital scanning generator provides enough time for collecting electron-acoustic signals and real-time improving signal-to-noise ratio. Due to the outstanding capabilities of collecting weak signals, electron beam intensity could be lowered, so that radioactive damage on samples could be reduced, and the application scope of SEAM could be extended to a broad new range. Furthermore, digital scanning generator makes low-frequency electron-acoustic imaging (less than 10 kHz) possible.

3. 5

Theory Studies of Electron-acoustic Imaging

Various theoretic models have been proposed for imaging mechanism and signal generation of SEAM. Thermal wave coupling model was firstly suggested by A Rosencwaig et al (1976). In 1982, one-dimension theoretical calculation with thermal conductive equation and thermal-elastic equation was performed by A Rosencwaig and J Opsal (1982), but provided a general description of electronacoustic signal generation. After that, D J Davies (1983) in 1983, W L J Holstein (1985) in 1985, A Rosencwaig and J Opsal (1986) in 1986 and L D Fevro et al (1986) in 1987 extended the thermal wave coupling theory to a three-dimension model. Favro indicated that scatters in specimen are the key factors changing thermal waves into elastic waves. The scatters could be any defects in ideal crystals. The model could explain the existence of shear and transverse waves, and further resolution difference due to different distance (the distance should be smaller than thermal wavelength) between scatters and incident electron beam. Because most materials have non-zero coefficient of linear expansion, thermal wave coupling model is a universal mechanism for electron-acoustic imagining. From carrier diffusion equation and the relationship between excess carrier concentration and strain similar to thermal elastic equation, N Kultscher and L J Balk (1983) deduced the elastic wave vibration equation of excess carrier mode. It could be found from this equation that strain is in square proportion to local electric field, which means that incident electron beam with f-frequency would generate electron-acoustic signals with 2f-frequency. Excess carrier model was firstly proposed by R G Sterns and G S Kino (1985) in the analysis of optical-


acoustic imaging mechanism, in which experiments confirmed that elastic waves originate from both thermal waves and excess carries. For piezoelectric materials, Balk et al (1988) proposed a possible piezoelectric coupling mechanism. Since no additional mechanism is required for the resulting acoustic waves from electron beams in this case, e.g. thermal waves from thermal wave coupling, acoustic wave resource mainly depends on the energy diffusion zone. For insulating materials, diffusion function D(r, z) determines the distribution of local electric field; for semiconductor materials, diffusion function D(r, z) plays a dominant role. Therefore under this mechanism spatial resolution does not depend on thermal properties and is irrelevant to the frequency of incident electron beam, which is different from that in the case of thermal wave coupling. Since the acoustic waves are directly generated from the incident beams without scatters, thus high resolution could be achieved with this kind of coupling mechanism even at low frequency. Such feature is very beneficial to observe finer structures of ferroelectric domains. Balk et al (1984) used SEAM to observe magnetic domains in Si-Fe alloy, and their studies indicated that magnetic domains could be visualized only in non-linear electron-acoustic images, and linear electron-acoustic images for sample grains. Based on their results, they proposed a magnetostrictive model to interpret SEAM imaging of magnetic domains. Since the magnetic field due to the electron beams is quite low magnetic field and anisotropic, magnetostriction is in square proportion to magnetic field intensity. This model is limited to qualitative analysis because of the complicated three-dimension calculation involved theoretically. These models for electron acoustic imaging are summarized in Table 3.1. Table 3.1

Theoretical model for electron acoustic imaging

Model

One-dimension wave equation

Acoustic generation mechanism

Force

Thermal wave coupling model

∂ 2u ∂ 2u αE ∂ T ( z , t ) = v 2 2 − cl ρ ∂z ∂ t2 ∂z

Periodically modulated thermal wave

∂ T ( z, t ) ∂z

Piezoelectric coupling model

∂ 2u ∂ 2u e ∂ E ( z, t ) = v2 − 2 ∂t ∂ z 2 ρ Cr ∂ z

Periodically electric field

∂ E ( z, t ) ∂z

Excess carrier coupling model

∂ 2u ∂ 2 u ε 0 − α ∂ E 2 ( z, t ) = v2 + 2 ∂z ∂t ∂ z 2 8πρ Cr

Magnetostrictive model

ü

Material parameters α

: thermal

expansion Ecl: elastic module e:

Piezoelectric stress tensor Cr: Stiffness

Periodically excess carriers

∂ E 2 ( z, t ) ∂z

ε 0: Dielectric constant Cr: Stiffness

Periodical magnetic field

ü

Magnetostrictive parameters


3. 6

SEAM Imaging of Ferroic and Other Materials

Scanning electron acoustic microscopy exhibits a powerful tool for observing microstructures in ferroic and other materials.Here it gives SEAM applications to imaging some typical material.

3. 6. 1 SEAM imaging features of ferroelectric domains SEAM imaging of ferroelectric domains was firstly presented in 1995 by Yin’s group in Shanghai Institute of Ceramics, Chinese Academy of Sciences. Since then substantial investigations of functional material and devices were performed by SEAM and many original results has been presented (Zhang, Jiang, Shi, et al, 1997; Yin, Liao, Jiang, et al, 1998). Previous observations of ferroelectric domains were based on their electrical and optical properties, not from mechanical property. According to the electron-acoustic imaging mechanism, the resulting acoustic signals are dependent on the piezoelectric coefficient and the vibration mode of energy dissipation zone. Thus the material’s mechanical properties are closely relevant to the domain imaging contrast. Due to this unique imaging mechanism, SEAM has some outstanding advantages in imaging domain structure over other microscopy techniques: 1. A variety of materials including ferroelectrics, ferroelastics and ferromagnetics, not limited to a certain material, could be measured by SEAM; 2. Non-destructive imaging capability with retaining the original morphology of domains, and enable real-time and dynamic observing of domain structures; 3. Sub-surface imaging ability with observing the buried domain configurations; 4. No special pre-treatment to the specimen preparation, thus trivial processes such as polishing, milling and etching could be neglected to greatly simplify experimental procedures for domain observation; 5. Simultaneous imaging both surface morphology (secondary electron image) and ferroelectric domains (electron acoustic image) to enable more comprehensive investigation on their one-to-one corresponding relationship between domains and microstructures like grains, grain boundaries and defects, etc. 6. Higher-resolution imaging capability for finer domain configurations. As a novel method for domain observation, SEAM has enriched this field so much. Further SEAM studies of ferroelectric domains enclose real-time observation of domain nucleation and growth under the externally applied electric field, dynamic domain behavior in response to the temperature and the externally applied stress, and the underlying origin for the electron-acoustic imaging.


3. 6. 2 Electron-acoustic imaging of ferroelectric materials Scanning acoustic electron microscopy exhibits unique advantage over other techniques in observing subsurface ferroelectric domains.Here SEAM it is used to image domain configurations of some typical ferroelectric ceramics and crystals. 3. 6. 2. 1 Ferroelectric BaTiO3 materials A lot of techniques have been developed for revealing domain structures such as chemical etching (Hooton, Merz, 1955; Kulcsar, 1956; Tennery, Anderson, 1958), electron microscopy (Ikegami, Veda, 1967; Bursill, Lin, 1984; Bihan, 1989), powder pattern decoration (Hatano, 1973), polarized light microscopy (Merz, 1954; Little, 1955), X-ray topography (Authier, 1964), liquid crystal method (Furuhata, 1973) and etc.. Among these methods, domain structure could only be observed after special treatment on specimen surface (especially for ceramics), which could damage the original domain structures, and brings about more inconveniences to domain observations. SEAM has provided an excellent approach to solve the above-mentioned problems. In this chapter, our electron-acoustic imaging study focus on the perovskite-type BaTiO3 ceramics, crystals and layer-structured Bi4Ti3O12 crystals as well. Due to its unique imaging mechanism, no special pre-treatment is required for the sample preparation. As a result, the initial status of specimen surface and internal structure could evade from damage in SEAM imaging process. In addition, the legible images of domain structure and the one-to-one relationship between domain structure and grains could be demonstrated clearly. Therefore, SEAM provides a practical, non-destructive method for observing domains in ferroelectric materials. 3. 6. 2. 2 Ferroelectric BaTiO3 single crystal

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BaTiO3 has the prototype cubic perovskite structure above 120 . Below 120 it transforms successively to three ferroelectric phases: first to 4mm tetragonal, then to mm orthorhombic at about 5 , and finally to a 3m trigonal phase below –90 . The polar axis in the three ferroelectric phase is [001], [011], and [111]. For tetragonal phase, the spontaneous polarization is along c axis, and c axis will be slightly prolonged with tetragonality c/a≈1.01. SEAM imaging of BaTiO3 crystal was usually conducted with 30 keV of electron beam. The time integration constant of lock-in amplifier was set to 100 μs or 200 μs, depending on the electron-acoustic signal intensity. The signal sensitivity is 100 μV. Electron-acoustic imaging experiments with frequency- modulated electron beams were also performed. Meanwhile, non-linear electron- acoustic signal at the double frequency (2f) were measured and found nothing. Fig.3.4(a) shows the secondary electron image (SEI), while Figs.3.4(b), (c), (d) show electron-acoustic images (EAI) under different modulation frequencies,

ć

ć

ć


and Fig.3.4 (e) shows the line-scanning magnitude of the electron acoustic

Fig.3.4 SEI of BaTiO3 single crystal (a), EAI at different modulation frequency of f =145.5 kHz(b), f =153.8 kHz (c), and f =334.1 kHz (d), respectively; the magnitude of electron acoustic signal amplitude scanning along the specimen surface (e); SEI (f) and the corresponding EAI (g) of domain structure in another BaTiO3 single crystal (Courtesy to Kohler, Schubert, 2002)


signal (Zhang, Jiang, Yang, et al, 1996). For further understanding, Figs.3.4(e), (f) demonstrate SEI and its corresponding EAI of ferroelectric domains in another BaTiO3 single crystal (Kohler, Schubert, 2002). Surface morphology could be observed from the SEI, but electron acoustic image provides totally different information from SEI. The EAI in Fig.3.4(b) shows domain configurations inside the crystal, and reflects some surface defects with varying the modulation frequency of electron beam. The alternating black and white strips in Fig.3.5 represent c-domain and a- domain.

Fig.3.5

c-a domain structure of BaTiO3 single crystal

For linear electron acoustic image, excess carrier coupling mechanism is not possible. In addition, the domain morphology tends to be clearer as the modulation frequency decreases while some surface defects appear with increasing the modulation frequency. Such phenomenon indicates that thermal wave coupling plays a certain role in the electron acoustic signal’s generation. At higher frequencies, defects at the sample surface become clearer on the electron acoustic images because the detection depth is closer to the specimen surface. While the domain patterns change blurry with unclear boundary due to the defect’s effect at the specimen surface. In the case of lower frequencies without surface effects, the bulk domain configurations are more resolved with increasing detection depth. From the above discussions, it could be found that the material’s electric parameters play a more significant role on electron acoustic imaging than the thermal parameters. Theory studies indicated that the electric field from spontaneous polarization would affect the electron acoustic imaging contrasts. For ferroelectric BaTiO3 crystal, the incident electrons in SEAM will cause changes to the internal polarization field, the internal stress and the strain in the crystal due to its piezoelectric effect. The deformation finally results in the electron acoustic signal in the piezoelectric detector attached beneath the sample. The spontaneous polarization vector and polarization magnitude are closely related to the phase and amplitude image of the electron-acoustic signal, respectively. Thus, c and a-


domains would be differentiated in the electron acoustic images. 3. 6. 2. 3 Ferroelectric BaTiO3 ceramics Due to its high dielectricity and piezoelectricity, BaTiO3 ceramics has been widely used in electric engineering and other fields such as capacitor, sonar transducer, electro-acoustic transducer and ceramic filter for communication. As a kind of polycrystalline material, BaTiO3 ceramic usually consists of a number of randomly oriented grains with multi-domain state. General techniques for observing domains in the opaque BaTiO3 ceramics are quite complex, mainly involved with polishing, etching, thinning or other processing to a specimen. As mentioned earlier, these processes will damage the original domain state. However, electron acoustic images of domain structure in BaTiO3 ceramic could be obtained from SEAM without any pre-processing, just similar to that of BaTiO3 single crystal. The domain distributions show a good corresponding relationship with individual BaTiO3 grain. Here the electron acoustic imaging mechanisms under different experimental conditions are presented. In the SEAM, the acceleration voltage for the electron beam is 30 kV, the integration time constant is 100 μs, detection sensitivity is 100 μV, and the modulation frequency varies from 60~400 kHz. The SEI image of Fig. 3.6(a) shows grain distribution on the specimen surface, and the corresponding EAI images in Figs.3.6 (b)~(d), measured respectively at 98.9 kHz, 114.7 kHz and 133.7 kHz, reveal various details related to domain structures. Grain distributions in SEI are well correlated with domain configurations in EAI. It can be seen that the domain arrangements at neighboring grains tend to be parallel, and various oriented domains coexist at the junctions of three or more grains. Such domain distribution could keep the whole system more stable with lower free energy. Further findings reveal that the integrity of ferroelectric domains with averaged 2 μm in width are closely relevant to the grain’s sintering quality of BaTiO3 ceramics. Several domain orientations can be found in the EAIs of BaTiO3 ceramics, as indicated by mark “S” in Fig.3.6(b), which suggests that the grain is actually twin crystals rather than a single crystal. Meanwhile, domains are found to cross several grains, as marked by “Q” in Fig.3.6(b). The rectangle and circle areas in Figs. 3.6 (b)~(d) show different domains in the same region at different modulation frequencies. In addition, domains are narrower for smaller grains and vice versa, i.e. domain width is related to the grain size to some extent. Figs. 3.6 (e)~(f) illustrate domain-crossing phenomena through grain boundaries. As a result, several characteristics about domain arrangement could be obtained from SEAM: 1. Domains in the two neighboring grains tend to align in parallel at the grain boundary. 2. Domain alignment tends to make a certain angle at the junction of several grains. 3. Various oriented domain could coexist in one grain.


4. Domain could cross several grains. 5. Domain width is related to grain size. In addition, the polarization and its orientation of BaTiO3 ceramics play a significant role in the electron-acoustic signal’s generation and the imaging contrast. The thermal wave coupling mechanism has also some effects on the grain boundary imaging and the acoustic detection depth (Zhang Jiang, Yang, et al, 1996; Liu, Balk, Zhang, et al, 1998).

Fig.3.6 SEI of BaTiO3 ceramics (a), EAI at different modulation frequencies of f=98.9 kHz (b),114.7 kHz (c) and 133.7 kHz (d), respectively, (e) and (f) showing domain structure crossing the grain boundary of BaTiO3 ceramics ( See Color Picture in Appendix)


3. 6. 3 Ferroelectric Bi4Ti3O12 single crystal Figs.3.7 (a)~(b) show SEI and EAI of Bi4Ti3O12 single crystal, which exhibit greatly different imaging contrast. Fig.3.7(a) only shows the surface morphology and some defects, while Fig.3.7(b) reveals stripped domain distribution with different width and orientation (Jiang, Kojima, Zhang, et al, 1999).

Fig.3.7

SEI (a) and EAI (b) of Bi4Ti3O12 single crystal

3. 6. 4 Ferroelasitc NdP5O6 single crystal Figs.3.8 (a)~(b) show SEI and EAI of NdP5 O6 single crystal, respectively. Fig.3.8 (a) only shows surface morphology and defects, and Fig.3.8 (b)


shows clearly the stripped ferroelastic domain structures. Since NdP5O6 single crystal is not piezoelectric, acoustic signals basically originate from the thermal elastic effects of specimen when the electrons beam incidents on the specimen. Therefore the domain imaging contrast mainly comes from locally elastic property variations of the specimen (Jiang, Kojima, Zhang, et al, 1999).

Fig.3.8

3. 7

SEI (a) and EAI (b) of NdP 5O6 single crystal

Magnetic Domains in Austenitic Steel

Figs.3.9(a) ~ (b) show SEM image (back scattered electron image) and SEAM image of austenitic steel with formation of martensite. The SEM image only shows the surface scratches, while various oriented grain structures are clearly seen in SEAM


image. The electron acoustic contrast mainly comes from the anisotropy of the thermal elastic properties (Kohler, Schubert, 2002).

Fig.3.9 SEM image (a) and the SEAM image (b) of the austenitic steel with formation of martensite (magnification 200 , image size 560 μm 560 μm) (Courtesy to Kohler, Schubert, 2002)

h

h

Fig.3.10 reveals the SEAM image of ferromagnetic domains in the austenitic steel containing martensite. It was the first time that magnetic domain structures in this kind of materials were observed with SEAM technique. The magnetic domains contrasts result from different thermal-elastic properties of different oriented ferromagnetic domains. In addition, when a specimen is subjected to the periodically modulated electron beam, the magnetic field produced by the induced


periodic current would generate magnetostriction, which also affects the magnetic domain imaging contrasts (Kohler, Schubert, 2002).

Fig.3.10 SEAM image of austenitic steel with formation of martensite (magnification 100 , image size 1120 μm 1120 μm) (Courtesy to Kohler, Schubert, 2002)

h

h

3. 8 Modulation Frequency Dependence of SEAM Imaging Domain Structures According to the thermal wave imaging, SEAM spatial resolution depends on the spot size of electron beam, diffusion depth of the incident electron beam and diffusion length of thermal wave. δ = d0 + ds + d r

(3-4)

where δ is spacial resolution, d0 is the spot size of incident electron beam, ds is the diffusion depth of incident electron beam, and dr is the diffusion length of thermal wave. The equation for thermal wave diffusion could be expressed as: di =

2k ωρ C

(3-5)

where k is thermal conductivity, ρ is specimen density, C is thermal capacity of specimen, and ω=2πf with f as the modulation frequency for electron beam. Thus for the electron acoustic images due to the thermal wave coupling mechanism, spatial resolution will be improved with increasing modulation frequency, and this criteria is often used to judge whether electron acoustic imaging contrast originates from thermal wave coupling or not. It could be found in the electron acoustic images (Fig.3.11) of 0.65PMN-0.35PT crystal that the resolution is improved at lower frequency of 130.3 kHz (Fig.3.11(a)), and more


details appear than those at higher frequency (Fig.3.11(b)). Their comparison reveals another feature about SEAM imaging domain, i.e. modulation frequency dependence of different domain groups. Such behaviors are clearly demonstrated in the electron acoustic images of Figs. 3.12 for BaTiO3 crystal at 79.9 kHz, 80.6 kHz and 132.7 kHz, respectively (Jiang, Kojima, Zhang, et al, 1998; Liao, 1999).

Fig.3.11

SEI of the 0.65PMN-0.35PT single crystal at different modulation frequency of 130.3 kHz (a) and 138.3 kHz (b)

For thermal wave coupling mechanism, the electron acoustic signal’s magnitude shows a linear decrease as increasing the modulation frequency. In contrast, electron acoustic images of domain structure in BaTiO3 crystal show distinctive frequency dependence; i.e. quite strong electron acoustic signals and clear EAI contrast could be obtained in some specific frequency range, but very weak in other range. In addition, the bright and dark contrasts could exchanges at the two sides of one peak frequency. Therefore, the EAI contrast of BaTiO3 crystal mainly originates from other mechanisms other than thermal wave coupling. However this does not exclude the thermal wave coupling effect because the inhomogeneity in thermal properties could still result in some contribution through thermal wave coupling, eg. the


imaging contrast of surface scratches and scraps actually comes from the thermal wave coupling. Theoretic analysis on the electron acoustic imaging mechanism of domain structure and some specific phenomena will be given later.

Fig.3.12

Modulation frequency dependence of SEAM imaging domain structure in BaTiO3 single crystal (a) f =79.7 kHz; (b) f =80.6 kHz; (c) f =132.7 kHz; (d) SEI

3. 9

Electric Field Dependence of SEAM Imaging Domains

The domain evolutions in 0.65PMN-0.35PT single crystal were investigated by SEAM under different electric field, as shown in Fig.3.13. It could be found that the domain evolution of PMN-PT crystal differs from the 90 domain movement. The original domain patterns prior to electric field are observed in Fig.3.13(a). Domain group 1 and 2 with different orientations could be clearly seen, and another group of domains with relatively faint contrast is found within the domain group 2. When the crystal is subjected to the electric field of 4.5 kV/cm, domain structures begin to evolve without sidewise growth of 90 domain at both sides (Fig.3.13(b)). The electric field-induced new domain growth emerges in both the bright and dark domains.

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Fig.3.13

EAI of domain evolution in 0.65PMN-0.35PT single crystal at different electric field (f =133.4 kHz)

(a) E=0 kV/cm; (b) E=4.5 kV/cm; (c) E=5.4 kV/cm; (d) E=7.2 kV/cm

When the electric field increases to 5.4 kV/cm, finer strips with remarkable brightness appear in Fig.3.13(c), and the bright strips start to broad at the expanse of the dark strips. Thus domain growth and coalescence are almost complete, leading to the specimen a quasi single domain state, as illustrated in Fig.3.13(d). The single domain state exhibited strong stability even after the electric field was removed (Yin, Zeng, Li, et al, 2003).

3.10 Temperature Dependence of Ferroelastic Domains in PMNPT Single Crystals In order to further reveal the ferroelectricity of the parallel, stripped domain structure with distinct verges, slowly annealing procedure was performed to the PMN-PT crystal. And SEAM was used to study ferroelastic domain behavior in the annealed sample. Luo et al (1998) found that serrated parallel stripped domains in the tetragonal PMN-PT phase would developed at the crystal surface if crystals are subjected to rapid cooling from paraelectric to ferroelectric phase. We have also observed 90 domain structure of tetragonal phase with SEAM

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(Fig.3.14) (Liao, 1999). If the crystal is annealed at 300 for 3 hours and then slowly cooled in furnace, 90 domain structure of tetragonal phase could still be observed in local regions but show different characteristics: 1. Smaller domain sizes, i.e. averaged domain width is far smaller than that before annealing. 2. Only 90 domains in some local regions, as illustrated in Fig.3.14(b). Only 90 domain structure of tetragonal phase could be shown prior to annealing (Fig.3.14(a)); after annealing (Fig.3.14(b)), many complicated domain structure emerge in sample except some small-sized domain structures in Fig.3.15 and 3.16. This kind of domain structure was not reported earlier. As shown in Fig.3.15(b), two-fold, symmetric domains, i.e. butterfly domains, could be clearly seen. Fig.3.16(b) shows butterfly domains in larger areas. Butterfly-shaped domains have the features of twin domains as described in usual domain study, and their sizes vary from a few micrometers to 50 micrometers. The domain structures with different contrast distributed radially around the symmetric center but the boundaries are not quite clear.

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Fig.3.14 The comparison of ferroelectric domain structures of 0.65PMN-0.35PT single crystals before and after annealing (a) Before annealing, f =154.5 kHz; (b) After annealing, f =154.8 kHz


Fig.3.15 Butterfly shaped domain structure in the annealed 0.65PMN-0.35PT single crystals (a) SEI; (b) EAI at f =154.8 kHz

In addition, the butterfly domain distribution is distinctly oriented, i.e. two-fold symmetrical butterfly-shaped domains are respectively parallel. According to Feng’s optical microscopy observation of in-situ crystal growth of the rhombohedral Ba2 NaNb5 O15 (BNN) crystal, a large amount of ferroelastic domains were found to exhibit typical cross-shaped morphology with specific orientation. Their analysis revealed that the cross-shaped domains strictly correspond to the etching pit of 1/2 [100] line dislocation, which reveals that the domains could be considered as ferroelastic domains induced by dislocation, and their orientation is governed by the oriented dislocations. Butterfly-shaped structure and cross-shaped ferroelastic domains are similar because both show some same characteristics including: 1. Biaxial symmetry.


2. Twin domain structure. 3. Specific orientation. In addition, M L Mulvihill et al (1997) also observed ferroelastic domains with optical microscope in PZN-PT material. The butterfly-shaped domain patterns were supposed to be ferroelastic domains induced by the presence of dislocations in the crystal during the slow annealing process. For crystals before annealing, no butterfly-shaped domains but tetragonal 90 domains are visualized because tetragonal 90 domains would inhibit the formation of butterfly-shaped ferroelastic domains during the rapid cooling process. Whereas in the slow annealing process ferroelastic domains besides tetragonal phase would develop in the stress concentration area with dislocations or defects.

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Fig.3.16

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Butterfly shaped domain structure in large area of the annealed 0.65PMN-0.35PT single crystals (a) SEI; (b) EAI at f =153.2 kHz


3. 11

SEAM imaging of Other Materials

In addition to observing domains in ferroic matetrial, SEAM also exhibit a powerful tool for imaging stress distribution, interior defects and other subsurface structures in ceramic coatings, metals, superconductor ceramics and MEMS devices.

3. 11. 1 Residual stress distribution in Ti3N4 coatings Ceramic coatings have been found used in many cases where their resistance to high temperature, wear and erosion are important. In the present work, we carried out in situ observations of the carbon steel surface with TiN coating to investigate the surface and subsurface structures of the ceramic coatings. The EAI reveals the residual stress distribution around the crack on the carbon steel surface. This experimental results show that SEAM is a promising way to understand the subsurface microstructure of coatings and the cohesion between the substrate and the coating. It is very important to investigate the situation between the substrate and the ceramic coating, especially, the stress situation. The area of stress concentration is often the origin of coating flaking and destruction. Ti, Ti(C, N) and TiN coatings were sputtered on carbon steel surface layer by layer, as schematically shown in Fig.3.17. Before spraying TiN coating, the carbon steel surface was roughened to enhance the bonding strength between coating and substrate, so the steel surface exhibited scratch traces. Fig.3.18(a) shows the SEI of sample surface, Fig.3.18(b) and (c) show the corresponding EAI at different frequency. Butterfly- shaped stress distribution could be clearly observed around defects in Fig.3.18(c). Since electron acoustic signals of stress distribution originate from thermal wave coupling, detection depth decreases and the contrast of stress distribution in Fig.3.18 (b)~(c) becomes more distinct and concentrated as frequency decreases. However, when the frequency changes to 150.6 kHz, the EAI in Fig.3.18(d) is close to the sample surface, thus no stress concentration and distribution could be observed. The stress distribution obtained from SEAM observation shows alternating zones of elastic and plastic deformation in subsurface specimen under the externally applied stress or internal stress.


Fig.3.17

TiN/Ti (C,N)/Ti (Carbon Stell) ceramic coating sample

The butterfly-shaped stress distribution in Fig.3.18(c) actually shows the relative value of stress. Its imaging comes from the thermal wave coupling, and the EAI contrast depends on linear expansion coefficient, elastic constant and the density of specimens. Stress distribution in different region will result in different thermal-elastic properties, and will consequently be demonstrated on EAI.

Fig.3.18 SEI (a) and EAI of TiN/Ti(C, N)/Ti ceramic coating at f=123.0 kHz (b), f=132.5 kHz (c) and f=150.6 kHz (d)

According to the detection depth equation from thermal wave coupling d t = (2κ ωρ C )1 2


where κ is thermal conductivity, ρ is density and C is thermal capacity. For TiN, κ=0.2888J·cm-1, ρ=5.4 g·cm-3 and C= 0.585 J·g-1·K-1 (or 0.14 cal·g-1·K-1). Thus the detection depth at different frequency in Fig.3.18 is dt =4.9 μm under f=123.0 kHz (Fig.3.18(b)); dt =4.7 μm under f=132.5 kHz (Fig.3.18(c)); and dt =4.4 μm at f=150.6 kHz (Fig.3.18(d)), respectively. So we could estimate the depth of the residual stress field existing below the surface of TiN coating to be 4.7 μm. Therefore given the ceramic coating thickness, the concentrated stress locations could be known.

3. 11. 2 Stress distribution in ferroelectric composites During preparation and investigation of ceramic materials, an important criterion to evaluate the sintered body is porosity and its distribution and morphology. All these parameters will affect mechanical and electrical properties of the sintered body. As far as we know, studies of natural surface pores and their surrounding stress distribution have not been reported so far. Fig.3.19(a) is the SEI of natural surface of a sintered PZT ceramics with

Fig.3.19

SEI (a) and EAI (b) of ferroelectric composites (f=257.2 kHz)


several pores. The EAI exposes residual stress fields at the subsurface around the pores. The stress distributions surrounding the pores exhibit the butterfly shape with 2-fold symmetry. Their spreading range and intensity are proportional to the pore size. In Fig.3.19(b), the stress fields marked by A, B correspond to two big pores in Fig.3.19(a), and two smaller stress fields by C, D for two smaller pores. In fact, the ceramic density increases as the material sintering proceeds, but the final body will always contain some porosity. The mechanical property of the special interface between the pores and the ceramic grains could be influenced by the expansion or contraction of pores during the sintering process, resulting in the residual stress distributions developed around the pores interface. It is interesting that the stress field surrounding the pores has not only the same pattern of plastic and elastic zones, but also the same orientation (Liao, Yang, Jiang, et al, 1999).

3. 11. 3 Stress distribution in Si3N4 and ZrSiO4 ceramics As a typical high-temperature structural material, Si3N4 ceramic possesses many excellent properties including high strength, wear resistance and thermal shock resistance, so this ceramic material has wide and promising applications. Another ceramic material is ZrSiO4 composite, which is also a promising high-temperature structural material due to its excellent properties including low linear expansion coefficient, low thermal conductivity, high chemical stability, excellent thermal shock resistance and desirable high-temperature mechanical properties. SEAM investigation on residual stress distribution induced by Vicker indentation on Si3N4 ceramic and ZrSiO4 composite will be discussed in this section. The EAI results indicate that the residual stress field in the specimen shows alternating distribution of elastic and plastic deformation (Zhang, Jiang, Shi, et al, 1997). The SEI and EAI of Vicker’s indentation on Si3N4 ceramic at different frequency are shown in Fig.3.20, similar imaging results for ZrSiO4 composite are seen in Fig.3.21. The accelerating voltage for electron beam was 30 kV, modulation frequency varies from 60 kHz to 400 kHz, time integration constant was set to 100 μs and the detection sensitivity was 100 μV. In Fig.3.20(a) and Fig.3.21(a), only crack’s morphology and length could be observed. While in the acoustic imaging of Figs. 3.20(b)~(d) and Figs. 3.21(b)~(d), residual stress distribution from Vicker’s indentation could be clearly seen. Alternating dark and bright phases could be observed from the EAIs, indicating that the residual stress field distribute with alternately elastic and plastic areas, and stress concentration (plastic deformation areas) at the indentation center and crack tips. The experimental results described here is consistent with the theoretic calculations (Qian, 1995).


Fig.3.20 SEI (a) and EAI of residual stress distributions from Vicker’s indentation (10 kg) in Si3N4 ceramic at f=160.4 kHz (b), f=132.3 kHz (c) and f=103.5 kHz (d)

Residual stress distributions at EAIs evolve with the change of the electron beam modulation frequency. At higher frequencies, residual stress field distribute well symmetrically around the indenting center in Fig.3.20(b)~(c) and Fig.3.21(b)~(c). While at lower frequencies, great changes happen in the images of residual stress distribution. Cloud morphology of defects appears on the EAI of Si3N4 ceramic (Fig.3.20(d)), and non-symmetric distribution occurs along the diagonal of the cracks on the EAI of ZrSiO4 composite (Fig.3.21(d)). Such asymmetric residual stress distribution in Fig.3.21(d) might originate from inhomogeneous composition. Further study reveals that EAIs are more legible with higher resolution at the higher frequencies, and the EAI morphology is more similar to SEI at the higher frequencies. The frequency dependence of EAIs is in well agreement with the frequency dependence of detection resolution and depth in thermal wave coupling mechanism. Therefore the thermal wave coupling is the predominant mechanism governing the electron acoustic signal magnitude for these ceramics.


Fig.3.21 SEI (a) and EAI of residual stress distributions from Vicker’s indentation (10 kg) in ZrSiO4 composite at f=312.9 kHz (b), f=112.0 kHz (c) and f=62.29 kHz (d)

3. 11. 4 Stress distribution of Al metal Fig.3.22(a) shows SEI of aluminum with only Vicker’s indentation morphology and micro-crack length, while the acoustic contrasts of the residual stress field are clearly seen in Fig.3.22(b). The acoustic image indicates that the indentation area consists of alternating elastic and plastic areas instead of purely plastic deformation area, which directly reveals the fluctuant residual stress distribution by indentation (Zhang, Jiang, Shi, et al, 1997). Quantitative analysis of stress distribution could be obtained in Ref. (Qian, 1995). The above results demonstrate that SEAM could be utilized to observe the residual stress field due to Vicker’s indentation on ceramics and metals. The residual stress field consists of alternating elastic and plastic deformation areas in surface and sub-surface, and the concentrated stress regions lie at the indentation center and the crack tips. At the surface layer or regions close to the surface, the indentation force is quite large so that the internal stress from the material’s inhomogeneity can not appear. But in deeper layers, the internal stress distribution from material’s inhomogeneity begins to occur in the acoustic image as the


indentation load decreases.

Fig.3.22

SEI (a) and EAI (b) of Al metal ( See Color Picture in Appendix)

3. 11. 5 Surface structures and internal defects in lead-free piezoelectric ceramics Fig.3.23(a) shows the SEI of lead-free piezoelectric NaBiTiO3-KBiTiO3 ceramic material, the corresponding EAI is observed in Fig.3.23(b). Both surface structures and defects are clearly shown in the SEI and EAI with nearly same resolution (Yin, Zeng, Li, 2003). Fig.3.24 (a) and (b) show the SEI and EAI of BaTiO3 crystal, respectively. Both figures show the surface scratches due to the specimen machining, but the EAI also exhibits a microcrack extending from top left to bottom right, as indicated by arrows in the figure. The microcrack is not found on SEI, so it’s from the subsurface, which means EAI could be used for non-destructive observation of sub-surface structures (Kohler, Schubert, 2002).


Fig.3.23

SEI (a) and EAI (b) of NaBiTiO3-KBiTiO3 ceramic


Fig.3.24 SEI (a) and EAI (b) of BaTiO3 single crystal (magnification 150 , image size 800 μm 800 μm, f=119.0 kHz) (Courtesy to Kohler, Schubert, 2002)

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3. 11. 6 Phase transitions in superconductor ceramics Fig.3.25 shows the temperature dependence of electron acoustic signal magnitude of superconductor ceramic Bi2Sr2CaCu2Ox, with solid line and dash line indicating heating and cooling processes, respectively. An abrupt jump of electron acoustic signal magnitude occurs in the temperature ranging from 200K and 230K, and exhibits the characteristics of first-order phase transition. This indicates that the electron acoustic techniques could be used to investigate the phase transition behavior for solid state materials (Urchulutegui, Piqueras, Liopis, 1989).

Fig.3.25 Temperature dependence of electron acoustic signal magnitude in Bi2Sr2CaCu2Ox superconductor ceramic(Courtesy to Urchulutegui, Piqueras, Liopis, 1989)


3. 11. 7 SEAM imaging of MEMS devices Fig.3.26 shows the SEI and EAI of MEMS device. The SEI in Fig.3.26 (a) exhibits only surface morphology of the device. while the EAI in Fig.3.26 (b) shows the mark number of “B2271” inside the device. (The structure and the material of the device were not open due to the commercial confidentiality by US manufacturer.). This observation reveals the unique ability for non-destructive imaging internal structures by SEAM. In fact, SEAM has also been successfully employed to observe the buried structures of other semiconductor devices (Takenoshita, 2000; Takenoshita, 2002).

Fig.3.26

3. 12

SEI (a) and EAI (b) of MEMS device

Scanning Probe Acoustic Microscopy

Scanning probe microcopy develops rapidly since the invention of scanning tunneling microscope (STM) by G Binnig and H Rohrer in 1981 (1982). With the advent of advanced sciences and technologies, nanoscale physical effects of functional ceramics, as well as miniaturization and integration of devices have become increasingly more important, thus understanding local properties of functional ceramics and devices at nanometer scale is need urgently. Recently, various SPM-based techniques for nanopolar regions in functional ceramics have been rapidly developed, in which SFM in piezoelectric mode has been widely used for observing domain structures and characterizing polarization reversal behavior (Auciello, Gruverman, Tokumoto et al, 1998).


Scanning near-field acoustic microscopy was firstly proposed by P Gunther et al (1989). Scanning near-field acoustic microscopy has two operation modes: specimen vibration and tip vibration. In the case of specimen vibration, the acoustic vibration generated by transducer attached to the specimen bottom is to be detected by the probe that contacts the up surface of specimen; while for tip vibration mode, an alternating voltage is applied between the probe and the specimen, and the transducer attached to the bottom surface detects the acoustic vibration. Since the imaging method is developed on the basis of scanning force microscopy, it is generally referred to as scanning probe acoustic microscopy (SPAM). Here we present the two acoustic microscopy imaging of ferroelectric materials.

3. 12. 1 Tip-vibration mode scanning probe acoustic microscope Fig.3.27 shows the schematic of tip-vibration mode scanning probe acoustic microscope. The SPAM was built based on the commercial atomic force microscope. Besides the basic parts of AFM, a signal generator has been mounted between the conducting probe and the bottom surface of specimen so as to produce acoustic vibration in the bulk specimen directly. A piezoelectric transducer attached to the bottom of specimen was used to detect the acoustic vibration and transfer it into the measured electrical signals, and then was processed by the lock-in magnifier connected to the transducer. Both surface morphology and the acoustic images could be observed by SPAM simultaneously (Liu, Heiderhoff, Abicht, et al, 2002).

Fig.3.27

Schematic of scanning probe acoustic microscope


3. 12. 1. 1 Operation principle In SPAM, an alternating voltage is applied between the conducting probe and the specimen, and acoustic vibration will be generated from functional ceramic due to the converse piezoelectric effect. The acoustic waves containing its local interaction with the sample is introduced to the piezoelectric transducer, and transferred into electrical signals. A lock-in magnifier will collect and process the electric signals and send to computer for imaging. During scanning point-by-point on x-y plane, acoustic waves will be reflected or refracted to generate acoustic-wave intensity variation when encountering micro-defects, local inhomogeneity, and different oriented domains and grains. The variation in acoustic wave intensity will lead to contrast on the acoustic image showing the sample’s microstructures. 3. 12. 1. 2 SPAM imaging of ferroelectric domain structures We firstly describe the applications of SPAM to observe domain configurations in BaTiO3 and PLZT ferroelectric ceramics. A. Ferroelectric BaTiO3 ceramics (Liu, Heiderhoff, Abicht, et al, 2002) Fig.3.28(a) and (b) show surface morphology and acoustic microscopy of BaTiO3 ceramic, respectively. Only surface topography with grains of various orientations and grain boundaries is visualized in Fig.3.28 (a), while the acoustic image (Fig.3.28 (b)) reveals several groups of ferroelectric domains, strip domain structures, and even serration-shaped domain structures.

Fig.3.28 The surface topography image(a) and acoustic microscopy image of ferroelectric domain (b) of BaTiO3 ceramics ( Courtesy to Liu,Heiderhoff,Abicht,et al,2002) (See Color Picture in Appendix)


B. Ferroelectric (Pb, La)(Zr, Ti)O3 (PLZT) ceramics (Yin, Li, Zeng, et al, 2004) Fig.3.29 (a) shows the topography of the PLZT ceramics obtained by SPAM. The main features in the topography are the surface scratches. No information associated with the domain structure was revealed. Fig.3.29 (b) is an acoustic image of the PLZT ceramics using the SFM acoustic mode, which clearly reveals fingerprint patterns related to domains structures with antiparallel polarization (Urohulutegui, 1993). These stripes show a pronounced and different contrast in the different areas (A, B, C) appearing in Fig.3.29 (b) due to different crystallographic orientations of the individual grains. The bend and split of the domain marked by arrows at the grain boundary regions may be attributed to the existence of inhomogeneous lattice distortions or spatial defects which destroyed the continuity of ferroelectric domains and minimized the elastic energy and depolarization fields at the grain boundaries. The fingerprint patterns in Fig.3.29 (b) are relatively regular and are almost periodically spaced in each individual grain. The stripe-structures in Fig.3.29 (b) are 300 nm in width. Fig.3.30 illustrates the spatial distribution of the phase difference signal cos φ (φ phase difference between the modulation voltage and the piezoresponse signal). The phase difference corresponds to the direction of the polarization vector (sgn(Pz)=cosφ). The phase images indicate that the bright and black domain images can be ascribed to the antiparallel polarization, i.e. -c and +c domains, respectively. They are consistent with the former analysis, and give confirmation of the observed 180 domain structure in the acoustic image. In addition, the domains appear as bending and splitting patterns at the grain boundaries, as shown in Fig.3.29 (b).

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Fig.3.29 The surface topography image (a) and the corresponding acoustic microscopy image (b) of PLZT transparent ceramics ( See Color Picture in Appendix)

Fig.3.31 shows a magnified acoustic imaging of the stripe-domains. A line scan (C-D) drawn in Fig.3.31 (a) is shown in Fig.3.31 (b). The amplitude distribution of the acoustic signal is periodically spaced, reflecting the periodical domain distribution and the homogeneous piezoelectricity in PLZT ceramics. In addition, by observing the


minimum resolvable peak separation marked by arrows in Fig.3.31(b), the estimated spatial resolution of acoustic mode SFM imaging in this work is about 70 nm.

Fig.3.30

The acoustic microscopy image of PLZT transparent ceramics

Fig.3.31 The acoustic microscopy image (a) and (b) the acoustic signal of line scan C-D in Fig.3.31(a) ( See Color Picture in Appendix)


3. 12. 1. 3 Ferroelectric domain imaging mechanism In SPAM, When an ac voltage is applied between the conductive tip and the rear surface of the sample, the tip acts as an exciting source and generates acoustic waves in the piezoelectric PLZT ceramics due to the converse piezoelectric effect. Therefore the generation of the acoustic signal in the SPAM is attributed to piezoelectric coupling. The tip has a spherical shape with a radius of curvature of 10 nm, and the excited local region under the tip is about 50 nm in lateral dimensions, 200 nm in depth (Balk, 1980) and is independent of the surrounding area. Each local region of the sample can be considered as an independent piezoelectric vibrator. The vibration model of each piezoelectric vibrator is related to the interaction between the applied electrical field and the polarization vector. Obviously, this kind of piezoelectric vibration is closely connected with the locally effective piezoelectric constant ( dijeff ) of the sample. The effective dijeff piezoelectric tensor is not only related to the piezoelectric coefficients d31, d33, d15, d22 of the PLZT sample which has a La/Zr/Ti ratio of 6/65/35, but also related to the deviation angle of the spontaneous polarization vector from the normal direction (upward) of the sample surface. This composition is in the rhombohedral ferroelectric phase and has eight domain states with spontaneous polarization vectors oriented in the directions of cube diagonals of the cubic phase. The multi-domain state system with different orientations will give different contribution to the domain contrast observed in the amplitude image in the SFM acoustic mode, which can be considered as an equivalent function of the local effective piezoelectric constants. Different dijeff will undoubtedly give rise to different domain contrast in the acoustic image. For a local area with the spontaneous polarization orientation close to the normal direction of the sample (i.e. nearly parallel to the ac-field direction), its larger dijeff will lead to stronger piezoelectric vibrations and corresponding bright domain tones as shown in the A area of Fig.3.29(b). In contrast, the local area with larger deviating angle θ has a smaller dijeff , thus causing weaker piezoelectric vibrations and showing a gray, dark domain tone, as reflected in the B and C area of Fig.3.29(b). So the differences in the effective piezoelectric constants of the local areas under the SPAM tip are the intrinsic characteristics resulting in the antiparallel domain contrasts of the PLZT ceramics in SPAM.

3. 12. 2 Sample-vibration mode scanning probe acoustic microscopy In the present study, low-frequency (below 100 kHz, even down to several kHz) near-field acoustic microscopy for imaging ferroelectric microstructure was successfully developed based on the commercial atomic force microscope,


which provides a subsurface depth profile information and as well as spatial resolution at the 10-nanometer scale. The domain configurations of ferroelectric material were clearly visualized by this newly, non-destructive, low-frequency acoustic mode scanning probe microscopy technique with high resolution. The acoustic response origin of ferroelectric domain and the acoustic imaging mechanisms are discussed in terms of interaction between the excited acoustic wave and ferroelectric domains. 3. 12. 2. 1 Introduction Nondestructive characterization of subsurface structure is of great interest for almost every area of science and engineering. The conventional far-field acoustic/optical microscopy has limited spatial resolution close to 100 μm due to classical diffraction. Scanning probe microscopy based on atomic force microscopy offers higher resolution to image depth information. Especially in recently reported scanning near-field ultrasound holography (Shekhawat, Dravid, 2005), which provides nanoscale-resolution images of internal substructures of microelectronic materials and malaria parasites in red blood cells. In ferroelectric fields, the acoustic mode scanning probe microscopy including atomic force acoustic microscopy, ultrasonic force microscopy, and scanning probe acoustic microscopy have the advantage of visualizing nanoscale domain structures, elastic properties or subsurface configurations. Although the AFM-based acoustic microscopy has the advantage of visualizing nanoscale structures inside the specimen, they are generally operated at high modulation frequency ranging from a few hundred kilo Hertz to one mega Hertz. 3.12.2.2 Low frequency scanning probe acoustic microscopy (LF-SPAM) Figure 3.32 shows the schematic sketch of the custom-built low-frequency near-field acoustic probe microscopy based on the commercial atomic force microscope (AFM) (Seiko Inc. Japan) operating in acoustic mode (a), in which a piezoelectric transducer was bonded onto a sample holder of an AFM, and the sample was glued on the transducer. An external function generator provides sinusoidal excitation signals applied to the piezoelectric transducer, which gives rise to flexural vibration and thereby emits acoustic waves into the sample, causing a periodic local deformation of the sample’s surface underneath the tip. The motion of the scanning force microscopy tip is detected optically by laser beam deflection from the backside of the cantilever. The detected signal is then fed to a lock-in amplifier for imaging. Figure 3.32 also includes the set-up of piezoresponse mode AFM (b), which was used to obtain the piezoresponse signal of ferroelectric domains with one-to-one corresponding relationship between the acoustic signal and piezoresponse signal.


Fig.3.32

3. 12. 2. 3

The schematic diagram of low-frequency scanning probe acoustic microscopy

LF-SPAM imaging of piezoelectric, ferroelectric microstructures

Here we introduce the applications of LF-SPAM in observing locally ferroelectric domains and elastic structures in lead-free piezoelectric ceramics and piezoelectric single crystals. A. Lead-free Bi4Ti3O12 piezoelectric ceramics (Zeng, Yu, Zhang, et al, 2005) The novel acoustic probe technique was used to perform imaging of ferroelectric microstructures in piezoelectric ceramic and ferroelectric single crystals. Figure 3.33 shows the topography image and its corresponding acoustic image of Nb-doped Bi4Ti3O12 piezoelectric ceramics at frequency of 30 kHz. Plate-like structure parallel to the long axis of the grain and small rod-like structure were clearly shown in individual grains (Fig.3.33(a)), while interesting superfine structures appear in Bi4Ti3O12 grains (Fig.3.33 (b)). Scale-like elastic structure perpendicular to the long axis of the plate-like middle grain can be clearly seen in the zoom image (Fig.3.33 (c)). The line scanning signal of line C-C’ in Fig.3.33(c) reveals the elastic structures arrange in a scale-like form as observed in Fig.3.33(d). The height distance between the scale structures is down to 12 nm. The scale structures of Fig.3.33(c) can be clearly seen in the corresponding three-dimensitional configuration of Fig.3.33(e), Noted the observed angle is rotated so as to get an optimum view effect. The observed elastic structure parallel and perpendicular to the long axis of plate-like grain are as a result of the crystal structure arrangement in the Bi4Ti3O12 grains consisting of layers of perovskite-like Bi2Ti3O10 units sandwiched between bismuth oxide (Bi2O2)2+ layers. The scale-like elastic structures in the acoustic image


Fig.3.33 The topography image (a) and the corresponding acoustic image (b) of Bi4Ti2.98Nb0.02O12 piezoelectric ceramics at modulation of 30 kHz. Fig.(c) is the zoom image of the selected area indicated by white square of Fig.(b). Fig.(d) is the line scanning signal of the white line as indicated in Fig.(c). Fig.(e) is the three dimensitional image of Fig.(c) with the rotated view angle. Fig.(f) and (g) are the topography image and the corresponding acoustic image of Bi4Ti2.985W0.015O12 ceramics ( See Color Picture in Appendix)


reflect the preferential grain growth behavior of the anisotropic particle evolving from a small plate parallel to the c axis into the final plate-like grain during sintering. In addition, our studies demonstrated that small amount of Bi2O3 crystllites occurs in the sintering process of the Bi4Ti3O12 ceramics. Figs. 3.33 (f) and (g) present the topography and the corresponding acoustic image of W-doped Bi4Ti3O12 ceramics, respectively. The topography exhibit smooth, plate-like grains, while inhomogeneous, elastic microstructures are shown clearly in the acoustic image (Fig.3.33(g)). In which 100~300 nm sized small particles appear inside the grain. These particles are directly related to low-melting Bi2O3 liquid phase forming in the sintering process. B. ZnO resistor ceramics Figure 3.34 gives high-resolution acoustic image of piezoelectric ZnO resistor. The smaller grains (100~200 nm in size) could be found to accumulate at the grain boundaries (Fig.3.34(a)) and very difficulty to be distinguished. In contrast, the corresponding acoustic image (Fig.3.34(b)) shows clearer individual boundary due to different elastic properties among them. The line scanning acoustic signal (Fig.3.34(d)) demonstrates that this kind of new acoustic technique approaches ultrahigh resolution up to 4 nm. Such ultrahigh resolution acoustic probe microscopy plays an important role in revealing the buried elastic structures of functional materials and devices. C. Relaxor-type PMN-PT crystals and PLZT ceramics (Zeng, Yu, Li, et al, 2005) The acoustic imaging ferroelectric microstructures could be operated at very low frequency, even down to several kHz and several hundreds Hz. Fig.3.35 gives an acoustic image of ferroelectric domains of relaxor-type PMN-PT single crystal. The Topography image shows only polished scratch feature (Fig.3.35 (a)) at the modulation frequency of 6.42 kHz, while finger-print domain configurations exhibit in the corresponding acoustic image (Fig.3.35 (b)). Interestingly at such low frequency, the piezoresponse image shows no any remarkable piezoresponse contrasts except noise signals. Fig.3.36 shows the piezoresponse and the acoustic images of ferroelectric PLZT ceramics at low frequency. The domain configurations were surprisingly visualized by the acoustic probe technique at such low frequency. The acoustic signal of ferroelectric domains becomes clearer when the modulation frequency is increased to 9.3 kHz (Fig.3.36(f)). The irregular patterns in the acoustic image are in well agreement with those in the corresponding piezoresponse image (Fig.3.36(c)), and reflect the ferroelectric features of the polar regions in the relaxor-type PLZT ceramics.


Fig.3.34 The topography image (a) and the corresponding acoustic image (b) of ZnO resistors. Fig.(b) is the zoomed image of the selected area indicated by white square in Fig.(b). Fig.(d) is the line-scanning signal of line A-A in Fig.(c)

Ą


Fig.3.35 The topography image (a) and the corresponding acoustic image (b) of PMN-PT single crystal at the modulation frequency of 6.42 kHz ( See Color Picture in Appendix)

Fig.3.36 The topography image (a), the corresponding piezoresponse image at 9.3 kHz (b), (c), the acoustic image at 0.5 kHz (d), (e), and at 9.3 kHz of PMN-PT single crystal (f) ( See Color Picture in Appendix)


3. 12. 2. 4 Modulation frequency dependence of acoustic imaging (Zeng, Yu, Zhang, et al, 2005) In the conventional acoustic imaging approaches, subsurface structure could be visualized by collecting acoustic signal at different modulation frequency. Our present near-field acoustic probe microscopy also bear this unique feature. Fig.3.37 presents topography image of Bi4Ti3O12 ceramics and its corresponding acoustic image at different modulation frequency of 30 kHz, 50 kHz, 70 kHz, 90 kHz, 120 kHz, respectively. Low-frequency (30 kHz) acoustic image in Fig.3.37(b) shows distinct features in individual grains, especially in grain 1 and 2. Slim, strip-like structure parallel to the long axis of the grain and small rod-like structure were clearly shown in these grains. Some grains exhibit circular structure, as seen in grain 3 and 4. The smallest elastic structure is down to 60 nm. With increasing the modulation frequency, the elastic contrasts become significant, the fine elastic structures are not as clear as that in Fig.3.37(b). The structures with different contrasts reflect different elasticity of domain structures in the grains. The bright parts represent stiffer regions, while the dark less stiff. In comparison of the elastic

Fig.3.37 The topography image (a) and the corresponding acoustic image of Bi4Ti2.98Nb0.02O12 piezoelectric ceramic at different modulations frequency of 30 kHz (b), 50 kHz (c), 70 kHz (d), 90 kHz (e), and 120 kHz (f), respectively ( See Color Picture in Appendix)


response image at 120 kHz modulation frequency in Fig.3.37(f) with that Fig.3.37(a), it can be found that no significant difference from the topography image except the enhanced contrasts of some edges. Thus we can conclude that the acoustic contrasts in Figs.3.37(b)~(f) exhibit a subsurface origin with different elasticity of domain structures. In addition, another interesting phenomenon could be found in our study. It has to be noted that the acoustic image shows more stereo than the corresponding piezoresponse one. Such features reveal that the acoustic signal at a low frequency is stronger than that at a high frequency due to remarkable decrease of ferroelectric indentation nearby the tip at the low frequency. While higher tip indentation forces possibly occur at high frequency, reducing the contrast of the acoustic image due to inhibiting acoustic signal excitation and measurement. The observed acoustic signal behavior is related to the contact stiffness response between the tip and sample surface at different modulation frequency, as discussed later. 3. 12. 2. 5 Multi-response of ferroelectric domains to local stress (Yu, Zeng, Ma, 2005) For most applications, PMN-PT single crystals are expected to bear electrically and thermally induced stresses as well as mechanical loads. However, the response of the 180 domain structure in the rhombohedral phase to mechanical stresses has not been reported due to its difficult visualization in the conventional techniques. Here we exploited the low frequency scanning probe acoustic microscopy to study the responses of domains in PMN-PT crystals to Vicker’s indentation. Typical topography (Fig.3.38(a)) and LF-SPAM image (Fig.3.38(b)) before Vicker’s indentation in the PMN-PT single crystal are displayed. The domain contrast mechanism is related to the different electric polarization of ferroelectrics.

e

Fig.3.38 Typical (a) topography and (b) domain structure image obtained by LF-SPAM in Pb(Mg1/3Nb2/3)O3–0.33PbTiO3 single crystal


e

The fingerprint 180 domain pattern appears on the left of Fig.3.38(b), and tweed, non-180 domains on the left. It is well known that the indentation-induced domain switching patterns are related to strains. For revealing the nature of the stress-related domain switching, the strain-free fingerprint pattern 180 domain areas were selected for indentation-induced response. The mechanical and electric response to Vicker’s indentation is shown in Fig.3.39, exhibiting three major features. Firstly, the indentation load produced plastic deformation underneath the indenter and along the diagonals, which left square indentation and protuberances on the crystal surface. Secondly, indentation

e

e

Fig.3.39 LF-SPAM images of Vicker’s indentation-induced crack pattern and domain pattern for Pb(Mg1/3Nb2/3)O3–0.33PbTiO3 single crystal under 50 g indentation load for 10 s. (a) two-dimensional and (b) three dimensional topographies, (c) the acoustic image, and (d) the illustration of the indentation-induced mechanical and electric response


induced microcracks extended along some preferential directions. Thirdly, new lamellar domains in directions perpendicular to the cracks near the indentation formed on the matrix domain structure. In Fig.3.39(a) and (c), it can be seen that indentation induced crack pattern did not emanate from the stress focused directions as diagonals of the indent as in Fig.3.39(a) and (b). Furthermore, less plastic deformation appeared in the crack regions. It is suggested that the crack directions are perhaps the special cleavage directions. According to the crystal symmetry, the other equivalent directions, 90 away from the crack directions, are available at these conditions. But cracking did not occur in the experiment. This result means that along the direction perpendicular to the crack a high toughness and fracture resistance exist. Furthermore, the domain structures switched in the directions also suggest that the direction probably possesses high piezoelectric properties.

e

3. 12. 2. 6 Low-frequency acoustic imaging mechanisms (Zeng, Yu, Li, et al, 2005) The visualized acoustic image contrasts could be conceptually understood from the point view of interaction between the excited acoustic wave and microstructures in the sample. In the low-frequency acoustic microscopy, the sample vibration is modulated sinusoidally in time by the underneath ac field-excited transducer. The tip vibrates in phase with the sample surface displacement due to low frequency modulation. The modulated tip amplitude reflects actually the tip-sample surface interaction stiffness (contact stiffness, k*). Supposing that the tip-surface contact obeying the Hertzian contact-mechanics model, k* can be expressed as k * = 3 6 E *2 RF , where R is the radius of the tip apex, F the applied load by tip to

the sample surface, E* is the reduced elastic modulus of the tip-sample surface system. E* could be expressed as

1 E*

=

1 − ν s2 1 − ν c2 , here E1, E2 and νs, νc are + E1 E2

the Young’s moduli and the Poisson’s ratio of the surface and the cantilever, respectively. It is clear that the value of k* depends not only on the applied static load and the tip geometry, but also on the elastic properties of the tip and the sample. Besides, especially for the ferroelectric material, k* is related to the interaction between the local ferroelectric microstructure (ferroelectric domain) beneath the tip and the excited acoustic wave in the bulk sample. It is well known that the bulk electric polarization due to a preferred macroscopic orientation of the domains influences the elastic response of a ferroelectric material when stresses are applied. Thus, the acoustic contrast reflects contact stiffness difference of ferroelectric domains depending on the local elastic fields-induced acoustic vibration in the sample during imaging at low frequency. However, an in-depth analysis concerning the physical mechanism of acoustic


imaging is necessary in the further work. Although the present low-frequency acoustic microscopy demonstrates high resolution imaging capability, it so far only limits to ferroelectrics. However, based on the interaction between the acoustic wave and microstructures in the specimen, other attempts to detect microelectronic structure defects and biology cells should be further expected for its potential applications in semiconductor systems and biology.

3. 13

Comparisons of SEAM with SPAM

Both SEAM and SPAM are near-field acoustic microcopies, in which imaging resolutions is irrelevant to acoustic wave length. No pre-treatment such as polishing and etching etc., is required for observing domain structures with these two methods. SEAM is based on Scanning Electron Microscope, while SPAM is based on Scanning Probe Microscope, but neither of the two microcopies requires high cost. In contrast, there are many differences in the two microcopies. SEAM takes use of the modulated electron beam as the excited source, and has a wide space for observation. For SEAM, it could non-destructively image the subsurface structure information with sub-micrometer resolution. However, SPAM could provides a more resolved acoustic image of superfine structures with ultrahigh resolution up to several nanometers. The similarities and differences between SEAM and SPAM are quite helpful for understanding material microstructures and imaging mechanisms. In practical applications, the two methods could supplement each other. In comparison with other near-field microcopies, SEAM and SPAM play a unique and significant role in imaging and characterizing material microstructures.

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3 Near-field Acoustic Microscopy of Functional Ceramics London Balk L J, Kultscher N (1984) Nonlinear scanning electron acoustic microscopy. Journal de Physique 45(C2):869–872 Bihan R L (1989) Study of ferroelectric and ferroelastic domain structures by scanning electron microscopy. Ferroelectrics 97:19–46 Binnig G, Rohrer H, Gerber Ch (1982) Tunneling through a controllable vacuum gap. Appl Phys Lett 40:178–180 Brandis E, Rosencwaig A (1980) Thermal-wave microscopy with electron beams. Appl Phys Lett 37:98–100 Bresse J F, Papadopoulo A C (1987) Be incorporation in heavily doped molecular beam epitaxy grown GaAs: Evidence of nonradiative behavior by cathodoluminescence and electron acoustic measurements. Appl Phys Lett 51:183–185 Bursill L A, Lin P J (1984) Microdomains observed at the ferroelectric/paraelectric phase transition of barium titanate. Nature 311:550–552 Cargill G S (1980) Ultrasonic imaging in scanning electron microscopy. Nature 286:691–693 Carslaw H S, Jaeger J C (1973) Conduction of Heat in Solids (the 2nd Ed). Oxford University Press, London Davies D J (1983) Scanning Electron Microscopy Part : 1163 Davies G, Howie A, Staveley S L (1982) Scanning electron acoustic microcopy. In:Proceeding of SPIE-The International Society for Optical Engineering 368:58 Furuhata Y, Toriyama K (1973) New liquid-crystal method for revealing ferroelectric domains. Appl Phys Lett 23:361–363 Gunther P, Fischer U Ch, Dransfeld K (1989) Scanning near-field acoustic microscopy. Appl Phys B 48(1):89–92 Hatano J, et al (1973) Improved powder-Pattern technique for delineating ferroelectric domains. Jpn J Appl Phys 12(10):1644–1645 Holstein W L (1985) Image formation in electron thermoelastic acoustic microscopy. J Appl Phys 58(5):2008–2021 Hooton J A, Merz W J (1955) Etch patterns and ferroelectric domains in BaTiO3 single crystals. Phys Rev 98(2):409–413 Ikegami S, Ueda I (1967) Mechanism of Aging in Polycrystalline BaTiO3. J Phys Soc Jpn 22:725–734 Jiang F M, Kojima S, Zhang B Y, et al (1999) Application of SEAM and SAM To Ferroelectric And Ferroelastic Crystals. Ferroelectrics 217: 335–341. Jiang F M, Kojima S, Zhang B Y, et al (1999) Application of SEAM and SAM To Ferroelectric and Ferroelastic Crystals. Ferroelectrics 222:237–242. Kirkendall T D, Remmel T P (1984) Journal de Physique 45(C2):877 Kohler B, Schubert F (2002) Some aspects of photo and particle acoustic methods. Molecular and Quantunm Acoustics 23:225–238 Kulcsar F (1956) A microstructure study of barium titanate ceramics. J Am Ceram Soc 3991:13–17

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Kultscher N, Balk L J (1983) J Scanning Electron Microscopy Part I:33 Lemons R A, Quate C F (1974) Acoustic microscope-scanning version. Appl Phys Lett 24:163–165 Liao J (1999) Ferroelectric domains and their dynamic behavior. Ph. D thesies, Shanghai Institute of Ceramics, Chinese Academy of Sciences Liao J, Yang Y, Jiang X P, et al (1999) Scanning Electron Acoustic Imaging of Residual Stress Distributions in Ceramic Coatings and Sintered Ceramics. Mater Lett 39: 335–337 Liu X X, Heiderhoff R, Abicht H P, et al (2002) Scanning near-field acoustic study of ferroelectric BaTiO3 ceramics. J Phys D: Appl Phys 35:74–87 Liu X X, Balk L J, Zhang B Y, et al (1998) Scanning electron acoustic microscopy for the evaluation of domain structures in BaTiO3 single crystal and ceramics. J Mater Sci 33:4543–4549 Little W A (1955) Dynamic behavior of domain walls in barium titanate. Phys Rev 98(4):978–984 Merz W J (1954) Domain formation and domain wall motions in ferroelectric BaTiO3 single crystals. Phys Rev 95(3):690–698 Mulvihill M L, Cross L E, Cao W W, et al (1997) Domain-related phase transitionlike behavior in lead zinc niobate relaxor ferroelectric single crystals. J Am Ceram Soc 80(6):1462–1468 Murata K, et al (1971) Monte carlo calculations on electron scattering in a solid target. Jpn J Appl Phys 10(6):678–686 Opsal J, Rosencwaig A (1982) Thermal-wave depth profiling: theory. J Appl Phys 53(6):4240–4246 Piqueras J (1994) Scanning electron acoustic microscopy of electronic materials. Materials Science and Engineering B 24(1–3):209–212 Qian M L (1995) Study of Characteristics of ultrasound pulse thermoelastically generated b a laser pulse. Acta Acousta 20(1): 1–10 Qian M L, Cantrell J H (1989) Signal generation in scanning electron acoustic microscopy. Mater Sci and Eng A 122(1):57–64 Rosencwaig A, Gersho A (1976) Theory of the photoacoustic effect with solids. J Appl Phys 47(1):64–69 Rosencwaig A, Opsal J (1985) Thermal Wave Imaging with Thermoacoustic Detection. IEEE UFFC 33 (5) 516–528 Shekhawat G S, Dravid V P (2005) Nanoscale imaging of buried structures via scanning near-field ultrasound holography. Science 310(5745):89–92 Takenoshita H (2000) Nondestructive internal observation of metal-oxide-semiconductor LSI designed by 0.8 µm rule. Jpn J Appl Phys A39 (9):5312–5315 Takenoshita H (2002) Comparative study of scanning electron microscopy and electronacoustic microscopy images. Jpn J Appl Phys 41(1):70–72 Tennery V J, Anderson F R (1958) Examination of the surface and domain structure in ceramic barium titanate. J Appl Phys 29:755–758

3 Near-field Acoustic Microscopy of Functional Ceramics Urchulutegui M, Piqueras J, Liopis J (1989) Scanning electron-acoustic microscopy of MgO crystals. J Appl Phys 65(7):2677–2680 Urchulutegui M, et al (1993) Scanning electron acoustic microscopy of misfit dislocations in InGaAs/GaAs superlattices. J Phys D: Appl Phys 26:1537–1539 Yin Q R (1985) Thermal wave microscope and its applications. Nature (in Chinese) 8:344–348 Yin Q R, Jiang F M, Hui S X (1994) Piezoelectric electron-acoustic probe (PEAP) and some applications. Ferroelectrics 151:97–102 Yin Q R, Li G R, Zeng H R, et al (2004) Ferroelectric domain structures in (Pb, La) (Zr, Ti) O3 ceramics observed by scanning force microscopy in acoustic mode. Appl Phys A 78:699–702 Yin Q R, Liao J, Jiang F M, et al (1999) Electron acoustic imaging of ferroelectric domain and mechanism analysis on BaTiO3 ceramics. Ferroelectrics 231:1–18 Yin Q R, Tang Z, Zhang H J, et al (1990) Scanning electron acoustic microscope. Journal of Electron Microscopy (in Chinese) 9:53–55 Yin Q R, Wang T, Qiang M L (1991) Techniques of optic-acoustic and optic-thermal and their applications. Science press, Beijing Yin Q R, Zeng H R, Li G R, et al (2003) Near-field acoustic microscopy of ferroelectrics and related materials. Mater Sci & Eng B 99:2–5 Zeng H R, Yu H F, Li G R, et al (2005) Local elasticity imaging of ferroelectric domains in PMN-PT single crystals by low-frequency atomic force acoustic microscopy. Sol. Sta. Comm. 133(8):521–525 Zeng H R, Yu H F, Zhang L N, et al (2005) Local elastic response of individual grains in lead-free Nb-doped Bi4Ti3O12 piezoelectric ceramics. Phys Stat Sol (a): Rapid Research Letters 202(4):R41–R43 Zhang B Y, Jiang F M, Hui S X, et al (1996) Signal generation of ferroelectric semiconductor ceramics in scanning electron-acoustic microscopy. Journal of Function Materials and Devices (in Chinese) 2:53–57 Zhang B Y, Jiang F M, Shi Y, et al (1997) Scanning electron-acoustic imaging of residual stress distributions in aluminum metal and ZrSiO4 multiphase ceramics. Appl Phys Lett 70:589–591 Zhang B Y, Jiang F M, Yang Y, et al (1996) Electron acoustic imaging of BaTiO3 single crystals. J Appl Phys 80(3):1916–1918 Zhang B Y, Jiang F M, Yang Y, et al (1996) Piezoelectric electron acoustic probe of domain structures in ferroelectric ceramics BaTiO3. Ferroelectric Letters 22:21–25 Zhang B Y, Jiang F M, Yin Q R (1995) Observation of growth defects with SEAM to GaAs/GaAs. In:Proc of Inter Sixth Beijing Conf and Exhi on Instru Analysis, A: Electron Microscopy, Beijing Zhang B Y, Jiang F M, Yin Q R (1996) Theory and applications of scanning electron acoustic microscope. J Inorg Mater (in Chinese) 11(2):207–213 Zhang B Y, Yin Q R (1996) Piezoelectric electron acoustic study of domain structures in ferroelectric ceramics BaTiO3. Ferroelectrics Lett 22:21–25

Piezoresponse Force Microscopy of Functional Ceramics

Piezoresponse force microscopy (PFM) was customer-built based on the commercial atomic force microscope and used to characterize ferroelectric domains of functional materials. The PFM imaging contrast mechanism, domain configuration and their evolution behavior under the inhomogeneous tip fields in ferroelectric thin film, lead-free piezoelectric ceramics and relaxor-type single crystals are presented in details in this chapter.

4. 1

Introduction

Ferroelectric materials exhibit a wide spectrum of physical effects including piezoelectricity, pyroelectricity, photoelectricity, photo-acoustic effect, photorefractive behavior, and non-linear optical activity. These properties offer great potential application in electronic devices such as non-volatile dynamic random access memories, micromotors, microactuators, thin film capacitors, pyroelectric infrared detectors, and ultimately optical waveguides. These unique properties are closely related to their microstructural domain patterns and dynamic response to mechanical, electrical and optical loads. The formation and kinetics of these domains are governed by the underlying atomistic structure, and their interaction with domains walls and material imperfections. Recent advances of ferroelectrics have resulted in the development of low-dimensitional ferroelectrics, such as one and two dimensitional ferroelectric nanostructures including ultrathin epitaxial films, nanoscale capacitors, nanowires, nanoshell tubes, nanorods and free-standing nanocrystals. These structures may spur the development of novel nanoscale devices. As the size of the devices reduces to submicro- and nano-scale, the motion and interaction of domain walls significantly influences the overall

4 Piezoresponse Force Microscopy of Functional Ceramics

physical properties of the devices. Thus it is very important to understand the physical phenomenological behavior of ferroelectric domains and their dynamic evolution in nanoscale volumes. Typical techniques for revealing domain structures include chemical etching, polarization microscopy, powder decoration, scanning electron microcopy, transmission electron microcopy, and scanning electron acoustic microscopy, etc. However, these techniques have some drawbacks to some extend such as complicate sample pretreatment, original domain damage and low imaging resolution. To overcome these limitations, one has to develop novel, high-resolution imaging technique to probe domain configurations and local properties at micro- and nanometer scales. At present, scanning probe microscopy (SPM)-based techniques provide a powerful tool to imaging and characterizing the static and dynamic properties of ferroelectrics. For the first time, SPM provides a non-destructive, ultrahigh resolution tool to visualize nanoscale domain structures and to evaluate local electromechanical and elastic properties. SPM has also opens a new window for high-density data storage in ferroelectric material via SPM-base nanoscale domain patterning. Currently, the most popular SPM method for imaging ferroelectric domains is piezoresponse force microscopy (PFM), based on detecting the piezoelectric vibration of ferroelectric surfaces induced by the ac field between the probing tip and the sample. Its high spatial resolution, easy implementation, effective manipulation of nanoscale domains and locally quantitative characterization capabilities make it a versatile tool for nanoscale ferroelectric studies. In this chapter, we describe the basic principle of PFM and its applications to characterizing ferroelectric domains of bulk ceramic, single crystals and ferroelectric thin films as well.

4. 2

History and Development of Scanning Probe Microcopy

The advent of scanning probe microscopy is the human’s persistent pursuit of ultrahigh resolution microscope. The first concept, scanning probe near-field optical microscope (SPNOM), was proposed by E. H. Synge in early 1928 in order to surpass the 0.2μm resolution limit of the conventional optical microscope due to light diffraction. The imaging resolution of SPNOM could reach up to 0.01 μm (Synge, 1928). In 1982, scanning tunneling microscope (STM) with atomic scale resolution was invented by G Binnig et al (1982). STM brought a milestone to the development of super-resolution microscopes, and revolutionized the conventional microscope. Since then, a large number of high-resolution probe microscopes with different modes have been developed. As an important branch of SPM, scanning force microscope (SFM) is termed as a generalization to atomic force microscope (AFM), in which it measures the

4. 3

Piezoresponse Force Microscopy

electrostatic and magnetostatic interactions, as well as long-range Van der Waals forces between the tip and the specimen (Binnig, Quate, 1986; Sarid, 1991). The most popular ones in SFM include friction force microscope (Mate, McClelland, Erlandsson, et al, 1987), tapping-mode AFM (Martin, Williams, Wickramasinghe, 1987), magnetic force microscope (Martin, Wickramasinghe, 1987), electrostatic force microscope (Martin, Abraham, Wickramasinghe, 1988), Kelvin probe force microscope (Nonnemacher, O’Boyle, Wickramasinghe, 1991), piezoresponse force microscope and other important microscopes. Piezoresponse force microscopy was firstly introduced in 1992 by P Guethner and K Dransfeld to detect domain structures in ferroelectric copolymer films (Birk, Glatz-Reichenbach, Li, et al, 1991; Güthner, Dransfeld, 1992). Soon after that, it became the most effective approach for nanoscale study and control of ferroelectric domains in bulk crystals and thin films. In 1994, K Franke et al (1992) firstly reported PFM imaging of PZT thin films, and in 1996, PFM measurement of local hysteresis loops of nanoscale domain reversal was presented by T Hidaka et al (1996). Since then, PFM has become a powerful tool to perform studies of nanoscale domain configurations, domain dynamic behavior including domain nucleation and growth, polarization and reversal, polarization relaxation and local hysteresis spectroscopy.

4. 3


Piezoresponse-mode scanning force microscopy, i.e. piezoresponse force microscopy has become an effective research instrument for characterization of domain structures and physical properties on nanoscale (Hong, Woo, Shin, et al, 2001; Gruverman, Pignolet, et al, 2000; Chu, Szafraniak, Scholz, et al, 2004).Here it describes PFM imaging principle and its imaging features.

4. 3. 1 Operation principle

ˈ

In PFM based on converse piezoelectric effects of specimens, imaging of piezoresponse-mode SFM is obtained by means of detecting local vibration of piezoelectrics under applied AC voltage. Fig.4.1 shows the schematic of PFM that is built up on the basis of commercial AFM. Its working mechanism includes the following processes: the PFM probe scans specimen surface in a contact mode, and AC voltage from signal generator is applied between the probe tip and bottom electrode of specimen; the laser intensity reflected by the back of PFM cantilever will be detected by a photodiode detector, and the collected signals actually reflect piezoelectric vibration of specimen; signals are amplified by a lock-in amplifier


connected with reference signals from the bottom electrode of the specimen.

Fig.4.1

Schematic diagram of piezoresponse force microscopy

During the imaging processes, an alternative voltage is applied on the top surface of specimen through the probe that actually acts as a mobile top electrode. The AC voltage frequency should be far lower than the resonant frequency of the cantilever in order to avoid imaging interference. The applied AC voltage (Acosωt) will induce the sample vibration with the same frequency due to the converse piezoelectric effect, and will also generate vibration of 2ω in frequency due to electrostrictive effect. Ferroelectric domains could be observed by controlling the first resonance (piezoresponse signals) Ipr, in which the amplitude could be expressed as: I pr = kd33V0 cos ϕ and d33 = 2Q11εε 0 P

(4.1)

Where k is a calibration constant (depend on the sensitivity of optical detector PSD of SFM), d33 is the piezoelectric coefficient, Q11 is the electrostrictive constant, ε is the dielectric constant, and ε 0 is the vacuum dielectric constant. Polarization orientation is determined by the phase difference between the piezoresponse signals and applied AC voltage. Since the piezoresponse amplitude is related to the piezoelectric coefficient d33 and the phase is related to polarization orientation, the regions with the opposite polarization orientation would vibrate out-of-phase under the external applied AC voltage, and consequently exhibit different contrasts on piezoresponse image, as shown in Fig.4.2. PFM imaging mechanism of ferroelectric domains could be further understood with detailed discussions on the deformation behavior of the


4. 3

cantilever system (Abplanalp M, Eng L M, Gunter P, 1998). The cantilever resonance originates from the elongation, contraction, shear and other deformation of the sample under the AC voltage V AC=V0 cos ( ω t) due to the converse piezoelectric effect. The thickness vibration could be expressed Δ t = dzzV0 cos ( ω t), where dzz is the piezoelectric coefficient perpendicular to specimen surface, and the transverse deformation of Δ x= –dxxz V0 cos ( ω t) = – d15 sgn (Px) V0 cos ( ω t). According to piezoelectric effect, vertical and horizontal vibration of micro-cantilever in PFM could be expressed as following equations: t

Δz vertical ∞ ∫ d33 EAC (t )dt

(4.2)

0

t

Δzhorizontal ∞∫ d15 EAC (t )dt 0

Therefore, the out-of-phase and the in-phase piezoresponse signal of ferroelectric domains could be collected, and more comprehensive information about domain configurations could be acquired.

Fig.4.2

Principle of piezoresponse SFM

(a) No topographic contrast if no voltage is applied; (b) A change in thickness occurs for a positive voltage applied to the tip; (c) Opposite thickness changes for a negative voltage

Fig.4.3 shows the schematic graph of detecting bending and torsion vibration for PFM cantilever. The two kinds of vibrations will be transformed into laser deformation signals and sent to quadrant photoelectric detector PSD, in which vibration signals could be divided into two parts: vertical bending vibration signal ((A+B)–(C+D)) and horizontal torsion vibration signal ((B+D)–(A+C)). Then the weak signals will be amplified by a lock-in amplifier with the applied voltage V(t) as the reference signal. Hence c domains corresponding to the vertical vibrations and a-domains from the horizontal torsion vibrations could be visualized with PFM, respectively.


Fig.4.3

Schematic of detecting bending and torsion vibration for PFM cantilever

4. 3. 2 PFM imaging features In piezoresponse force microscopy,the effective piezorespone volume is related to some imaging features like the effective depth of electric field,resolution and imaging parameters like tip load,as discussed below. 4. 3. 2. 1

Effective depth of electric field

During piezoresponse imaging, effective depth of electric field generated by the alternating voltage applied on probe tip is the response depth of detecting ferroelectric domains, i.e. Z-sensitivity of Piezoresponse Force Microscope. The sensitivity also depends on the specific materials. For example, the theoretical calculations show that 90% of the alternating voltage reaches the vertical area 30% of specimen thickness for ferroelectric material with dielectric constant εr of 100. However, for both 120 μm-thickness BaTiO3 single crystal film and 400 nm-thickness PZT 20/80 epitaxial film, the effective piezoresponse layer is about 30 nm according to the finite element analysis (Abplanalp, Günter, 1998). The PZT/SrRuO3(110)/SrTiO3(001) layered film by Ahn et al revealed that the SrRuO3 layer presented a steady variation of electrical resistance during polarization for 400 nm-thickness PZT layer, which means that the ferroelectric domains reversal induced from film surface could reach the sample bottom, and reversibly change carrier concentration. According to C Hamagea et al (2001), two factors should be considered for the effective electric field depth. 1. Dielectric constant of materials. Since electric field intensity is in reverse proportion to dielectric constant, high dielectric constant leads to low electric field intensity within specimen, thus piezoresponse area will be quite shallow in specimens with high dielectric constant, i.e. piezoresponse almost lies at the surface layer. For example, the effective field depth of c domains is 20 times that of a domain in BaTiO3 single crystal. 2. Contact area/film thickness ratio. Larger contact area (e.g. blunt probe tip) and

4. 3


thinner ferroelectric films helps to decrease the inhomogeneity of electric field induced by PFM probe, thus piezoresponse region of thin films is larger than that of thick films. Actually it could be deemed that the piezoresponse of thick films (thicker than 1μm) reflects domain structures at the surface layer, while the piezoresponse of thin films (thinner than 500 nm) originates from the whole contribution across the sample thickness. 4. 3. 2. 2 Resolution Imaging resolution is a vital parameter for any microscopic technique. For PFM, its resolution determines the smallest domain size to be detected, and it depends on many factors, e.g. PFM probe diameter, interaction volume between tip and the specimen, contact area, contact force and tip characteristics. PFM resolution is generally defined as the distance between two neighboring antiparallel domains. Up to now, the highest transverse resolution of ferroelectric domain structure detected with PFM is about 8~10 nm (Dunn, Shaw, Huang, et al, 2002). 4. 3. 2. 3 Experimental parameters During PFM imaging, the interactive forces between PFM probe and specimen consist of three terms (Hong, Woo, Shin, et al, 2001; Hong, Shin, Woo, 2002): F = Fel + Fpiezo + Fcap

(4.3)

Where Fel is the electrostatic force, Fpiezo is the piezoresponse component, and Fcap is surface capacitance force between micro-cantilever and specimen surface. Thus piezoresponse contrasts of ferroelectric domains actually include both the piezoresponse contrast and electrostatic contrast. Obviously the Fpiezo component needs to be increased to obtain desirable piezoresponse image contrast. PFM resolution strongly depends on probe, especially the elastic properties of microcantilever material. Meanwhile, it is also closely related to the experimental parameters including probe diameter, interactive force, contract area (i.e. effective area). On the basis of comprehensive deliberation of dependence of piezoresponse contrast on experimental parameters, S V Kalinin and D A Bonnell (2002) established the piezoresponse contrast mechanism map. Their studies indicated that the tip-surface contact area is smaller and the local electrostatic force is far higher than piezoresponse force under smaller interactive force; while larger tip-surface contact area in the case of the stronger indenting force, which promotes piezoresponse. Stiffer micro-cantilever and larger interactive force is lucrative for piezoresponse, while limper probe facilitates electrostatic contribution. However, polarization of material under probe tip would be restrained and plastic deformation would be produced when excessively large interactive force is applied. In one word, proper imaging parameters are crucial for acquirement of high-quality piezoresponse images in practical experiments.


4. 4

PFM Imaging of Ferroelectric Domains

PFM is increasingly becoming a powerful microscopy to imaging nanoscale domains in ferroelectrics.Here we introduce PFM applications to ferroelectric ceramics, thin films and single crystal.

4. 4. 1 Ferroelectric thin films As an important kind of information and functional materials, ferroelectric films have a significant and promising application future in the fields, such as microelectronics due to their outstanding properties of ferroelectricity, piezoelectricity, pyroelectricity, and non-linear optical effects. Currently numerous significant progresses in the research and development of ferroelectric films have been achieved, e.g. unit size of high density ferroelectric memory has been reduced to 30~100 nm (Alexe, Harnagea, Erfurth, et al, 2000). Therefore, local studies of ferroelectric domains, mechanisms of polarization and reversal, polarization fatigue and polarization retention, are important for FeRAM and MEMS devices. PFM is undoubtedly becoming a powerful tool for the present ferroelectric study. 4. 4. 1. 1 Perovskite structure Figs.4.4 show the topography and the corresponding piezoresponse image of ferroelectric PZT thin film. As indicated by Fig.4.4(a), each grain presents multi-domain structure instead of single-domain state. Domain structures in thin films basically originate from the misfit strain between the film and the substrate, and internal stresses among grains (Pertsev, Zembilgotor, Tagantsev, 1998), so that the spontaneous polarization within a grain is randomly oriented and causes a complicate contrasted piezoresponse image with bright, dark and gray tone. The regions with distinct bright and dark contrasts are evidently 180 domains with antiparallel polarization orientation. However, most grains show gray contrast in the piezoresponse image. A. Gruverman suggested several possible reasons (Gruverman, Auciello, Tokumoto, 1998): 1. possibly vertical stack of several grains with random polarization orientations on the film surface; 2. a-domains with polarization vectors parallel to film surface; 3. domains with polarization vectors away from the normal direction of film surface; 4. non-crystal or non-ferroelectric structure; For PZT ferroelectric thin films with complete multi-crystalline perovskite structure, the piezoresponse image contrast mainly comes from the random orientation of grains since different grain orientations lead to different polarization

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4. 4 PFM Imaging of Ferroelectric Domains

orientations and resultantly different contribution to piezoresponse contrast. Therefore, the complicated domain contrast on piezoresponse image of PZT thin film could be attributed to crystal orientation and domain configuration within grains (Zeng, 2003; Zeng, Li, Yin, et al, 2003b).

Fig.4.4

Topography image (a) and the corresponding piezoresponse image (b) of ferroelectric PZT thin film

Fig.4.5 shows the topography and the corresponding piezoresponse image of another region on the PZT ferroelectric thin film. It could be found that the


Fig.4.5

Topography image (a) and the corresponding piezoresponse image (b) of ferroelectric PZT thin film

piezoresponse image of ferroelectric domains exhibits complicated contrasts, but some grains show characteristics of layered domains, which are 90 domains. Such domains are clearer in Fig.4.6. This kind of domain structures is extraordinarily obvious in Fig.4.7, which shows three-dimensional polarization distribution of ferroelectric domain. It could be found that c domains present quite strong contrast in vertical piezoresponse image of Fig.4.7(b), and presents weak contrast in lateral piezoresponse image of Fig.4.7(c), but a-domains

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present just the opposite variation. Fig.4.7 has demonstrated a very important feature of 90 domain structures within various grains, i.e. periodic distribution of domain structure with about 30 nm in size, which originate from the periodic misfit strain between film and substrate.

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Fig.4.6 Topography image (a) and the corresponding piezoresponse image (b) of ferroelectric PZT thin film ( See Color Picture in Appendix)


Fig.4.7 Topography image (a), vertical piezoresponse image (b) and the lateral piezoresponse image (c) of ferroelectric PZT thin film

4. 4. 1. 2 Bismuth layer structure Bismuth layer-structured ferroelectric thin films such as SrBi2Ta2 O9 (SBT) and Bi 4Ti3 O12 are actively studied currently due to their excellent properties of fatigue resistance, e.g. no sign of fatigue after 1012 cycles of polarization switching. In addition, these materials are lead-free, which is quite beneficial for environment protection. In spite of lower spontaneous polarization than PZT-based materials, these materials are promising to replace PZT-based


materials in microelectronic devices due to the two advantages mentioned above to promote development of MEMS devices from environmentally friendly materials. The topography images of the epitaxial Bi4Ti3O12/LaNiO3/SrTiO3 thin film prepared by pulse laser deposition (PLD) (Harnagea, 2001) and reveal several rectangular grains; the piezoresponse image shows strong domain contrast within the grains, and bright, dark and gray contrasts are alternately distributed along long axis of grains, with domain walls parallel to c-axis of grains.

4. 4. 2 Ferroelectric ceramics Transparent and lead-free piezoelectric ceramics are two types of potential ceramics in micromechanical and electro-optical fields.Their domain configurations visualized by piezoresponse force microscopy are given as below. 4. 4. 2. 1 Transparent materials With their outstanding electrooptical and ferroelectric characteristics, transparent ferroelectric ceramics of PLZT and novel PZN-PLZT have found wide applications in electronic and optical devices. Properties of PLZT and PZN-PLZT ceramics are closely related with their ferroelectric domains, which could be usually observed with optical microscope due to their transparency (Zhu, Ao, Yao, et al, 1990). However, optical microscopy has its limitations. Primarily its resolution is insufficient so that fine structures of ferroelectric domains can not be reflected. Fig.4.8 shows the topography image and piezoresponse image of transparent PLZT ceramic. The piezoresponse image presents fingerprint-like domain structures. For


Fig.4.8 The topography image (a) and the piezoresponse image (b) of transparent PLZT ceramics ( See Color Picture in Appendix)

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more legible domain structures, Fig.4.9 shows the piezoresponse image of 180 antiparallel domains structures with average size of about 200 nm in a smaller scanning area. Fig.4.10 shows the piezoresponse image in another region of the same specimen. It could be found that there’s no piezoresponse signal in some small regions as indicated by arrows, and these non-ferroelectric regions demonstrate the inhomogeneity in microstructures and properties of PLZT ceramic.

Fig.4.9

Piezoresponse image of transparent PLZT ceramics


Fig.4.10

Piezoresponse image of transparent PLZT ceramics

Fig.4.11 shows the topography image and the piezoresponse image with a scanning area of 15 15 μm2 of PZN-PLZT electrooptic ceramic. The piezoresponse image shows two types of 180 antiparallel domains structures with strong contrasts. One is the fingerprint shaped domain structure similar to that in PLZT ceramic, and the other is the domain structure with bright and dark contrasts in large areas.

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Fig.4.11 The topography image (a) and piezoresponse image (b) of PZN-PZT transparent ceramics

Actually the large bright and dark domains consist of many small fingerprint shaped domains. Thus it could be deduced that the large domains come from the mergence of the small domains, and domain orientations and sizes are related with grain orientations and sizes. Furthermore, domain structures of PZN-PLZT ceramic might be also related to the local composition undulations. As some experiments indicate, domain sizes of PLZT ceramic decrease when increasing lanthanum concentration (Xu, Kim, Li, et al, 1996). 4. 4. 2. 2

Lead-free piezoelectric ceramics

As a kind of environmentally friendly materials, lead-free doped Bi4Ti3O12 piezoelectric ceramics have been attached with increasingly more importance (Noguchi, Miyayama, 2001). Among all bismuth layer-structured ferroelectrics, Bi4Ti3O12 has the highest Curie point TC=675 . At temperatures below TC, the material presents a ferroelectric phase with the symmetry class of monoclinic point group m and at temperatures above TC, the material exhibits a parallelelectric phase with tetragonal system. Bi4Ti3O12 has greatly anisotropic properties due to its layered structure. The material shows a quite strong spontaneous polarization Ps =50 10–2C/m2 along the (010) plane of monoclinic system, i.e. the a plane, and makes a angle of 4.5 with the a-axis of the monoclinic system. Thus Ps has a polarization component of 50 10–2 C/m2 along c axis. Various kinds of domains walls will emerge in the crystal due to its symmetry, thus the lead-free Bi4Ti3O12 piezoelectric ceramic presents a complicated

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domain patterns. Fig.4.12 (a) is the topography of niobium doped Bi4Ti3O12 piezoelectric ceramic, showing long plate shaped grains similar to the structure of the film with the same composition. For more distinct piezoresponse image, Figs. 4.12(b) and (c) show the topography and piezoresponse images of the grain as framed in Fig.4.12(a). It could be seen that bright, dark and gray domain contrasts are alternately distributed with the smallest domain of around 70 nm. Our studies show that the niobium doping could greatly affects the domain configurations in the Bi4Ti3O12 piezoelectric

Fig.4.12 Topography image of niobium doped Bi4Ti3O12 piezoelectric ceramic (a), topography image (b) and the corresponding piezoresponse image (c) of the grain framed in (a)


ceramic, as shown in Fig.4.13. Fig.4.13 shows the typical surface morphology (microstructures) (Figs. 1(a) and (c)) and the corresponding domain structures (Figs. 13(b) and (d)) of Bi4Ti2.98Nb0.02O12 and Bi4Ti2.92Nb0.08O12 ceramics. Our present studies show that the average grain size decreases with Nb-doping, and the domain structures in Bi4Ti2.98Nb0.02O12 and Bi4Ti2.98Nb0.08O12 ceramics (Figs.4.13 (c) and (d)) are quite different. In Bi4Ti2.98Nb0.02O12 ceramics, most of domain walls are parallel to the c-axis (see Fig.1(b)). Whereas, in Bi4Ti2.92Nb0.08O12 ceramics, domain walls are mainly perpendicular to the c-axis and parallel to the a–b plane (Fig.4.13(d)) (Chu, Zhang, Xu, et al, 2004).

Fig.4.13 (a) Typical surface morphology and (b) the corresponding domain structures of Bi4Ti2.98Nb0.02O12 ceramics. (c) The typical surface morphology (microstructures) and (d) the corresponding domain structures of Bi4Ti2.92Nb0.08O12 ceramics


4. 4. 3 Ferroelectric single crystals Solid-solution relaxor ferroelectric single crystals Pb(Mg1/3Nb2/3)O3-PbTiO3 (PMN-PT) and Pb(Zn1/3Nb2/3)O3-PbTiO3 (PZN-PT) have become a research hotspot among ferroelectric materials due to their excellent electromechanical coupling properties (Service, 1997). The material with the composition of morphotropic phase boundary has unique properties, i.e. in spite of spontaneous polarization along direction, its optimal properties are presented along direction; its electromechanical coupling coefficient k33 could reach 94%, piezoelectric coefficient d33>2500 pC/N, and its maximum strain could reach 1.7%. Applications of the materials could bring revolutionary breakthroughs to the fields including medical ultrasonic imaging, sonar communication, and ultrasonic transducer. From the first-principle study, Fu and Cohen have revealed that the exceptional properties of these materials come from polarization deflection induced by electric field (Fu, Cohen, 2000). At present, a large amount of research work has been focused to understand the physical essences of PMN-PT or PZN-PT and its dynamic behaviors and evolution trends under applied electric field. PFM is undoubtedly a powerful instrument for the research and observation of nano-scale domain structure of PMN-PT single crystal.

4. 4. 3. 1 PMN-PT single crystal Ferroelectric domains in PMN-PT single crystal exhibit very interesting phenomena.We introduce such new findings from the following five aspects. A. Antiparallel ferroelectric domains Figs.4.14(a) and (b) show the topography image and the corresponding piezoresponse image of PMN-30%PT, and Fig.4.14(c) shows the piezoresponse result of a smaller scanning area. The distinct fingerprint domain structures with bright and dark contrast are a typical image of antiparallel polarized domain structure. The piezoresponse images also show inhomogeneity of domain regions since these domain regions consist of many smaller domains with opposite contrast, in which the smallest domain width is about 100 nm. The domain structure of PMN-PT is obviously different from the strip domain structure of ordinary ferroelectrics such as BaTiO3 and TGS. The unique domain structures of these materials might result from polar nanoregions in the relaxor ferroelectrics. Polar nanoregions developed at temperature far above the dielectric peak temperature (Tm), and would interact with each other during the cooling process under zero electric field to form complicated domain patterns


and lead to the dielectric relaxation at temperatures around Tm (Zeng, 2003).

Fig.4.14 The large scanning topography image (a), the corresponding piezoresponse image (b) of relaxtor-type PMN-PT ferroelectric single crystals, (c) the piezoresponse image in a small scanning areas

B. Hierarchy domain structures Fig.4.15 show typical topography and piezoresponse images of (110)-oriented PMN–33% PT single crystals. Though the contrast is not so strong, it can be seen in Fig.4.15(b) that the scanned area is distributed with parallel micron-sized domains, which are about 20~30μm in width. This is in agreement with the results obtained by another method with low resolution, in which the large domain contrast is strong while


the fine domain structure is obscured. This may be a consequence of the fluctuations in the components and structure under cooling. In Fig.4.15(c), which is magnified from part of the region scanned in Fig.4.15(b), it is clearly seen that large micron-sized domains are not uniform and actually consist of either small fingerprint pattern domains (FPDs) or tweed pattern domains (TPDs), in which sizes are down to about 1μm and about 0.5μm, respectively. This is probably corresponding to the rhombohedral and tetragonal phases coexisting in the PMN–33% PT single crystal.

Fig.4.15 Topography image (a) and piezoresponse image (b) of the (110)-oriented PMN–33% PT single crystal in the same region. The piezoresponse image (c) magnified from part of (b).(d) Linescan image of the topography with the direction marked in Fig.(a) ( See Color Picture in Appendix)


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The FPDs or 180 antiparallel domains are only ferroelectric, not ferroelastic, because no strain is involved during domain formation or switching. On the other hand, the TPDs or non-180 (71/109 and 90 ) domains are not only ferroelectric but also ferroelastic because these domains can also be switched or reoriented by an applied stress, and strain is involved during domain formation, switching or reorientation. There are two types of linear domain walls as illustrated by the solid line and dashed line in Fig.4.15(b). The former separates the regions with the same TPDs and the latter separates the regions with different TPDs and FPDs. The line scan image of the topography with the direction marked in Fig.4.15(a) is shown in Fig.4.15(d). It shows that the fall in height at the boundary of the adjacent TPD regions is obvious while the others are not so obvious. And, the result also demonstrates that the region of the TPDs has a tilt perpendicular to the long direction in the topography, and the region of FPDs is flat in the topography. Other typical TPD structures have been imaged in other scanning areas of the same PMN–33% PT single crystal and magnified in Figs.4.16(a)~(c). It is a result of the strain developing in the course of domain formation. Additionally, it should be noted that the TPDs are split into opposite contrast stripes. The reason will be given in detail in the subsequent discussion. Fig.4.17(a) shows another area of the piezoresponse image of the (110)-oriented PMN–33% PT single crystal. It is seen that the FPDs and TPDs coexist in the same micron-sized domain and a series of gradual transformation regions have been clearly ‘frozen’. It is reasonable to anticipate the FPDs and TPDs to be intrinsically related. A piezoresponse image of the extended region on the right side of the scanned area in Fig.4.17(a) was also recorded, as illustrated in Fig.4.17(b). It shows that the TPDs are discontinuous as in Fig.4.15(c). 180 domain walls and non-180 domain walls are shown in Fig.4.17(c), whose centre is marked with ‘C1’ in Fig.4.17 (b). The 180 domain walls are for reducing the depolarization energy, and the non-180 domain walls are for reducing the strain energy that is produced in the phase transformation. During the cooling process from high temperatures, PMN–33% PT undergoes a paraelectric-to-ferroelectric phase transformation. Large-sized domains are formed with the fluctuating of components and structure as shown in Fig.4.15(b). At that time, random fields induced by built-in charge disorder already have emerged, and are responsible for the formation of nanopolar clusters. With further decrease in the temperature, the interaction between nanopolar clusters was enhanced, which could have changed the short-range order state into the longrange order state, and thus promoted the reorientation and mergence of nanoclusters into irregular FPD, as shown in Fig.4.15(c). In this manner, the elastic strain energy of the ferroelectric transformation could have been stored in atomic level shuffles. The growth of the FPDs would not stop until the ferroelectric interactions between the nanopolar clusters and random field were almost equivalent. With the transformation, the lattice deformation became larger and larger. The strain

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at the atomic level could not be contained any more. For strain accommodation, the deformation in topography occurs as a result of the following processes. The crack pattern domains begin to nucleate and grow in the FPD to increase the domain wall density, which makes the strain compatible. And the domain structure becomes more complicated and irregular. It is well known that 180 domain walls have an arbitrary orientation, while non-180 domain walls have perfect orientation, (100) or/and (110) in PMN–PT single crystals. With the increase of the ferroelastic TPD wall density, the orientations became more and more obvious through the domain walls’ splitting, reorientation and movement as displayed in Fig.4.17 (c) (Yu, Zeng, Chu, et al, 2004).

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Fig.4.16 Piezoresponse image of (110)-oriented PMN–33% PT single crystal (a) and its magnified images (b) and (c) with different magnification ( See Color Picture in Appendix)


Fig.4.17 Piezoresponse images of (110)-oriented PMN–33% PT single crystal. Tweed pattern and fingerprint pattern subdomains coexist in the same high hierarchy domain (a). The gradual transformation region between tweed pattern and fingerprint pattern subdomains (b), piezoresponse image (c); the scanning centre point corresponds to ‘C1’, marked in (b) ( See Color Picture in Appendix)

C. Domain inhomogeneity Fig.4.18 (a) shows the domain images of the 15×15 μm2 scanning areas in the investigated PMN-30%PT single crystals. The corresponding topography image of Fig.4.18(a) is neglected due to only representing the surface image of the (001)oriented single crystal without any additional information. The piezoresponse image contrast in Fig.4.18 (a) is determined by the out-of-plane component of polarization, where black and white areas correspond to the opposite polarization directions. Irregular, fingerprint-like domain patterns with average size of 200 nm


were clearly observed in Fig 4.18 (a). The remarkable contrasted domain regions are not uniform but a mixture of small domain regions with opposite contrasts. The irregular domain patterns are not the unique features of the domain structures visualized in the sample. Another type of regular domain structure was surprisingly imaged in other scanning areas of the same PMN-30%PT single crystal, as shown in Fig.4.18 (b). Fig.4.18 (b) shows the regular, narrow strip-like domain patterns as compared with irregular domains in Fig.4.18 (a). The large strip domain patterns in Fig.4.18 (b) actually consist of much finer domain structures down to 30 nm. The complex, spatial inhomogeneity of domain structure of the (001)-oriented PMN-PT single crystal are ascribed to the relaxor nature of the unpoled PMN-PT single crystals with a rhombohedral ferroelectric state. Usually, the ferroelectric domain state in relaxor-type ferroelectrics is dependent on the random internal field and the interaction between nanopolar clusters. The random field was induced by structural disorder, compositional fluctuation and other defects, as shown clearly by HRTEM observations. For the PMN-PT family, the PMN is a typical relaxor in which random fields prevent the formation of long-range ferroelectric order on cooling under zero electric field. The PT addition favors the long-range ferroelectric order; but the relaxor features still remain in zero-field-cooled (ZFC) PMN-30%PT single crystals when the random field arising from structure irregularity dominates over the ferroelectric interaction between nanopolar clusters. Thus, the fingerprintlike domain pattern reflects the existence of short-range order manifested in formation of nanoscale polar clusters, which is the principal feature of the relaxor PMN-PT single crystal in the ZFC state. However, when the spatial inhomogeneity of local structure decreases in PMN-PT single crystals, the influence of random internal field was inhibited, and the ferroelectric interaction between nanopolar clusters would be enhanced, which can induce the short-range order rhombohedral ferroelectric state into the rhombohedral ferroelectric state with long-range ordering, thus, promoting reorientation and merging of nanoclusters into regular, strip-like micron-sized domains upon cooling as shown in Fig.4.18 (b). The presence of nanoscale inclusions inside these domains of Fig.4.18 (b) may represent remains of the polar clusters in which random fields were especially strong to prevent their reorientation and growth. Assuming that there are local areas where the random fields are almost equivalent to the ferroelectric interaction between nanoclusters in (001)-oriented PMN-PT single crystals, it is possible to observe shorter strip-like domain structures and fingerprint-like domain structures coexisting in these local areas. We actually obtained such an experimental result, as shown in Fig.4.18(c). The expected shorter strip domain patterns cover large scanning areas of Fig.4.18(c) with small areas of fingerprint domain patterns appearing in the down left corner of Fig.4.18(c). Thus we believed that random fields from nanoscale structure irregularity affect greatly the domain arrangement of PMN-PT single crystals. Such local random fields can


serve as the heterogeneous nucleation centers for polarization switching. It undoubtedly plays a crucial role upon the ferroelectric phase stability and superior properties of the (001)-oriented PMN-PT single crystals (Zeng, Yu, Chu, et al, 2005).

Fig.4.18

The piezoresponse image of three different scanning areas (15 ×15 μm2) in (001) PMN-30%PT single crystal

D. Three-dimension polarization distribution Fig.4.19 shows the piezoresponse image of the three-dimensional polarization distribution of ferroelectric domains in PMN-PT single crystal. Figs. 4.19(a), (b) and (c) respectively show topography image, vertical piezoresponse image and lateral piezoresponse image along direction of PMN-30%PT. Figs. 4.19(d), (e), (f) show the results of another region on the same specimen. The piezoresponse


Fig.4.19 The topography image (a,d), vertical piezoresponse image (b,e) and lateral piezoresponse image (c,f) of (001)-oriented PMN-PT single crystals


image clearly shows the antiparallel domain structure. PMN-30%PT relaxor single crystal is rhombohedra phase at room temperature, and rhombohedra ferroelectrics have four independent piezoelectric coefficients, i.e. d33, d22, d31 and d15. During the piezoresponse imaging, the piezoelectric vibration signals are directly proportional to the effective piezoelectric coefficient along different component directions. The vertical and lateral effective piezoelectric coefficients are: 2 3 2 2 eff d 33 = d15 − d 22 + d 31 + d 33 3 3 3 6 3 3 3 3 1 4 2 2 eff d15 = d15 + d 22 − d 31 + d 33 3 3 3 6 3 3 3 3 Since the spontaneous polarization of rhombohedral phase is along direction, it has eight domain situations, i.e. Ps[111] , Ps[1 1 1] , Ps[ 1 1 1] , Ps[ 1 11] , Ps[ 1 1 1 ] , Ps[ 1 1 1 ] , Ps[11 1 ] , Ps[1 1 1 ] . On the basis of the vertical piezoresponse

vibration and lateral torsion vibration as well as the bright and dark piezoresponse contrasts, it could be deduced that polarization in Fig.4.19(a), (b), (c) is likely to be Ps[111] , Ps[ 1 1 1 ] , Ps[ 1 11] , and Ps[1 1 1 ] while the polarization in Fig.4.19 (d), (e), (f) is likely to be Ps[1 1 1] , Ps[ 1 1 1 ] , Ps[ 1 1 1] , Ps[11 1 ] . Therefore the three-dimensional distribution of polarization is markedly significant for understanding of polarization orientations of ferroelectric domains (Zeng, Yu, Chu, et al, 2004).

edomain wall for tetragonal PMN-PT

E. Unusual piezoresponse at the 180 crystal

Fig.4.20 shows the typical topography and its corresponding domain structures on the (001) surface of PMN–40% PT single crystal imaged simultaneously by piezoresponse force microscopy. In the topography (Fig.4.20(a)), there is no additional information except for some surface scratches. The domain patterns are hierarchical. Two sets of {110} oriented, large domain structures are located in the miniature fingerprint domain structure. The angle between the two sets oriented domain structures is roughly 90 in Fig.4.20 (b). In the tetragonal phase ferroelectrics, the {110} oriented domains are compatible with stresses formed in the course of phase transformation. So the different {110} oriented domains can coexist in the same surface and grow, respectively. While they meet, they can intersect and stop there (see the right side in Fig.4.20 (b)) or grow through each other (see

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the left side in Fig.4.20 (b)). Fig.4.21 shows the topography, piezoresponse, phase and amplitude images obtained from the same region of the PMN–40% PT single crystal. The domains are 0.7 mm in width, which are nearly equal to those of the miniature fingerprint domain displayed in Fig.4.20(b). It is shown that they are continuous and oriented. From the line scan in the phase image, the adjacent domains have a 180 change in the phase. So it can be identified that the domain patterns are antiparallel domains. It is interesting to find that the amplitude values in the adjacent domains are nearly equal, but it has a sharp decrease at the domain wall regions in the amplitude line scan image, shown in Fig.4.21(e). This indicates that the piezoelectricities in the adjacent domains are almost equal but an obvious low in the domain wall (Yu, Zeng, Wang, 2005).

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Fig.4.20 The typical topography (a) and the domain structures (b) at the same area in (001) surface of PMN–40% PT single crystal. The angle between the two sets of oriented domain structures is roughly 90

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Fig.4.21 The topography (a), the piezoresponse (Acosφ) image (b), the phase (φ) image (c) and the amplitude (A) image (d) of the 180 domains at the same region in (001) surface of PMN–40% PT single crystals. The line scans of the Acosφ, φ, A images in the same direction, which was marked in (b), is displayed in (e)

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4. 4. 3. 2 PZN-PT single crystal The relaxor ferroelectric Pb(Zn1/3Nb2/3)O3–PbTiO3, PZN-PT) single crystals with the composition near the morphotropic phase boundary (MPB, 9.5%PT) are currently under investigation due to their excellent electromechanical properties. PZN-PT single crystal systems near the MPB composition exhibit strong property fluctuations. It has been shown that the piezoelectric strain constant (d33) varied from 2000 to 4000 pC/N in the same PZN-8.0%PT single crystal boule, which makes it unsuitable for many practical device applications. But the piezoelectric and dielectric property fluctuation of PZN-7.0%PT single crystal systems are much more stable than those of PZN-8.0%PT single crystal due to a little away from the morphotropic boundary composition. Thus PZN-7.0%PT single crystal


system could be the best candidate for piezoelectric transducers like medical imaging, active vibration damping, and underwater sensors. It is very clear that the inherent domain arrangement and movement in PZN-7%PT contribute to their finally stable electromechanical performance. Thus it is very necessary and important to visualize submiro-, even nanoscale domain structures and interpret their relationships with stable properties in PZN-7%PT crystals. Fig.4.22(a) shows polarized optical microscope image for (001)-oriented PZN-7%PT single crystal, which reveals rhombohedral macrosized domains with 100-150 μm in width. Inside the macrodomain plates, submicro-, even nanoscale domains were found in the various length scale of piezoresponse image of PZT-7% PT crystals, as shown in Fig.4.22 (b, c). The imaging contrast is determined by the

Fig.4.22 The polarization light image (a) and the piezoresponse image of ferroelectric domain in (001)-oriented PZN-7%PT single crystal, (b) 20×20 μm2 and (c) 8×8 μm2 scanning area ( See Color Picture in Appendix)


out-of-plane component of piezoresponse signal, in which black and white areas correspond to the opposite polarization orientation. It must be noted that the corresponding topography is not shown in this figure due to only the scratch features without any other information in the (001)-oriented crystal surface. The domain configurations in Fig.4.22 (b) seems discontinuous, irregular pattern, but in fact demonstrates some regular features with alternative bright and dark contrasts regions. This behavior is clearer in the magnification image of Fig.4.22 (c), and the regular domain is averaged 2 μm in width. It is very interesting to note that many contrast inclusion are enclosed in the seemingly same contrast matrix. i.e. much finer “bright” inclusions in the dark contrast areas and vice versa. Such domain phenomenon is typical in PZN-7%PT single crystal. Fig.4.23 illustrates this kind of

Fig.4.23

The piezoresponse image of ferroelectric domain in (001)-oriented PZN-7%PT single crystal. (a) 15×15 μm2 ; (b) 10×10 μm2 and (c) 1×1 μm2 scanning area

4. 5

Dynamic Behavior of Nanoscale Domain Structure

domain arrangement in another scanning area of the same single crystal. The average size of finer domains is about 70 nm as shown in Fig.4.23 (b). The complicated domain configurations in Fig.4.22 and Fig.4.23 reflect the relaxor nature in PZN-7%PT single crystal. It is well known that ferroelectric state in the rhombohedral ferroelectrics (like PZN) is governed by the random internal field arising from the structural disorder, compositional fluctuation and other defects, which prevents long-scale ferroelectric order during zero-field cooling (ZFC). Although normal ferroelectrics (like tetragonal PbTiO3 phase) incorporation facilitates the long-range ferroelectric order, the short-range ferroelectric order still remains with manifestation of nanoscale polar clusters in the zero-field-cooled specimen. In our present study, the polar clusters size in PZN-7%PT single crystal is larger than that of PZN-4.5%PT crystal. The latter is down to 20 nm according to Bdikin et al’s reports. The underlying mechanism for difference of domain regularity and polar cluster size between the PZN-7%PT and PZN-4.5%PT is possibly related to the role of titanium ion in the polar clusters interaction. It is expected that highly polarizable Ti4+ ions will play an important role in promoting interaction between clusters. Thus more titanium ions in PZN-7%PT would participate in facilitating long-range ferroelectric order formation. Such process undoubtedly accelerates growth and coalescence of polar clusters into microsized domains upon cooling. As a result, ferroelectric domains could be more regular at micro-, even nanometer scale in the PZN-7%PT crystal than in the PZN-4.5%PT crystal, which makes its macro physical property more stable than any other PZN-PT composition near MPB (Zeng, Shimamura, Villora, et al, 2007).

4. 5


Physical properties of ferroelectric materials are closely related with dynamic behaviors of ferroelectric domains under various applied fields,such as electric,stress and temperature fields.Piezorespone force microscopy provides a vital tool to investigate dynamic behavior of ferroelectric domains at nanometer scale.

4. 5. 1 Domain writing Ferroelectric thin films are promising candidate materials for high-density ferroelectric memory devices. In ferroelectric memory devices, information unit (one bit) could be defined as two state of “+”and “–”, which corresponds to spontaneous polarization and its reversed polarization. Thus a single domain could be deemed as one bit of data unit, and data density increases as domain sizes decrease. SFM tip could generate an intensive local electric field on specimen


surface under quite low voltage, so that nanoscale ferroelectric domains could be induced. As reported by C H Ahn et al, 50 nm-domain was obtained by applying impulse voltage to PZT53/47-RuO2 material of 0.18μm in thickness with biased SFM tip (Ahn, Tybell, Antognazza, et al, 1997), and domain size of 50 nm corresponds to a data storage density of 40 Gbit/cm2. The reversed domain size is related with impulse width and pulse time. A Gruverman et al achieved partial reversion of domains within one single grain on PZT20/80-LSCO film by altering SFM pulse width to realize domain reversion smaller than 30 nm (Gruverman, Auciello, Tokumoto, 1998). Nanoscale ferroelectric domain reversion by SFM tip bias could not only facilitate further study on physical origins of nanoscale ferroelectric domains, but also pioneer the near-field, SFM-based high-density non-optical data storage technologies.

4. 5. 2 Domain nucleation and reversal Actually responses of ferroelectric domains under applied fields are the foundation of ferroelectric material performance and their applications. Among these responses, ferroelectric domain nucleation and reversal under applied fields are exceptionally significant. A typical process of polarization and reversal including following stages: nucleation of wedge domains near electrodes (reverse polarization) longitudinal growth of domain nucleus accompanied by little lateral growth to form 180 domain structure lateral expansion of reverse domains through lateral movement of 180 domain walls mergence of reverse domains to complete reverse polarization. Although the processes of polarization and reversal described in the classical theory have already been observed in some transparent ferroelectric crystals, the universal validity of the polarization and reversal theory has been challenged by some recent experimental results. Yin and Cao investigated the average domain orientations in multi-domain single-crystal of PZN-PT with ultrasonic methods. They found that its reversal process was finished through local polarization deflection, i.e. the deflection went through 0 71 180 in stead of 0 180 (Yin, Cao, 2001). With increasingly further study on nano-scale ferroelectric domains with piezoresponse microscopy, nucleation and reversal of nano-scale ferroelectric domains under applied fields could be further understood and be added with new meaning and significance.

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4. 5. 2. 1 Polarization reversal behavior of ferroelectric thin film In ferroelectric films, domain growth parallel to applied electric field (longitudinal growth) is usually much faster than that vertical to applied electric field (lateral growth). According to the classical theory, this process is dominantly controlled by lateral growth (Tagantsev, 1996; Chai., et al, 1997; Gopalan, Mitchell, 1998).

4. 5


However, our PFM studies proposed the opposite situation, i.e. longitudinal growth is much slower than lateral growth in ferroelectric film domains so that domains represent longitudinal growth mode (Zeng, Li, Yin, et al, 2003a; Zeng, Yin, Li, 2003). Figs. 4.24 (a), (b) and (c) show respectively topography, piezoresponse images with SFM probe applied with reversing voltage of -1.5 V. Great differences in domain contrasts could be found by comparing Fig.4.24(b) and Fig.4.24(c). For example, domain contrast in lower part of grain A has been reversed to form distinct bright and dark domain strips under the bias voltage of +1.5 V. Obviously these bright and dark domain strips correspond to regions with

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Fig.4.24 Topography image (a) and piezoresponse image of 5 5μm2 regions scanned with the tip biased at -1.5 V (b) and +1.5 V (c) in PZT thin film


opposite polarization, and the reversed domains are about 70 nm in width. But the upper part of grain A shows no obvious change in domain contrast during reversal because of strong pining effect on domain reversal. Similar regions without reversal could also be found in other grains. The pinning effect is mainly resulted from defect pining and space charge effect in grains. It is obvious that SFM could not only be used to observe the inhibition on domain reversal due to pinning effect but also provide experimental basis for understanding the performance degradation of ferroelectric film during polarization reversal. Another obvious difference between Fig.4.18(b) and Fig.4.18(c) is the step like structure for domains with sizes reaching 30 nm shown in Fig.4.18(c). The step like domain structure reflects the polarization reversal of ferroelectric film. The equidistant step like domain structure indicates that domain lateral growth is faster than longitudinal so that each step width is constant during longitudinal growth, which is exactly the feature of longitudinal growth proposed by Kuroda. Thus longitudinal growth is predominant during reversal of PZT film with (111) preferred orientation. Ferroelectric domain longitudinal growth is the decisive process in (111) preferred orientation PZT thin film, and it controls the domain nucleation and growth during reversal of the oriented film (Hong, Setter, 2002). 4. 5. 2. 2 Sidewise domain growth dynamics Fig.4.25 shows the piezoresponse image of an array of nine domains reversed by applying negative 190 V of various pulse durations in the range from 40 to 120 s. The hexagonal, micrometer size domains are in good agreement with the as-grown isolated domain, and demonstrates strong stability during two weeks. The reversed domains strongly depend on the amplitude or duration of the applied voltage. No domain was found to switch when the pulse duration is shorter than 1 s. The smallest stable domain written by AFM tip bias is about 200 nm (by a 150 V, 1.5s pulse). This is likely the minimum size of the inverted domain by the AFM tip field. According to M Molotskii (2003), the domain radius under the inhomogeneous AFM tip field depends significantly on the spontaneous polarization value (P s). The equilibrium radius r m strongly decreases with P s increase in the form of rm ∝ 1 Ps2 / 3 . Thus nanoscale domains (below 100 nm) for Barium

magnesium fluoride ( BMF) crystal under AFM field should be very difficult to be expected due to its smaller P s value of 7.7 μ C/cm 2 , not like nanodomain engineering in the case of LiNbO 3 (P s =70 μ C/cm 2 ) and RbTiOPO 4 (P s 30 μ C/cm 2 ).

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Fig.4.25 The piezoresponse image of ferroelectric domains fabricated by -190 V of various pulse durations for ferroelectric BaMgF4 crystal

Fig.4.26 shows the dependence of the domain radius on the pulse magnitude for 30 s pulse duration. Each data point was calculated as an average over three domains written at different places on the sample. The radius initially increases with applied voltage up to -160 V, and then the lateral size reaches the equilibrium state and keeps nearly constant, which is similar to that in bulk lithium niobate crystals exhibiting a power law voltage dependence of domain size with subsequent saturation. The domain saturation might be due to the suppression of the polarization reversal arising from the bulk screening in which redistribution of space charges or reorientation of dipolar defects compensate depolarization field. Such case was earlier reported in ferroelectric thin film (Colla, Taylor, Tagantsev, et al, 1998). It is interesting to note that the domain radius in the pulse voltage of -175 V is over 13 μm (Fig.4.26), far larger than those at the other voltages. Such unusual phenomenon should not be ascribed to the nonequilibrium domain transition to the equilibrium state with sharp increase of domain size due to the same applying pulse time for individual pulse voltage. This behaviour is probably related exclusively to the anisotropic domain wall movement, causing the domain wall lateral propagation along the crystallographic b-axis. Such domain wall motion is governed by the processing of its pinning-depinning from defects. The defect pinning centres serve as strong energy barriers for the domain lateral growth, and effectively inhibit domain wall move freely, thus give rise to asymmetry domain wall movement. According to T Tybell (2002) and P Paruch et al (2006), the defects will induce a spatial varying pinning potential and the


random field behaviour in the system. As a result, the domain wall motion demonstrates the character of creep behaviour in the bulk BMF crystal under the inhomogeneous AFM-tip fields, just similar to that of PZT thin film (Tybell, Paruch, Giamarchi, 2002). However, the origin of the domain lateral overgrowth behaviour, occurring at tip voltage of -175 V but not in the case of the other tip voltage above -175 V, is not clear and additional experiments are necessary to understand this abnormal phenomenon.

Fig.4.26

Domain radius vs pulse magnitude for 30s pulse durations for BaMgF4 crystal

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The time dependence of the domain radius for the same pulse voltage (Vtip= 190 V) is shown in Fig.4.27(a). The domain lateral size firstly increase (t1000 and good time stability. Such material could be used for thickness and tangent frequency devices. 3. Lead free power piezoelectric material with Curie temperature above 300 and large piezoelectric coefficient. 4. Process technology in manufacturing textured lead free piezoelectric ceramic. Polycrystalline materials with highly oriented grains can be obtained by template grain growth(Seabaugh, 1997) method and added crystal seeds, and they display strong piezoelectricity in a certain crystallographic direction.

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6. 5. 5 Thermoelectric ceramics For the purpose of environment protection and energy saving, more and more attention has been paid how to effectively utilize the industrial heat which would otherwise be wasted. Thermoelectric is a promising materials, which can turn heat into electricity, generating electricity clean, no waste produced, no noise and high stability. In recent years, the rapid development of communication and information technologies requires its components and devices to continuously shrink in their sizes, and these shrinking components and devices continuously set higher requirements on working temperature. Without rotary parts, thermoelectric cooling components have the advantages of small dimensions, long life time, no noise, no pollution and more precise in temperature control, and are promising material used for high precision local temperature control. The key in thermoelectric technology lies in the search for excellent thermoelectric

6. 5

Future Development of Functional Ceramics

materials. The materials found in the past have the disadvantage of instability, e.g. easy evaporation, surface oxidation and easy decomposition and melting at high temperatures. However, metal oxide thermoelectric materials are more stable. Performance of thermoelectric materials could be characterized with merit Z value expressed as followed: Z = S 2 / ρκ Where, S is thermoelectric power;ρ is electric resistance; κ is heat conductivity. Good thermoelectric materials should possess large S, small ρ and small κ, and resulted high Z. The efficiency of thermoelectric transformation is characterized by ZT where T is temperature. It is desired to have ZT above 1. In 1997, in superconductor related research, metal oxide NaCoO4 was found to have excellent stability and thermoelectric capability that was even better than the typical thermoelectric material B2Te3, the representative thermoelectric material at that time. In the past 3 to 5 years, more metal oxide material have been discovered, such as NaCo2O4, [CaCoO3]xCoO2, [Bi2Sr2O4]xCoO2, Ni1-xLixO-Ba1-xSrxPbO3, Bi2.3-xPbxSr2.6Co2Oy etc. Besides these oxides, CoSb3 and Zn4Sb3 have also been discovered. In the research on thermoelectric materials, advanced materials with superlattice have been further studied. The strong electron correlation within these material structures is the main reason for the remarkable thermoelectric effect (Takahata, Iguchi, Tanaka, et al, 2000; Ando, Miyamoto, 1999). The low dimensional quantum effect of nano structure is another approach to increase thermoelectric properties because density of states near Fermi level will be increased so as to improve conductivity and Seebeck coefficient. Besides, the interface phonon scattering effect strengthened by nano effect will decrease thermal conductivity of the system. All these factors contribute to the improvement of thermoelectric properties. In recent years, nano wires have become a hot spot in the research of thermoelectric materials. Future research directions for thermoelectric materials mainly include: 1. Preparation of high strength Bi2Te3 based materials and fabrication of microdevices. At present, the widely used Bi2Te3 based materials have low strength and poor machinability, so that they can not meet the requirement of producing miniaturized components; 2. Novel superlattice thermoelectric film. The research will promote the miniaturization of high performance and intelligence devices. 3. Nano wire thermoelectric materials. Research on preparation process and characterization methods of nano wire thermoelectric materials and its application in nano cooler. 4. Except for single crystal production, secondary sintering is required for thermoelectric ceramics to prevent Na vaporizing. Sometimes, oriented grain preparation, spark plasma or hot press sintering are used to improve

6 Review and Prospect of Functional Ceramics

thermoelectric performance. It should be pointed out that electricity generation with waste heat is not far away anymore. Actually sample thermoelectric devices have already been fabricated, e.g. a radio with power supply from a heater, and a watch driven by human body heat (Kishi, 2000).Generating energy from waste heat has significant importance. Taking mobile heat source of automobiles as an example, the potential thermoelectric electricity generation could reach a significant value, e.g. hundreds of thousands watt. Thus the application of low cost, safe and light oxide based thermoelectric materials will be the hope for alternative energy resources in the 21st century. It has been suggested that where there’s waste heat, there’s thermoelectricity, which also indicates the promising applications of the materials.

6. 5. 6 Functional ceramic films Information technology had developed at an unprecedented speed during the late half of the 20th century. The invention of transistor in 1940 and the appearance of IC technology in 1960 have brought revolutionary changes to every aspects of human society. It is globally accepted that micro-electro-mechanical systems (MEMS) is the key technology to combine micro-mechanic sensors and actuators (executive component) with traditional IC. The technology will be the core of technology revolution in the first half of the 21st century, while sensors and actuators will be the core of the technology revolution. Materials for MEMS components include Si, SiO2, Si3O4, polymer, ferroelectric ceramic, shape memory alloy and chemically sensitive material etc. Ferroelectric ceramics are ideal materials for micro sensors and actuators in MEMS system because of its excellent piezoelectric and pyoelectric effects. Besides, these ferroelectric materials have strengths of low loss, low noise and large energy density. Due to its large sizes and high working voltage, bulk materials can not satisfy the demand in continuously miniaturized components. Therefore, there is urgent need to develop film materials to meet such requirement. Hereinto ceramic films with thickness over 1~10 μm or even thicker have become a hot topic in recent research. Compared with thin films, thick films can improve actuator’s force output per unit volume, reduce noises of sensors; while compared with bulk materials, thick films have the merits of lower working voltage, lightness, small volume, low cost and good compatibilities with VLSI technology. Furthermore, with the rapid development of large-scale integrated circuits, component size is getting smaller and operating speed is getting higher. To prevent leakage current of CMOS component and reduce time constant of linkage among circuits ( “parasitic capacitance between interconnects” ), the development of high K and low K ceramic dielectric films has become a key technical subject in the

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semiconductor industry. In the processing research on dynamic random access memories (DRAM), the main objective is to reduce the memory cell size and to increase its storage charge. To achieve larger storage efficiency, area of dielectric layers in capacitor should be enlarged while its thickness should be reduced. Therefore high dielectric coefficient thin film or ferroelectric materials could be used for memory devices. For example, high ε BST (BaSrTiO3) material could be used for DRAM, PZT and SBT(SrBi2Ta2O9) could be used for NVRAM (non-volatile random access memories). Multilayer ferroelectric thin films or superlattice materials are often applied to reinforce the spontaneous polarization. High dielectricity thin films or ferroelectric films are used for their switching properties, the control of grain size, and effects of grain boundary and other interface boundaries are key factors and need to be fully understood. The property of thin films differ from their bulk material of the same composition due to three major aspects: interface space charge, lattice misfit and huge surface/volume. In recent years, great efforts have been made in the research of the electrical (due to surface and interface space charge) and mechanical (due to lattice misfit and thermal expansion difference between film and substrate) effects on film properties in order to gain instructive principles to improve film performance. Important material systems include: PZT, bismuth containing layered perovskite SBT (e.g. SrBi2Ta2O9, SrBi2NbTaO9, SrBi4Ta4O15) and ion-replacement materials such as BLT (La0.75Bi3.25Ti3O12) etc(Park, Kang, Bu, 1999). If platinum electrodes are used for PZT devices, Pr will be dramatically decreased to cause severe aging after several times of switching. This problem could be solved with oxide conductor as electrodes or doping into PZT ceramics. However, SBT films show no sign of aging after 1212 cycles of switching because the self-regulating structure could compensate the negative effect of space charge. Besides the SBT films could avoid switching aging(Paz, et al, 1995), the material also has the advantage of low leakage current, but it requires high deposition temperature (750~800 ). While BLT material has attracted attentions, because it could be prepared at a relatively low temperature (650~700 ), and it has a high Pr but almost no aging. For manufacture of microelectronic components, integrated ferroelectric films and high ε materials could be directly deposited on silicon surface in nano scale to replace conventional gate oxide SiO2. SiO2 has been used as gate dielectric materials for many years for silicon based integrated circuits. However, as recent development of IC industry requires multifunctional IC devices, higher circuit density, and smaller devise size, so that the requirement for gate dielectric thickness has been greatly reduced. Thus a resulted problem associated with the shrinkage is the leakage current becomes larger. For example, when the thickness of SiO2 layer is decreased to 1 nm, the leakage current could reach 100 A/cm2 so that many problems including energy loss and power rise would occur. In order to

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meet the demands from development of future CMOS (complementary metal oxide semiconductor) components, materials with higher ε, high energy barrier, and high stability when deposited on Silicon are developed to replace the conventional SiO2 as dielectric material. Moreover, when integrated into devices, the materials should satisfy various requirements of integration, interface property, compatibility and reliability(Wallace,Wilk, 2002). At present, many efforts have been focused on ZrO2(Ramanathan, 2002) and HfO2 based materials. For example, HfO2/HfSixOy material has shown outstanding advantages as gate dielectric material to replace SiO2; the field effect transistors with ZrO2 as gate dielectric material have been fabricated and their size is quite small (0.5μm 0.5μm). Piezoelectric thin films could serve as the actuating source for MEMS (micro-electro-mechanical system), and commonly used films are of microns in thickness and multilayer films. MEMS are usually applied in sensors and actuators, e.g. film sensor, acceleration sensor and micro-motor. Multilayer piezoelectric devices, e.g. ultrasonic motor or actuator, have been rapidly developed in applications. For example, operating forceps with various freedom degrees could be manufactured with the actuation of ultrasonic motor(Takemura, 2003). The device can accomplish complex operations with endoscopes through a tiny cut (less than 10 mm wide), so it is beneficial for the recovery of patients and could satisfy the miniaturization development of medical treatment. Besides, the device can also be used for micro-operations such as cell separation and categorizing. It can also be used for making micro-pump. Multilayer piezoelectric components can be used for sound insulation board and vibration control system for skyscrapers. PZT ceramics are the major piezoelectric materials, but relaxed ferroelectric material (PbMg1/3Nb2/3-PbTiO3 system) can also be considered as a good alternative. Surface acoustic wave devices are usually made of bulk single crystal, but might be replaced by ZnO and AlN films. Development in thin film process technologies: Scientists and researchers from various disciplines have joined to develop ceramic thin films and have made numerous achievements. For example, in order to meet the demands of miniaturization, advanced functional films have been developed, such as high dielectric constant gate oxide film, ferroelectric film and superlattice material applied for FeRAM (ferroelectric random access memories). Rapid progresses have been achieved in thin film technology and refined processes in recent years. The massive research in superconductive oxide in the 1980’s and the research on DRAM and FeRAM after 1990 had greatly promoted the rapid development of film preparing techniques and equipments. At present, sputtering and chemical vapor deposition are the most often used methods for film fabrication. In the future, MOCVD will be an important thin film process because of its excellent step coverage, essential for three-dimensional devices. Currently there is still a lack of equipments qualified for high quality film production. So manufacturers of

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materials, devices and equipments need to closely cooperate to develop relevant equipments. Although film materials are made by variously process technologies, they have some common characteristics that are different from bulk materials. Due to the planar restriction from substrate, there’s internal stress along film surface but no internal stress along thickness direction; due to the difference between fabricating temperature and working temperature, there’s residual stress caused by difference in thermal expansion; due to different lattice parameters between substrate and film, there’s interface stress; in addition, films are thermodynamically unstable because of a large specific surface. However, different film process technologies will result in different internal stress status leading to different film performance. Effects of substrates on films: Films and bulk material are very much different. For example, in order to obtain semiconductor property, the amount of niobium dopant in BaTiO3 film is 10 times that in bulk material (Kagano, 1998). Ba0.5Sr0.5TiO3 film via epitaxial growth on SrTiO3 substrate is under strong compressive stress in the planar direction. The film is ferroelectric at room temperature due to a two-dimension stress between film and substrate, while bulk material of the same composition is paraelectric at room temperature. It has also been proved that BaTiO3 film remains tetragonal system above Curie temperature because of space charge effect. A buffer layer could be introduced between PZT dielectric film and substrate to adjust the internal stress. By introducing crystal seeds to the substrate surface, crystallization temperature can be reduced and film microstructure of oriented grains could be obtained. Techniques of molecular beam epitaxy (MBE) and pulse laser deposition (PLD) can be used to achieve films with atoms regularly arrayed on substrate surface. Nowadays these film process technologies are often used. In the past, lattice mismatch was discussed to explain film properties. Actually, comprehensive considerations should be taken to include, i.e. crystal structure, ion arrangement, attraction or repulsion of substrate atom or ion on films. A buffer layer between film and substrate or template crystal seeds could be used to adjust grain orientation and internal stress. In a report in the issue of November 2004 of American Science magazine, 13 scientists from US, Germany and China declared jointly a new finding, i.e. with proper substrate and film thickness, film strain could be adjusted so that Curie temperature (TC) of BaTiO3 film could be increased from 120 to 400~540 and its remnant polarization Pt could be enhanced to 50~70 μC/cm2. This discovery offers a promising future for lead free system ferroelectric film in applications of data storage, memory and photoelectric components. The finding is a momentous achievement of ferroelectric film materials in the past decade. Noteworthy research directions in the future mainly include: 1. Processing techniques for fabricating thick ceramic films. Physical state evolution

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and establishment of stresses during film processing need to be investigated; problems of film fracture, property none-uniformity and stress control should be solved; 2. Relationship between processing techniques of thick ceramic films and microstructure, spontaneous polarization, dielectricity and piezoelectricity etc.; 3. Research on the electron-mechanical and optical-mechanical features and displacement and force output under effects of light or electric field; 4. Performance degradation of ceramic films under effects of DC, AC current and light and corresponding mechanisms; 5. Design and fabrication of thick ceramic film devices; 6. Design, fabrication and application of high K (>12) and low K (

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