rare earth permanent-magnet alloysâ€™ high temperature phase transformation: in situ and...

Shuming Pan Rare Earth Permanent-Magnet Alloys’ High Temperature Phase Transformation In Situ and Dynamic Observation and Its Application in Material Design

Shuming Pan

Rare Earth Permanent- Magnet Alloys’ High Temperature Phase Transformation In Situ and Dynamic Observation and Its Application in Material Design

With 157 figures

Author Shuming Pan General Research Institute for Nonferrous Metal, No.2, Xinjiekouwai Street, Haidian District, Beijing, 100088, China

Based on an originanal Chinese edition: (Xitu Yongci Hejin Gaowen Xiangbian Jiqi Yingyong),

Metallurgical Industry Press, 2013.

ISBN 978-3-642-36387-0 ISBN 978-3-642-36388-7 (eBook) Springer Heidelberg Dordrecht London New York

Library of Congress Control Number:2013930302

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© Metallurgical Industry Press, Beijing and Springer-Verlag Berlin Heidelberg 2013

Foreword

Foreword

Material is one of the three primary mainstenance of the modern civilization and new material is taken for the bases and precursor of new technology revolution. Extent of production, development and application of the permanent-magnetic material is one of the indicatings for the extend of development of the contempo-rary national economy. Average family use of the permanent-magnetic material is also regarded as a measure of the live standard of the modern countries.

In resent years the requirement of the world for the rare earth permanent-magnet material has grown by 30% annually, synchronously, the kinds of this material have been developed, quality of products of the material has been im-proved and variety of the products has been extended continually in despite of impact the financial crisis on economies of many countries. Scientific and techno-logical researches and industrialization of the rare earth permanent-magnet mate-rials in China have achieved joyful result. “Rare Earth Permanent-Magnet Alloys’ High Temperature Phase Transformation: In Situ and Dynamic Observation and Its Application in Material Design” is a monograph written by Professor Shuming Pan, the General Research Institute for Nonferrous Metals, China. This book in-cludes study results of the author for more than thirty years, and researches ther-modynamics, kinetics and metallography of phase transformation related to the rare earth permanent-magnet alloys by the current solid phase transformation the-ory, especially driving force and resistant of phase transformation, homogeneous and inhomogeneous nucleation, the law of growth up of new phase, internal en-ergy change between phase transformation, the free enthalpy of phase transforma-tion, diffusion type continual phase transformation, etc. This book analyzes and discusses in depth magnetism of the first and the second generations rare earth permanent-magnet materials at 1.5K, introduces magnetism and its variation curves at –196 to 200 , and introduces the law of phase transformation of the third generation rare earth permanent-magnet alloys at temperature from room temperature to 960 and experimental video record of the whole process of high temperature phase transformation. The book also discusses manufacturing tech-nique, principle and composition of the rare earth permanent-magnet alloys, and

Foreword

the effect of above factors and on microstructure and performance of the alloy. These experiment results may be called very precious and provides important reference for study, education, production and development of the permanent-magnet materials.

China is a country with plenty rare earth resource. To transform advantage of resource into higher economic benefit it needs more efforts of scientific technical staffs to promote continual development. For this purpose Professor Shuming Pan has made his contribution. I, hereby, wish him to obtain new achievement.

I congratulate the publish of this book.

April, 2010 Changxu Shi

Preface

Rare earth permanent-magnet alloys are new function materials developed in the 1960s. With the progress of science and technology, human has stridden forward the information age. Permanent-magnet alloys have been rapidly developed. Level of their applications has been one of the measures of people’s life standards. The permanent-magnet alloys have been one of substantial bases for developing modern science and technology, such as computer, space and aviation, communi-cations, metallurgy, chemical engineering and medical protection. Meanwhile, the rare earth permanent-magnet alloy has developed gradually from one of common members to dominant role in the family of permanent-magnet alloys. By 2010, it has occupied 55% of the permanent-magnet market. The twenty first century will be the age of great developments for the rare earth permanent-magnet alloys.

The author has studied the first generation of rare earth permanent-magnet al-loys, SmCo5, the second generation Sm(Co, Cu, Fe, Zr)7.4 with high coercivity, and the third generation NdFeB series for more than twenty years. The studies include the law of phase transition from room temperature to 1000 for the rare earth permanent-magnet alloys, based on modern theory of solid phase transition. The author studied the driving force and resistant force of the phase transition, homogeneous and non-homogeneous nucleation, growth of new phase, nucleation ratio, free enthalpy, and diffusively continuous phase transition. The phase transi-tion is important in materials science. Theory of solid phase transition is a gold key to open the door of materials science: “you would not understand metal mate-rials without understanding the solid phase transition.” It is expected that the book is of benefit to science researchers, producers, teachers and students working in rare earth permanent-magnet and other materials, magnetism, metallurgy and chemical engineering. Therefore, the objectives of the book are to investigate the phase transition of rare earth permanent-magnet alloys from room temperature to high temperature and its correlation with magnetic properties and to summarize the key technology of fabrication, in order to improve the properties of the alloys, create new materials and new process, and accelerate development of the alloys.

Parts of achievements about high temperature phase transition in the book have

Preface

been reported on Chinese Science (in English and Chinese), Journal of Physics (in Chinese), Journal of Non-ferrous Metals (in Chinese), Journal of Metal (in Chi-nese), Journal of Rare Earth (in Chinese), and international journals: Journal of Applied Physics, IEEE Transaction Magnetics, and Journal of Magnetism and Magnetic Materials. The results have attracted much attention of scholars in the international conferences. Professor K. J. Strnant, pioneer in the first generation rare earth permanent-magnet alloys, attended an international conference hold in China in 1983. We showed him the image-recording of high temperature phase transition for rare earth permanent-magnet alloys. When he saw the Sm2Co7 phase separated from matrix phase and the process of new phase transition in the image-recording, Professor Strnant was excited and praised that the achievement was the advanced level in the world. He said: “I proposed that the coercivity of SmCo5 is determined by the pinning of thin layer Sm2Co7 at crystal boundary. However, Sm2Co7 was not observed directly. Now we see it in China. You should report on this new discovery in journal as soon as possible.” Later, the phase was also ob-served in American Laboratory. Professor Fidler, the famous scientist in micro-structure and magnetic properties, indicated that in situ dynamic observation for diffusively continuous phase transition was the advanced work in the world.

While writing this book, the author recalled many scholars and specialists, who collaborated with the author in the studies of rare earth permanent-magnet alloys and high temperature phase transition during more than thirty years. Here the au-thor expresses heartfelt thanks for their kind helps. They are Fengzuo Tian, An-sheng Liu, Guocheng Zhang, Jiguang Sun, Chengzhou Yu, Qiming Ying and Yujiu Liu in General Research Institute for Nonferrous Metal, Ruzhang Ma, Zuxiong Xu, Jueyun Ping and Zhengwen Li in Department of Materials Physics, Beijing University of Science and Technology, Yuefu Xiao, Shouzeng Zhou, Maocai Zhang, Zhengwen Li, Zhijun Zuo and Jianjun Tian in Department of Materials and Engineering, University of Science and Technology Beijing, and Baogen Shen, Guodong Li, Fuming Yang and Helie Luo in Institute of Physics Chinese Academy of Sciences, Yingchang Yang and Wending Zhong in Peking University, and Wei Li, Jinfang Liu and Youmei Li in General Research Institute of Iron and Steel, Zhenxi Wang, Boping Hu and Yang Luo in Sanhuan Company of Chinese Academy of Sciences, Hanming Jing in University of Jilin, Daku Sun and Guohua Chen in South-West Institute of Physics.

The author kindly appreciates Academician Changxu Shi, Chinese Academy of Engineering, making foreword in the pressing affairs and giving their support and approval to the author’s work. The author also appreciates Academician Jun Ke and Jimei Xiao, Chinese Academy of Science, Academician Dianzuo Wang, Chi-nese Academy of Engineering, Professor Ruzhang Ma, Zhengwen Li, Shengen Zhang, Engineer Wenke Li, Feng Pan, Doctor Chao Wang, Zhijun Zuo and Jian-jun Tian, who encouraged author and gave a lot of pertinent common and sugges-

Preface

tion while writing the book. Academician Guocheng Zhang, Professor Yang Luo and Shouzeng Zhou communicated with the author and proposed their viewpoints, and thus improving and richening the book. The editors of Metallurgical Industry Press, Xiaofeng Liu, Xiying Zhang and Yuan Zeng, do their conscientious work during publishing. General Research Institute for Nonferrous Metal and Metallur-gical Industry Press and Springer Press offered their supports and encouragement during writing and publishing the book. The author must express deep gratitude to above all because their contributions are involved in the book.

March, 2012 Shuming Pan

Contents

Contents

Chapter 1 Introduction ···················································································· 1

1.1 Rare Earth Permanent-magnet Alloys ······················································2 1.1.1 Rare earth ·························································································2 1.1.2 Classification and development of rare earth permanent-magnet

alloys ································································································2 1.1.3 Crystal structure of rare earth permanent-magnet alloys ···················4 1.1.4 Magnetic parameters of rare earth permanent-magnet alloys ············7 1.1.5 Criterion of permanent-magnet alloys (materials)·····························8

1.2 Principle for Alloy Phase and Phase Transformation and Growth Rule of New Phase ··················································································8

1.2.1 Phase ································································································9 1.2.2 Phase transformation ········································································9 1.2.3 Alloy·································································································9 1.2.4 Material ··························································································10 1.2.5 Alloy phase·····················································································10 1.2.6 Solid solution··················································································10 1.2.7 Exsolution precipitation··································································10 1.2.8 Thermodynamic bases for phase transformation and

classification···················································································10 1.2.9 Single crystal ··················································································17 1.2.10 Single crystal superalloy ·······························································17 1.2.11 Enthalpy························································································18 1.2.12 Entropy·························································································18 1.2.13 Latent heat of phase transformation ··············································19 1.2.14 Driving force of phase transformation ··········································19 1.2.15 Rule of growing up of new phase··················································20

1.3 Research Methods of the Magnetic Properties of Rare Earth Permanent Magnets ···············································································21

Contents

References ······································································································25

Chapter 2 The First Generation Rare Earth Permanent-magnet Alloys···· 27

2.1 High temperature Phase Transition and Magnetic Properties of SmCo5 Permanent-magnet Alloys··························································28

2.2 The in Situ and Dynamic Observation on High Temperature Phase Transformation of SmCo5 Permanent Magnetic Alloy at 25-750 ···························································································35

2.2.1 Magnetic measurement ···································································36 2.2.2 Sample preparation and experiment method ···································38 2.2.3 Influence of annealing treated specimen on coercivity····················39 2.2.4 The in situ and dynamic observation by 1000 kV HVEM under

heating condition ············································································41 2.2.5 Discussion ······················································································50

2.3 Magnetism and the in Situ and Dynamic Observation of Permanent Magnetic Alloy of SmCo5 by Annealing at 600-1000 ························54

2.3.1 Specimen preparation and experimental method·····························55 2.3.2 Analysis on chemical composition of the SmCo5 permanent

magnetic alloy ················································································55 2.3.3 Magnetic measurement ···································································55 2.3.4 Structure of magnetic domain ·························································56 2.3.5 Irreversible loss of SmCo5 permanent magnetic alloy after

annealing at 25-1000 ···································································57 2.3.6 Electronic energy spectrum experiment and analysis of SmCo5

permanent magnetic alloy ·······························································58 2.3.7 The in situ and dynamic observation on eutectoid

decomposition of SmCo5 by electronic microscope ························61 2.3.8 The in situ and dynamic observation of SmCo5 in thermal

state using transmission electronic microscope·······························62 2.3.9 The in situ and dynamic observation on SmCo5 in thermal

condition of 750-960 by Transmission Electronic Microscope······63 2.3.10 Discussion ····················································································65 2.3.11 Conclusions ··················································································67

2.4 Analysis on Variation of the Coercivity and Phase Transformation········68 2.4.1 Specimen preparation and experimental method·····························68 2.4.2 Experimental result and discussion ·················································69 2.4.3 Conclusions ····················································································74

2.5 The Optic-electronic Spectrum Study on SmCo5 Permanent

Contents

Magnetic Alloy ····················································································74 2.5.1 Specimen preparation technique and experimental condition of

optic-electronic energy spectrum ····················································75 2.5.2 Investigation on surface composition of SmCo5······························75 2.5.3 Atoms concentration variation of elements of samarium,

cobalt and oxygen from surface to depth ········································76 2.5.4 Surface compound ··········································································77 2.5.5 Conclusions ····················································································77

2.6 Analysis on Magnetic Hysteresis Loop of SmCo5 Permanent Magnetic Alloy ······················································································77

2.6.1 Specimen preparation technique, magnetic measurement and transmission microscope condition and experimental method ········78

2.6.2 Analysis on chemical composition of three kinds of specimens······78 2.6.3 Analysis on preparation technique ··················································80 2.6.4 Curve of magnetic performance and analysis at 77-550K···············80 2.6.5 Observation and analysis on specimen using TEM ·························81 2.6.6 Conclusions ····················································································82

2.7 Magnetism of SmCo5 Permanent Alloy at 1.5-523 K·····························82 2.7.1 Specimen preparation technique, magnetic measurement

apparatus and experimental method················································82 2.7.2 Magnetism measurement and curve of SmCo5 permanent

magnetic alloy at 1.5 K and 40 K····················································83 2.7.3 Measurement of demagnetization curve and value of magnetic

parameter at –196-250 by magnetic parameter measurement apparatus···································································85

2.7.4 Reversible temperature coefficient of SmCo5 at –196-250 ··········87 2.7.5 Coercivity of SmCo5 at 475-1000 ···············································87 2.7.6 Discussion ······················································································89 2.7.7 Conclusions ····················································································90

References ······································································································90

Chapter 3 The Second Generation Rare Earth Permanent-magnet Alloys ···························································································· 95

3.1 Phase Precipitation, Phase Transformation at High Temperature and Magnetism of High Coercivity Sm(Co, Cu, Fe, Zr)7.4 ···················96

3.1.1 Specimen preparation process and experimental method ················97 3.1.2 Results of specimen magnetic measurement ···································98 3.1.3 Microtexture of the alloy at room temperature································99

Contents

3.1.4 The in situ and dynamic observation on precipitation, growth up and high temperature phase transformation of cellular structure from room temperature to high temperature ···················101

3.1.5 Conclusions ·················································································· 116 3.2 Function of Zirconium on Sm(Co, Cu, Fe, Zr)7.4 Permanent Magnetic

Alloy & Observation and Analysis by Electron Microscope ················ 118 3.2.1 Specimen preparation and experimental method··························· 118 3.2.2 Research on function of Zirconium··············································· 118 3.2.3 Conclusions ··················································································121

3.3 Magnetism of High Coercivity Sm(Co, Cu, Fe, M)7.4 Permanent Magnetic Alloy at 1.5-523K·································································121

3.3.1 Preparation of specimen and magnetism measurement apparatus and measurement method ·····························································122

3.3.2 Measurement results and discussion ·············································122 3.3.3 Conclusions ··················································································125

References ····································································································126

Chapter 4 The Third Generation Rare Earth Permanent Magnet ··········· 129

4.1 Improvement of the Properties of NdFeB Permanent Magnets Due to Element Substitutions·······························································129

4.2 Magnetic Properties and the Occupancy of Co and Ga Atoms for NdFe(Co, Al, Ga)B Permanent-Magnetic Alloys ·································133

4.2.1 Preparation and method ································································133 4.2.2 Nd16Fe77-xCoxB7 alloy ···································································134 4.2.3 Nd16Co10Fe67-yAlyB7 and Nd16Co16Fe61-yAlyB7 alloys·····················136 4.2.4 Nd16Co16Fe61-xGaxB7 alloy·····························································138 4.2.5 Conclusions ··················································································144

4.3 The Studies of Main Phase Nd2Fe14B and Nd2(Fe,Co)14B in NdFeB Permanent-magnet Alloys····································································145

4.3.1 The preparation of samples and experimental methods·················145 4.3.2 SEM analysis ················································································145 4.3.3 The formation of Nd2Fe14B···························································146 4.3.4 Mössbauer spectra at room temperature········································146 4.3.5 Composition analysis and the studies of Mössbauer spectra

for Nd2(Fe, Co)14B········································································147 4.3.6 In situ and dynamic observation of TEM on Nd2Fe14B and

Nd2(Fe, Co)14B ·············································································147 4.3.7 Conclusions ··················································································150

Contents

4.4 Studies on B-rich Phase in NdFeB Alloy ·············································150 4.4.1 Preparation process and experimental method ······························151 4.4.2 The in situ and dynamic observation of Nd1.11Fe4B4 by TEM ·······151 4.4.3 Study on Nd1+�Fe4B4 by X-Ray diffraction and Mössbauer

effect·····························································································153 4.4.4 Analysis on Nd1.1Fe4B4 phase·······················································153 4.4.5 Relationship between B-rich phase and coercivity························155 4.4.6 Conclusions ··················································································156

4.5 Influence of Boron Content in NdFeB on Nd2Fe14B Phase and Magnetic Property ···············································································156

4.5.1 Specimen preparation process and experimental method ··············157 4.5.2 Influence of boron content on alloy magnetic property and

phase structure ··············································································157 4.5.3 Conclusions ··················································································161

4.6 High Curie Temperature NdFeCoGaB Permanent Magnetic Alloy ······162 4.6.1 Preparation process and experiment method ·································162 4.6.2 Using cobalt to replace part of iron···············································163 4.6.3 Use Ga to replace part of iron in NdFeCoB alloy··························163 4.6.4 Conclusions ··················································································168

4.7 Influence of Adding Element Dysprosium on Performance of NdFeB Alloy························································································168

4.7.1 Specimen preparation process and experimental method ··············169 4.7.2 Experiment result using SEM ·······················································169 4.7.3 Measurement of magnetism··························································170 4.7.4 Experiment result using transmission microscope·························171 4.7.5 Distribution of Dy2O3 ···································································171 4.7.6 Conclusions ··················································································174

4.8 Nanocrystalline Microstructure and Coercivity Mechanism Model of NdFeB Alloys with Nb and Ga ············································174

4.8.1 Experimental procedure································································175 4.8.2 Magnetic properties measuring ·····················································175 4.8.3 Study of Mössbauer effect ····························································176 4.8.4 Study of nano-microstructure of NdFeB alloys with Nb ···············178 4.8.5 Dynamic cross and microstructure of the NdFeB alloys with

Nb and Dy ····················································································178 4.8.6 Dynamic cross and microstructure of the NdFeB alloys with

Nb, Ga, Co and Dy ·······································································179 4.8.7 Curie temperature of the NdFeB alloys with Nb ···························180

Contents

4.8.8 New coercivity mechanism model of multi-component NdFeB alloys ····························································································182

4.8.9 Conclusions ··················································································183 4.9 In Situ and Dynamic Observation on Magnetic and Phase

Transformation of Nd15Fe78B7 Permanent Magnet at High Temperature ························································································184

4.9.1 Preparation process of specimen and experiment method ·············185 4.9.2 Microstructure and phase in crystal boundary of NdFeB

permanent magnet·········································································185 4.9.3 Phase transformation of microstructure of B-rich phase at

high temperature ···········································································189 4.9.4 Phase transformation of microstructure of Nd-rich filmy belt

in Nd15Fe78B7 crystal boundary at high temperature ·····················190 4.9.5 Phase transformation of Nd2Fe14B base phase of Nd15Fe78B7

alloy at high temperature ······························································193 4.9.6 Conclusions ··················································································194

4.10 In Situ and Dynamic Observation on High Temperature Phase Transformation and Magnetism of Nd16Fe77B7 Permanent Magnetic Alloy ································································195

4.10.1 Samples preparation process and experimental method ··············196 4.10.2 The in situ and dynamic observation on nanometer

microstructure and high temperature phase transformation ········196 4.10.3 Function of cobalt in NdFeCoB alloy ·········································201 4.10.4 Magnetic characteristic measurement result and analysis ···········201 4.10.5 Curie temperature measurement result ········································202 4.10.6 Phase analysis by X-ray diffraction, lattice constant and cell

volume························································································202 4.10.7 Relationship between aging temperature and coercivity

of Nd16Fe69Co8B7········································································203 4.10.8 Conclusions ················································································204

4.11 Analysis on Lamella Phase of Grain Boundary in Microstructure of NdFeB Permanent Magnetic Alloy················································204

4.11.1 Experimental method ··································································204 4.11.2 Magnetism measurement ····························································205 4.11.3 Analysis on result of the in situ and dynamic observation of

samples·······················································································205 4.12 Quick Quenched NdFeB Permanent Magnetic Alloy ·························215

4.12.1 Sample preparation technique and experimental method ············215

Contents

4.12.2 Measurement result of quick quenched magnet ··························215 4.12.3 Relationship between crystallization temperature and

coercivity ···················································································216 4.12.4 Microstructure at room temperature············································216 4.12.5 The in situ and dynamic observation on the non-crystal sample

transferring to micro-crystal by HVEM ······································216 4.12.6 Conclusions ················································································217

4.13 Stability of the Rare Earth Permanent Magnetic Alloy·······················217 4.13.1 Stability on temperature······························································217 4.13.2 Time stability··············································································224 4.13.3 Chemical stability ·······································································224 4.13.4 Conclusions ················································································224

References ····································································································225

Chapter 5 Developments and Prospect of the Rare Earth Permanent- magnet Alloys·············································································· 231

5.1 Overseas General Development···························································232 5.2 Domestic General Development ··························································238 5.3 Development Survey of Preparation Technology ·································243 5.4 Application and Expectation································································248 References ····································································································253

Appendix ········································································································· 257

Appendix 1 The Structure of Outer Electrons for Rare Earths ···················257 Appendix 2 Atomic and Ionic Radius of Rare Earths ································258 Appendix 3 Physical Properties of Rare Earths··········································259 Appendix 4 Fundamental Physical Constants ············································261 Appendix 5 Conversion of magnetic quantity between SI and

Gaussian units ········································································262

Index ··············································································································· 265

List of Figures

List of Figures

Fig. 1.1 CaCu5-type crystal lattice 5 Fig. 1.2 Th2Zn17-type rhombic crystal lattice 5 Fig. 1.3 Nd2Fe14B-compound crystal lattice and B-contained triangular

prism in Nd2Fe14B crystal lattice 6 Fig. 1.4 Demagnetization curve and magnetic energy curve 8 Fig. 1.5 Relationship of phase transformation and system free energy 11 Fig. 1.6 Variation of free enthalpy, entropy, volume and heat capacity

when second-order phase transformation occurs 16 Fig. 1.7 Gliding interface constituted by Shockley displacement 20 Fig. 2.1 The relation curve of coercivity of SmCo5 permanent magnet

alloy at room temperature vs annealing temperature 36 Fig. 2.2 The relation curves of coercivity of SmCo5 alloy vs annealing

temperature 37 Fig. 2.3 The diagram of the definition of S-factor of Doppler broadening 39 Fig. 2.4 Annealing time dependence of iHc, S and �Br/(Br+�Br), in

which annealing temperature is 750 40 Fig. 2.5 The relation of the intrinsic coercivity iHc vs annealing

temperature 40 Fig. 2.6 The microstructure of SmCo5 permanent magnet alloy at room

temperature 41 Fig. 2.7 The microstructure of SmCo5 specimen under JEM-1000

ultra-high voltage electron microscope 42 Fig. 2.8 The heating-up speed curve of the 1000kV HVEM

observation specimen 43 Fig. 2.9 Electron micrographs and electron diffraction patterns of the

precipitation process 48 Fig. 2.10 The coercivity of SmCo5 specimen after 1h annealing at

different annealing temperatures 56 Fig. 2.11 Domain structure (600×) at thermal demagnetization 56 Fig. 2.12 Domain structure (600×) after magnetization with 3T

List of Figures

magnetic field strength 57 Fig. 2.13 Domain structure (600×) of the optimized state after

magnetization with 3T magnetic field strength 57 Fig. 2.14 The curve of the magnetic irreversible loss of SmCo5 alloy after

1h heating at 25-1000 58 Fig. 2.15 The study result of the surface of SmCo5 which contains Sm

of 36.5%(wt.) by means of photoelectron energy spectrum 59 Fig. 2.16 The study result of the surface of SmCo5 which contains Sm

of 36.5%(wt.) by means of photoelectron energy spectrum 60 Fig. 2.17 The study result of the surface of SmCo5 which contains Sm

of 37%(wt.) by means of photoelectron energy spectrum 60 Fig. 2.18 The study result of the surface of SmCo5 which contains Sm

of 37%(wt.) by means of photoelectron energy spectrum 60 Fig. 2.19 The electron micrograph (the left figure) of SmCo5 specimen

at 500 and P area diffraction pattern 61 Fig. 2.20 The defect in Sm2Co17 precipitated from SmCo5 after

20min annealing at 750 (electron micrograph) 62 Fig. 2.21 The electron micrograph of SmCo5 permanent magnet alloy

after 50min annealing at 750 (light field graph) 63 Fig. 2.22 The electron micrograph of SmCo5 permanent magnet alloy

after 60min annealing at 750 (light field graph) 63 Fig. 2.23 The result of in situ and dynamic observation of the film of

SmCo5 alloy (cut perpendicular to c axis) under 1000kV ultra high voltage electron microscope (heating at 950 for 1h) electron micrograph 64

Fig. 2.24 The diagram of precipitation of Sm2Co17 from SmCo5 alloy annealing at 750-960 64

Fig. 2.25 The peak of X-ray diffraction of SmCo5 alloy heating to 950 and then quick cool down to room temperature under 1000kV ultra high voltage electron microscope 65

Fig. 2.26 The variation of coercivity of SmCo5 specimen annealing below 600 69

Fig. 2.27 The phase diagram of Sm-Co 70 Fig. 2.28 The electron micrograph of SmCo5 at room temperature 71 Fig. 2.29 The domain pattern of the specimen magnetized at room

temperature (600×) 71 Fig. 2.30 The upgrowth, coarsening, gathering and joining of the

precipitation of Sm2Co17 in SmCo5 at 420 72 Fig. 2.31 The precipitation of Sm2Co17 and Sm2Co7 in SmCo5 72 Fig. 2.32 The photoelectron energy spectrum (AES) of SmCo5 at room

temperature 76

List of Figures

Fig. 2.33 The magnetic hysteresis loop of SmCo5 specimen 79 Fig. 2.34 The magnetic hysteresis loop of SmCo5 specimen 79 Fig. 2.35 The magnetic hysteresis loop of SmCo5 specimen 79 Fig. 2.36 The variation of jHc of SmCo5 vs temperature 80 Fig. 2.37 The variation of 4�Mr of SmCo5 vs temperature 80 Fig. 2.38 Dislocation and precipitation particle in microstructure of

SmCo5 specimen (No.2) at room temperature 81 Fig. 2.39 Microstructure of SmCo5 specimen (No.3) at room

temperature, dislocation and precipitation phase 81 Fig. 2.40 Demagnetization curve of SmCo5 permanent magnet alloy

(25 ) 83 Fig. 2.41 Demagnetization curve of SmCo5 permanent magnet alloy at

temperature of 1.5K 84 Fig. 2.42 Measuring curve of SmCo5 specimen No.2 at temperature of

1.5K and 40K 85 Fig. 2.43 Demagnetization curves of SmCo5 permanent magnet alloy

at –196-250 86 Fig. 3.1 Demagnetization curves of specimen of Sm(Co,Cu,Fe,Zr)7.4

alloy 98 Fig. 3.2 Heating curve of the film surface of Sm(Co, Cu, Fe, Zr)7.4 alloy

perpendicular to c axis 98 Fig. 3.3 Cellular microstructure of Sm(Co,Cu,Fe,Zr)7.4 alloy at

room temperature 99 Fig. 3.4 Electron diffraction pattern of Sm(Co,Cu,Fe,Zr)7.4 alloy at

room temperature 99 Fig. 3.5 Microstructure of Sm(Co, Cu, Fe, Zr)7.4 alloy at room

temperature 100 Fig. 3.6 The electron micrograph of Sm(Co,Cu,Fe,Zr)7.4 alloy at

room temperature 100 Fig. 3.7 The relation of coercivity of precipitation hardening 2:17 type

alloy vs the size of the crystal cell 100 Fig. 3.8 Electron micrograph of the cellular structure of

Sm(Co,Cu,Fe,Zr)7.4 heating to 460 102 Fig. 3.9 Electron micrograph of the cellular structure of

Sm(Co, Cu, Fe, Zr)7.4 heating to 500 102 Fig. 3.10 Electron micrograph of the cellular structure of

Sm(Co, Cu, Fe, Zr)7.4 heating to 700 102 Fig. 3.11 Electron micrograph of strip structure of Sm(Co, Cu, Fe, Zr)7.4

perpendicular to c axis heating to 780 103 Fig. 3.12 Electron micrograph of strip structure of Sm(Co, Cu, Fe, Zr)7.4

perpendicular to c axis heating to 785 103

List of Figures

Fig. 3.13 Electron micrograph of strip structure of Sm(Co, Cu, Fe, Zr)7.4 perpendicular to c axis heating to 790 103

Fig. 3.14 Electron micrograph of strip structure of Sm(Co, Cu, Fe, Zr)7.4 perpendicular to c axis heating to 810 104

Fig. 3.15 Electron micrograph of strip structure at 820 104 Fig. 3.16 Electron micrograph of strip structure at 830 105 Fig. 3.17 Electron micrograph of strip structure of Sm(Co,Cu,Fe,Zr)7.4

heating at 840 for 40min 105 Fig. 3.18 The relation of peak coercivity of Sm(Co, Cu, Fe, Zr)7.4 vs

annealing temperature 106 Fig. 3.19 Demagnetization curves of Sm2(Co,Cu,Fe,Zr)17 and high

coercivity Sm2(Co,Cu,Fe,Zr)17 after different heat treatment 106 Fig. 3.20 Curve of 2:17 type Sm-Co alloy corresponding to annealing

process and curve corresponding to annealing process 107 Fig. 3.21 The phase diagram of Sm-Co-13Cu-10Fe (Cu>Fe) in

lengthwise section 108 Fig. 3.22 The diagram of the solubility of � in � which decreases when

the temperature goes down 108 Fig. 3.23 The diagram showing the forming process of the cellular

structure of 2:17 type permanent magnet alloy 109 Fig. 3.24 Electron micrograph of Sm(Co, Cu, Fe, Zr)7.4 heating at 840

for 40min 111 Fig. 3.25 Electron micrograph of Sm(Co, Cu, Fe, Zr)7.4 heating at 840




for 10min 113 Fig. 3.29 Electron micrograph of Sm(Co, Cu, Fe, Zr)7.4 alloy heating at

840 for 80min then cooling down to room temperature 113 Fig. 3.30 The relation of mHc, Hk/mHc vs Zr content of

25.5Sm-Co-6Cu-15Fe-Zr alloy 120 Fig. 3.31 The demagnetization curves of high coercivity

Sm(Co, Cu, Fe, Zr)7.4 permanent magnet alloy measured from 200 to �196 123

Fig. 3.32 Demagnetization curve of Sm(Co, Cu, Fe, Zr)7.4 at 1.5K 123 Fig. 4.1 The Mössbauer spectra of NdFeCoB and the occupation

fractions of Co atoms in the tetragonal phase Nd2(Fe, Co)14B 135 Fig. 4.2 Curves for relation of remnant magnetization Br and intrinsic

List of Figures

coercivity jHc to Al content x 137 Fig. 4.3 Demagnetization curves of Nd16Fe61-xAlxCo16B7 alloys 137 Fig. 4.4 The Mössbauer spectra of NdFeCoAlB and the occupation

fractions of Al atoms in the tetragonal phase Nd2(Fe, Co, Al)14B 138 Fig. 4.5 Occupation probability of Ga atoms via Ga content x 140 Fig. 4.6 Electron micrograph of atom crystal lattice of

Nd15Co16Fe60Ga2B7 alloy 141 Fig. 4.7 Electron micrograph of Nd15Co16Fe60Ga2B7 alloy 142 Fig. 4.8 Electron micrograph of Nd15Co16Fe61GaB7 alloy 142 Fig. 4.9 Electron micrograph of Nd15Co16Fe61GaB7 alloy 143 Fig. 4.10 Electron micrograph of Nd15Co16Fe61GaB7 alloy 143 Fig. 4.11 Electron micrograph of (Nd0.9Dy0.1)16Fe75Nb2B7 alloy at room

temperature 146 Fig. 4.12 Electron micrograph of Nd15Fe78B7 permanent magnet alloy at

room temperature and electron diffraction of B-rich phase 147 Fig. 4.13 Electron micrograph of Nd15Fe78B7 permanent magnet alloy at

500 148 Fig. 4.14 Electron micrograph of Nd15Fe70Co8B7 permanent magnet alloy

at 500 148 Fig. 4.15 Electron micrograph of Nd15Fe78B7 permanent magnet alloy at

600 149 Fig. 4.16 Electron diffraction pattern of precipitation phase in Nd2Fe14B

of Nd15Fe78B7 permanent magnet alloy at 600 149 Fig. 4.17 Electron diffraction pattern of precipitates in the matrix phase

of Nd2Fe14B in Nd15Fe78B7 permanent magnetic alloy at 600 149 Fig. 4.18 Electron micrograph of Nd15Fe78B7 permanent magnet alloy at

room temperature 151 Fig. 4.19 Electron micrograph of Nd15Fe78B7 at 322 152 Fig. 4.20 Electron micrograph of Nd15Fe78B7 alloy at 500 152 Fig. 4.21 Mössbauer spectrum of B-rich phase (Nd1.1Fe4B4) in

casted Nd15Fe78B7 alloy at room temperature 153 Fig. 4.22 The influence of boron content of Nd-Fe-B on its magnetic

characteristic 158 Fig. 4.23 The result of X-ray diffraction of casted Nd15Fe85-xBx alloy 159 Fig. 4.24 Relation of Ga atom’s probability of occupying of crystal

lattice place in tetragonal phase vs Ga content x in Nd16Co16Fe60GaB7 164

Fig. 4.25 The X-ray diffraction pattern of Nd16Co16Fe60GaB7 powder 165 Fig. 4.26 Electron micrograph of Nd16Co16Fe59Ga2B7 alloy at room

temperature 166 Fig. 4.27 Electron micrograph of Nd16Co16Fe59Ga2B7 alloy at room

List of Figures

temperature, observed Ga-rich phase (Ga2Nd) between Matrix phase and crystal boundary phase 166

Fig. 4.28 Electron micrograph of Nd16Co16Fe59Ga2B7 alloy at room temperature, observed Ga-rich phase precipitated in matrix phase Nd2Fe14B 167

Fig. 4.29 Electron micrograph of Nd16Co16Fe59Ga2B7 alloy at room temperature, observed Ga-rich phase (Ga2Nd) precipitated in Nd2Fe14B phase 167

Fig. 4.30 Electron micrograph of Nd16Co16Fe59Ga2B7 alloy at room temperature, observed block of Ga2Nd phase precipitated in Nd2Fe14B phase 168

Fig. 4.31 SEM structure of NdFeB permanent magnet alloy and SEM structure of NdFeB with addition of Dy2O3 170

Fig. 4.32 Demagnetization curves of the specimen with addition of

Dy2O3 and specimen without addition 171 Fig. 4.33 Microstructure of NdFeB prepared with hydrogen pulverization 172 Fig. 4.34 Electron micrograph of (Nd0.9Dy0.1)15Fe76Nb2B7 permanent

magnet alloy at room temperature, observed Laves phase Fe2Nb precipitated in Nd2Fe14B phase 179

Fig. 4.35 Electron micrograph of Nd15Co16Fe60Ga2B7 permanent magnet alloy at room temperature, observed Ga-rich phase (Ga2Nd), Nd2Fe14B phase and grain boundary between Nd2Fe14B crystal boundaries 180

Fig. 4.36 Electron micrograph of Nd15Fe78B7 permanent magnet alloy at room temperature, observed Nd-rich phase, Nd2Fe14B phase in the triangle grain boundary of Nd2Fe14B 181

Fig. 4.37 Electron micrograph of Nd15Fe78B7 permanent magnet alloy at room temperature, observed Nd-rich phase and favorable grain boundary after annealing in the triangle grain boundary of Nd2Fe14B 186

Fig. 4.38 Electron micrograph of Nd15Fe78B7 permanent magnet alloy at room temperature, observed Nd-rich phase in different pattern at the join point of crystal boundary of Nd2Fe14B, and observed crystal boundary of the specimen annealed not at optimized temperature 186

Fig. 4.39 Electron micrograph of Nd15Fe78B7 permanent magnet alloy at room temperature, observed clean and clear Nd2Fe14B phase and B-rich phase with concentrate stacking fault, Nd-rich phase at join point of B-rich and Nd2Fe14B 187

Fig. 4.40 Electron micrograph of Nd15Co16Fe60Ga2B7 permanent magnet alloy at room temperature, observed Ga2Nd phase, plain and

List of Figures

straight and clear grain boundary, clear Nd2Fe14B phase 188 Fig. 4.41 Electron micrograph of Nd15Co16Fe60Ga2B7 permanent magnet

alloy at room temperature, observed grain boundary with boundary angle of 120º between matrix phase of Nd2Fe14B and clean Nd2Fe14B 188

Fig. 4.42 Electron micrograph of Nd15Co15Fe61Ga2B7 permanent magnet alloy at room temperature, observed plain and straight grain boundary in Nd-Fe-Ga-B alloy with Ga addition, and observed complete and clear microstructure of Nd2Fe14B 189

Fig. 4.43 Electron micrograph of Nd15Fe78B7 permanent magnet alloy at room temperature, observed three main phases of the alloy: matrix phase Nd2Fe14B, Nd-rich phase, B-rich phase, and electron diffraction pattern of select area in matrix phase, electron diffraction pattern of select area in B-rich phase 190

Fig. 4.44 Electron micrograph of Nd15Fe78B7 permanent magnetic alloy at 322 191

Fig. 4.45 Electron micrograph of Nd15Fe78B7 permanent magnetic alloy at room temperature 191

Fig. 4.46 Electron micrograph of Nd15Fe78B7 permanent magnet alloy at 140 192


Fig. 4.48 Electron micrograph of Nd15Fe78B7 permanent magnet alloy heating at 312 for 30min 193


Fig. 4.50 Electron micrograph of Nd2(Fe, Co)14B at room temperature 197 Fig. 4.51 Electron micrograph of Nd2(Fe, Co)14B at 500 198 Fig. 4.52 Electron micrograph of Nd2(Fe, Co)14B at 700 198 Fig. 4.53 The structure of two kinds of crystal grain boundary

in Nd16Fe69Co8B7 permanent magnet alloy at room temperature 199 Fig. 4.54 Electron micrograph of triangle grain boundary in

Nd16Fe69Co8B7 permanent magnet alloy at 700 200 Fig. 4.55 Demagnetization curve of Nd15Fe69Co8B7 alloy 202 Fig. 4.56 Measurement curve of Curie temperature of Nd16Fe69Co8B7 202 Fig. 4.57 The X-ray diffraction pattern of Nd16Fe69Co8B7 203 Fig. 4.58 Relationship of coercivity of Nd16Fe69Co8B7 permanent magnet

alloy vs aging temperature 203 Fig. 4.59 bcc thin layer and impurity in the grain boundary of Nd15Fe77B8

and bcc thin layer on the boundary of matrix body 205 Fig. 4.60 bcc thin layer between Nd15Fe77B8 alloy and Nd-rich phase 206

List of Figures

Fig. 4.61 Amplified graph of Fig. 4.60: thin layer of Nd15Fe77B8 alloy 207 Fig. 4.62 The pattern of the join point of three bcc thin layers in

Nd15Fe77B8 alloy, at the join point there is Nd-rich phase 207 Fig. 4.63 Electron micrograph of Nd15Fe77B8 permanent magnet alloy at

room temperature 208 Fig. 4.64 Filmy belt between the matrix phase and inclusions of

Nd15Fe78B7 alloy observed at 280 208 Fig. 4.65 The thin layer between matrix and impurity in

Nd15Fe77B8 alloy (C) 208 Fig. 4.66 The thin layer structure in Nd16Co16Fe57Ga4B7 specimen 209 Fig. 4.67 Electron micrograph of broadened thin layer in Nd15Fe77B8

alloy observed at 450 210 Fig. 4.68 Electron micrograph of broadened thin layer in Nd15Fe77B8

alloy observed at 600 210 Fig. 4.69 Electron micrograph of Nd15Fe77B8 alloy observed at room

temperature 211 Fig. 4.70 Electron micrograph of Nd15Fe77B8 alloy at 312 211 Fig. 4.71 Electron diffraction of Nd-rich phase 212 Fig. 4.72 Photoelectron energy analysis of NdFeCoGaB 212 Fig. 4.73 B-rich phase in Nd15Fe77B8 alloy 214 Fig. 4.74 Diagram of variation of remanence of open circuit vs

temperature 219 Fig. 5.1 Flow chart of manufacture of NdFeB permanent magnet 243

List of Tables

List of Tables Table 1.1 The crystal structure of rare earth compounds (RExTMy) 7 Table 2.1 Opinions of researchers on decline of Hc of SmCo5 at 750

from 1972 to 1995 28 Table 2.2 Nucleation rate of new phase changes with time 44 Table 2.3 Chemical composition of SmCo5 permanent magnetic alloy 55 Table 2.4 Magnetic performance of SmCo5 permanent magnetic alloy 56 Table 2.5 Concentration change of elements Sm, Co and O from surface

to depth 76 Table 2.6 Analysis on chemical composition of three types of different

samples 79 Table 2.7 Magnetism of SmCo5 at 1.5K (1.49-1.55K) 84 Table 2.8 Magnetism of SmCo5 at 1.5K (1.49-1.51K) 85 Table 2.9 Magnetic performance of SmCo5 at temperature range

from –196 to 250 86 Table 2.10 Reversible temperature coefficient of SmCo5 at –196-

250 87 Table 2.11 Coercivity value of SmCo5 at 475-1000 87 Table 3.1 Relationship between demagnetization temperature and

intrinsic coercivity 101 Table 3.2 List of time needed for precipitation and growth up

(coarsening) of cell texture of Sm(Co, Cu, Fe, Zr)7.4 109 Table 3.3 Variation of lattice constant and lattice mismatch in the matrix

of 2:17 phase caused by precipitation of 1:5 phase 110 Table 3.4 Magnetic performance of Sm(Co, Cu, Fe, Zr)7.4 at 1.48-

1.53K 124 Table 3.5 Magnetic performance of Sm(Co, Cu, Fe, Zr)7.4 at 1.48-

1.55K 125 Table 4.1 Magnetic performance and Tc of alloy Nd16Fe77-xCoxB7 135 Table 4.2 Magnetic performance and Tc of alloy Nd16Co10Fe67-yAlyB7 136 Table 4.3 Magnetic performance and Tc of alloy Nd16Co16Fe61-xAlxB7 136 Table 4.4 Magnetic performance and values of Tc of alloy

List of Tables

Nd16Co16Fe61-xGaxB7 corresponding to Ga content 139 Table 4.5 Lattice constants of the matrix phase, Nd-rich phase and the

Laves phase in alloy Nd16Co16Fe61-xGaxB7 139 Table 4.6 Ga-rich phase in NdFeCoGaB 144 Table 4.7 Fe crystal sites and neighbor circumstance of Nd2Fe14B

tetragonal phase 147 Table 4.8 Relationship of magnetic performance of Nd15Fe85 - xBx alloy

and Boron content 155 Table 4.9 Relationship between magnetic performance and B content 158 Table 4.10 Quantity analysis of phase composition in NdFeB alloy 159 Table 4.11 Relationship between gallium content and lattice parameters 165 Table 4.12 Magnetic parameters of NdFeB alloy before and after

adding Dy2O3 170 Table 4.13 Probe analysis on B-rich phase of NdFeB alloy after

adding Dy2O3 172 Table 4.14 Probe analysis on the matrix phase of NdFeB alloy after

adding Dy2O3 173 Table 4.15 Neighbor atoms number of RE2Fe14B 177 Table 4.16 Occupying probability of Fe and Nb in tetragonal phases

of Nd15(Fe1-xNbx)78B7 alloys 177 Table 4.17 Magnetic properties of NdFeB with Nb and Dy 178 Table 4.18 Hyperfine field parameters of Nd2Fe12-xCo2NbxB(x=0, 0.2)

alloys 181 Table 4.19 Lattice constants of the tetragonal phase and cell volume of

alloys Nd16Fe77B7 and Nd16Fe69Co8B7 203 Table 4.20 Magnetic measuring result of quick quenched magnet 215 Table 4.21 Relationship between crystallization temperature and

coercivity 216 Table 4.22 Properties of the matrix phase 221 Table 4.23 Magnetic performance and Tc of alloy (Nd1-x Dyx)16Fe77B7 222 Table 4.24 Magnetic performance and Tc of alloy (Nd1-x Nbx)16Fe77B7 222 Table 4.25 Magnetic performance and Tc of alloy NdFeB by using Co,

Nb and Ga to substitute part of Fe 224 Table 5.1 Performance of SmCo5 permanent magnetic alloy 233 Table 5.2 Performance of the alloy Sm(Co, Cu, Fe, Zr)z (z=7-8.5) 233 Table 5.3 Magnetic energy product and density of permanent magnetic

alloy with different bonding methods 247 Table 5.4 Comparison of magnetic performance of the permanent

magnetic alloy between anisotropy and isotropy bonding 247 Table 5.5 The highest trademarks of sintered NdFeB permanent

magnetic alloy manufactured in Japan Sumitomo Metal

List of Tables

Corporate 248 Table 5.6 Demands for NdFeB alloy of industrial developed countries

as per application fields from 1990 to 1995 249 Table 5.7 Application distribution and the variation of NdFeB alloy in

China since 1988 250 Table 5.8 Yield of sintered and bond NdFeB permanent magnetic alloy

from 1985 to 2008 253 Table A.1 The structure of outer electrons for rare earths 257 Table A.2 Atomic and ionic radius of rare earths 258 Table A.3 Physical properties of rare earths 259 Table A.4 Fundamental physical constants 261 Table A.5 Conversion of magnetic quantity between SI and Gaussian

units 262

Chapter 1 Introduction

This chapter aims to make an introduction on rare earth permanent magnetic al-loys and to summarize the principles of phase and phase transformation, essential rules and the application of high temperature phase transformation of rare earth permanent magnetic alloys.

Main contents in this chapter include: rare earth elements, rare earth permanent magnetic alloy classification, crystal structure, magnetic parameter, coercivity, intrinsic coercivity, magnetic remanence, magnetic energy product, magnetic field, magnetization intensity, Curie temperature, permanent magnetic material, initial magnetization curve and judgment standard for rare earth permanent magnetic alloys. The methods used to research the magnetic performance of the rare earth permanent magnetic alloys include: the application of dynamic cross manufactur-ing process (melting, milling and sintering) of the alloy, application in alloy com-position with improving performance of the alloy, and application in research subject selection and development of new material.

The first generation of permanent magnetic alloys, such as the SmCo5, contain the feature of high anisotropy (K1 = (15-19)�103 kJ/m3, where K1 is the magneto-crystalline anisotropy constant). These alloys also have a high magnetocrystalline anisotropy field (HA = 31,840 kA/m), low temperature coefficient and high Curie temperature (Tc = 720 ).

The second generation is the 2:17 type rare earth cobalt permanent magnetic al-loy (RE2Co17). Their characteristics include a high Curie temperature (Tc = 850 ), high intrinsic saturated magnetic induction intensity and high theoretical maxi-mum magnetic energy product in comparison with the RECo5.

The third generation is the rare earth-ferrous base series (RE-Fe-B) permanent magnetic alloy. Their main advantages include a record-high magnetic energy product and richness in raw material natural resources.

There are other rare earth permanent magnetic alloys such as RE-Fe-N and nanocrystal compound exchange coupling permanent magnetic alloys in addition to the aforementioned rare earth permanent magnetic alloys.

Rare Earth Permanent-Magnet Alloys’ High Temperature Phase Transformation

1.1 Rare Earth Permanent-magnet Alloys

1.1.1 Rare earth

Rare earths are elements with atomic numbers between 57 to 71 of the third sub-group in the Element Periodic Table: La (57), Ce (58), Pr (59), Nd (60), Pm (61), Sm (62), Eu (63), Gd (64), Tb (65), Dy (66), Ho (67), Er (68), Tm (69), Yb (70) and Lu (71), additive Sc (21) and Y (39) with similar chemical properties and electronic structure.

The light rare earths are La, Ce, Pr, Nd, Pm, Sm, Eu, and the heavy rare earths are Ga, Tb, Dy, Ho, Er, Tm, Yb and Lu. The rare earth permanent-magnet alloys are known as rare earth permanent-magnet materials. Their hard magnetic proper-ties have the origin in some particular metallic compound formed from rare earth and 3d transition metals with single and multi-phase structures (Xu, 1995).

1.1.2 Classification and development of rare earth permanent-magnet alloys

Based on composition, rare earth permanent-magnet alloys can be classified into three types: rare earth cobalt permanent-magnet alloys, rare earth iron permanent-magnet alloys, rare earth iron nitride and iron carbide permanent-magnet alloys. Based on development generation, rare earth permanent-magnet alloys can be classified into three generations:

The first generation, 1:5 type rare earth cobalt permanent-magnet alloys, is represented by SmCo5 with excellent magnetic properties. Later, PrCo5, (Sm, Pr)Co5, MMCo5 (MM is the mixture of rare earths) have been developed.

SmCo5 is divided into three types, based on their magnetic properties: (1) high coercivity (Hc), linear B-H demagnetizing curve and almost the same coercivity and remanence (Hc-Br); (2) low coercivity, non-linear B-H demagnetizing curve and coercivity smaller than remanence (Hc<Br); and (3) low temperature coeffi-cient of magnetic induce, ��%/ )�0 (Zhou, 1990).

The characteristics of SmCo5 are high magnetocrystalline anisotropy (K1=(15- 19)×103 kJ/m3), high anisotropy field (HA=31,840 kA/m), low temperature coeffi-cient (as compared to rare earth iron permanent-magnet alloys), and high Curie temperature (Tc=720 , where Tc is the Curie temperature; Tc of SmCo5 is almost doubles, as compared to 312 of Nd15Fe77B8).

The maximum magnetic energy product of SmCo5 predicted by theory can reach (BH)max=244.9 kJ/m3 (31MGs�Oe). (BH)max of SmCo5 product is 130-160 kJ/m3, which is much lower than the predicted value. The theoretical value of coercivity is Hc=31,840kA/m (400 kOe), while the practical coercivity is Hc=1592-2388 kA/m (20-30 kOe) in product and Hc=3980-4776 kA/m (50-60


kOe) in laboratory (Kumar, Das, Wettstein, 1978). The remanence of SmCo5 is Br=0.8-0.96 T (8.0-9.5 kGs).

For the second generation, 2:17 type rare earth cobalt permanent-magnet alloys (RE2Co17), the characteristics are high Curie temperature (maximum Tc=850 ), higher intrinsic saturation magnetization higher than that for RECo5, and high maximum energy product in theory ((BH)max=525.4 kJ/m3 (66MGs�Oe)). The typical alloys is Sm(Co, Cu, Fe, Zr)z (z=7-8). The alloys have cellular structure, which comprises the main phase Sm2(Co, Cu, Fe, M)17 with size of about 50 nm is surrounded by the boundary phase Sm(Co, Cu, Fe, M) with the thickness of about 10 nm. The coercivity is not related to the size of grains, but is determined by the microstructure of the two phases, namely 2:17 and 1:5 phases. Due to different domain wall energies between 2:17 and 1:5 phases, the wall of 2:17 phase is pinned by 1:5 phase, and thus leading to large coercivity, for the magnetized process or reversely magnetized process (Yu, Zhang, Li, 1997). The key to obtain large coer-civity is heating treatment technology, such as solid solution and isothermal aging.

For the third generation, rare earth iron permanent-magnet alloys, the charac-teristics are: (1) the maximum energy product sets a record, (2) raw materials is sufficient and cheap ( sufficient Fe replaces deficient Co in resource, and cheap and sufficient Nd replace expensive and relatively deficient Sm needed in the first and second generations). The magnetic properties and price is much superior to the first and second generation permanent-magnet alloys. In addition, the third generation alloy can partly replace other permanent alloys in some applications. Therefore, the third generation permanent-magnet alloys develop fast. The draw-back is lower Curie temperature as compared to the first and second generations, and poor temperature coefficient and anti-corrupt. Now, the ways to overcome these drawbacks have been found (Pan, Zhang, 1990). There are many types of RE-Fe-B series, such as ternary Nd-Fe-B, Pr-Fe-B and RE-Fe-B (RE=La, Ce, MM, ), quadruple Nd-Fe(M)-B, NdFeM1M2-B.

In addition, the permanent-magnet composites comprising rare earth iron ni-trides and nanocrystals have been developed.

For rare earth iron nitride (RE-Fe-N) permanent-magnet alloys, the characteristic is the existence of N, which is different from the first, second and third generation rare earth permanent-magnet alloys. Sm2Fe17Nx (x=0-3) is the interstitial metallic compound and is prepared by heating the powders of Sm2Fe17 in N2 gas at 450-550 . A cell of Sm2Fe17Nx has three 9e sites. After N atoms occupy the 9e sites, Sm2Fe17Nx retains the Th2Zn17 structure. However, the lattice parameters and cell volumes increase, and thus significantly modifying the magnetic properties. As compared to Sm2Fe17, the anisotropy is modified from c-plane and c-axis, the magnetocrystalline anisotropy constant K1 increases from 1.2 MJ/m3 to 8.4 MJ/m3, the anisotropy field reaches 16,800 kA/m, the saturation magnetization acquires 1.54 T, the specific magnetization increases by 78% (from 100J/ (T·kg) to 178 J/(T�kg)), and the Curie temperature is elevated 343 K (from 389 K to 782 K).


The maximum energy of Sm2Fe17Nx is (BH)max=198.9 kJ/m3 and the remanence is 1.19 T, while the coercivity is rather small, Hc=640 kA/m (Coey J M D, Sun Hong, 1990; Zhou, et al, 1992). The magnetic properties of Ga substituted Sm2(Fe0.983Ga0.017)17Nx are: Br=1.27 T (12.7 kGs), iHc=2000 kA/m (25 kOe) (Zhou, et al, 1992). The other characteristics of Sm2Fe17Nx are good stability and anti-wear. The drawback is that the alloy is only used as adhesive magnets, since Sm2Fe17Nx is decomposed above 600 .

Two phase nanocrystalline composites comprise soft and hard magnetic phases with the grain size of nanometer (Coehoorn, et al, 1988; Withanawasam, et al, 1994; Ding, Mccormick, street, 1993). The main phase can be hard or soft mag-netic property. The composites take advantages of large saturation magnetizations for soft magnetic phase and high magnetocrystalline anisotropy for hard magnetic phase. The two phases are compounded in the nanometer size, the boundary is co-lattice in crystallography, and there exists the exchange coupling in the boundary. The composites have the characteristics as same to the single ferromagnetic phase in magnetized and reversibly magnetized processes. In an applied magnetic field, the magnet moments of soft magnetic phase rotate synchronously with that of hard magnetic phase. The coercivity is determined by the strength of exchange coupling at the boundary of the two phases. The enhanced effect of remanence is found for the isotropic permanent-magnet alloys prepared. The magnetic moments of the soft magnetic phase are located in the direction of averaged moments of the hard mag-netic phase. The soft magnetic phase is about 5-10 nm in sizes. The magnetic ani-sotropy of the hard magnetic phase determines the strength of exchange interactions. For example, for a rare earth permanent-magnet alloy, in which the hard magnetic properties has its origin in the exchange coupling at the boundary between Nd2Fe14B hard magnetic phase and Fe3B soft magnetic phase, the magnetic proper-ties are: Ms=1.6 T, remanence Br=0.75 Ms, iHc=238.8 kA/m (Withanawasam, et al, 1994). In addition, other permanent-magnet composites are Sm2Co17/�-Fe, SmCo5/�-Fe, Sm2Fe17N3/�-Fe, Nd(Fe,M)12/�-Fe and Sm2Fe17N/ Fe-Co, et al.

1.1.3 Crystal structure of rare earth permanent-magnet alloys

For the first rare earth permanent-magnet alloys SmCo5, the crystal structure of magnetic phase is CaCu5 type structure with space group Pb/mmm, as shown in Fig. 1.1 (Nesbbit, 1973). For the second rare earth permanent-magnet alloys, 2:17 type rare earth permanent-magnet alloys has Th2Ni17 type structure at high tem-perature, and is modified into Th2Zn17 type structure at low temperature. Th2Ni17 belongs to the rhombohedra system in crystal structure with space group P63/mmc, as shown in Fig. 1.2. For the third generation permanent-magnet alloys, the crystal structure belongs to tetragonal system with space group P42/mnm, as shown in Fig. 1.3 (Herbst, 1984). The forth rare earth permanent-magnet alloys,


Fig. 1.1 CaCu5 -type crystal lattice

(a) Crystal lattice; (b) CaCu5-type crystal lattice unit; (c) The projection of atoms of CaCu5-type

crystal lattice on plane (0001)

Fig. 1.2 Th2Zn17-type rhombic crystal lattice


Fig. 1.3 Nd2Fe14B-compound crystal lattice (a) and B-contained triangular prism in

Nd2Fe14B crystal lattice (b)

the structure of RE2Fe17Nx is the rhombohedral and hexagonal structures, where N is in the crystal lattice as interstitial atoms.

Th2Zn17 type crystal structure is one of the basic structures in rare earth perma-nent-magnet compounds. Most of RE2Co17 and RE2Fe17 compounds have the Th2Zn17 type structure at low temperature. Th2Ni17 is the isomeric with Th2Zn17; both have the similar structures. The crystal structure of Th2Zn17 is shown in Fig. 1.2. Th2Zn17 belongs to the rhombohedral system in crystal structure with space group R3m. There are three Th2Zn17 formulae and fifty-seven atoms in a cell. Six Th (or RE) atoms occupy the c sites. Among and fifty-one Zn (or Co and Fe) atoms, nine atoms occupy the d sites, eighteen the f sites, eighteen the h sites and six the c sites. The coordinates of each atom are as follows:

1 1 2 2 1 2 2 1 2 1 2 16 : (0,0, ), ( , ,0), ( , , ), (0,0, ), ( , , ), ( , , 0)3 3 3 3 3 3 3 3 3 3 3 3

1 1 1 1 1 1 1 5 2 1 1 1 1 5 1 19 : ( ,0, ), (0, , ), ( , , ), ( , , ), ( , , ), ( , , ),2 2 2 2 2 2 2 6 3 6 3 6 6 6 6 61 1 5 2 5 5 1 5 5( , , ), ( , , ), ( , , )6 3 6 3 6 6 6 6 61 1 2 2 218 : ( ,0,0), (0, ,0), ( , ,0), ( ,0,3 3 3 3 3

c

d

f 2 2 2 20), (0, ,0), ( , , ),3 3 3 3

1 1 1 2 1 2 2 2 1 1 2 2 2 ( , ,0), ( ,0, ), (0, , ), (0, , ), ( , , ), ( ,0, ),3 3 3 3 3 3 3 3 3 3 3 3 3


1 1 2 2 1 1 1 1 1 1 2 1 2 1 (0, , ), ( , , ), ( ,0, ), ( , , ), ( , 0, ), (0, , )3 3 3 3 3 3 3 3 3 3 3 3 3 3

1 1 1 1 1 1 1 1 1 5 1 5 1 518( ) : ( , , ), ( ,0, ), (0, , ), ( , , ), ( ,0, ), (0, , ),2 2 6 2 6 2 6 2 2 6 2 6 2 65 1 5 5 2 5 1 1 5 5 1 1 5 ( , , ), ( , , ), ( , , ), ( , , ), (6 2 6 6 3 6 3 6 6 6 6 2 6

h

2 1 1 1 1, , ), ( , , ),3 2 3 6 2

1 5 1 1 1 1 2 5 1 1 5 1 1 1 1 2 5 1 ( , , ), ( , , ), ( , , ), ( , , ), ( , , ), ( , , )6 6 2 6 3 2 3 6 2 6 6 6 6 3 6 3 6 6

1 2 26( ) : 0,0,0.097 , 0,0, 0.097 , ( , , 0.097),3 3 3

1 2 2 2 1 1 2 1 1 ( , , 0.097), ( , , 0.097), ( , , 0.0973 3 3 3 3 3 3 3 3

c � �

� � � )

The crystal structure of rare earth compounds (RExTMy) is shown in Table 1.1.

Table 1.1 The crystal structure of rare earth compounds (RExTMy)

Compounds Structure Symmetry Space group

Sites of RE

atom Sites of TM atom Elements on

the TM sites

RE3TM Fe3C orh Pnma 4c, 8d 4c Co, Ni

RETM2 MgCu2 cub Fd3m 8a 16d Pt, Co, Fe, Mn, Mg, Pb, Al, Ni,

Ir RETM3 PuNi3 rh R3m 3a, 6c 3b, 6c, 18h Co, Ni, Fe

Ce2Ni7 hex P63/mmc 4f1, 4f2 2a, 4e, 4f, 6h, 12k Co, Ni RE2TM7

Gd2Co7 rh R3m 6c1, 6c23b, 6c1, 6c2, 9e,

12h Co, Ni

RETM5 CaCu5 hex Pb/mnm 1a 2c, 3g Co, Ni, Cu, Ag, Au, Zn, Pt

Th2Zn17 rh R3m 6c 6c, 9d, 18f, 18h Ni, Co, Fe, Zn RE2TM17

Th2Ni17 hex P63/mmc 2b, 2d 4f, 6g, 12j, 12k Ni, Co, Fe, Mg

RE2Fe14B tetra

Note: 1. orh orthorhombic, cub cubic, rh—rhombohedral, hex—hexagonal, tetra—tetragonal;

2. TM—Co, Ni, Fe, et al. transition metal elements.

1.1.4 Magnetic parameters of rare earth permanent-magnet al-loys

Magnetic parameters of rare earth permanent-magnet alloys include: 1. Coercivity. Coercivity is known as the strength of magnetic field at zero

magnetic density. Intrinsic coercivity, Hcm, Hci or mHc, is known as the magnetic field required at magnetization 4M to zero in the reversibly magnetized process.

2. Remanence. Br is known as the remnant magnetic density when the applied magnetic fields monotonously decrease from the magnetic saturation state of ma-terial to zero magnetic fields.


4Mr or Br is defined as remnant magnetization after the permanent-magnet al-loys are magnetized to saturation followed by removing the magnetic field. Br=4Mr in CGS unit can be obtained from 4M=B�H, and Br=0Mr in SI unit can be obtained from 0M=B�0H.

3. Magnetic energy product. (BH) represents the product of Bm and Hm, and is the energy density at the gap in magnetic field H, as shown in Fig. 1.4. In the fig-ure, a point g corresponds to maximum product of Bg and Hg, which is known as maximum magnetic energy product (BH)max.

Fig. 1.4 Demagnetization curve (2) and magnetic energy curve (1)

4. Magnetic field. 5. Magnetization M. The vector field that expresses the density of permanent or

induced magnetic dipole moments in a magnetic material. 6. Curie temperature. Ferromagnetic materials below the temperature and para-

magnetic materials above the temperature. 7. Permanent-magnet alloy. Magnetic materials with large coercivity. 8. Initial magnetization curve. Magnetizations with increasing magnetic fields

from zero for thermally demagnetized materials.

1.1.5 Criterion of permanent-magnet alloys (materials)

The criterion of permanent-magnet alloys in performances is: (1) anisotropy, (2) Curie temperature, and (3) magnetization.

The main factors whether or not permanent-magnet alloys have in great future are: (1) good magnetic properties, (2) plenty of main materials, and (3) cost and price.

1.2 Principle for Alloy Phase and Phase Transformation and Growth Rule of New Phase

The magnetic properties of rare earth permanent magnetic alloys are influenced


by multiple factors. Coercivity and Curie temperature can be improved by adding Co and Al. Sm2Co17 phase segregated from the parent phase SmCo5 during tem-perature elevating from 400 to 750 accompanied by the precipitation of Sm2Co7 phase.

Sm2Co17 and Sm2Co7 phases interact with each other during the growth process and the phase transformation is reversible when lower the temperature.

1.2.1 Phase

Phase is a region with homogeneous physical properties, and there is an interface to fence out it with other parts, a system contains multi-phase is a diphase system under a certain temperature and pressure, such as water and vapor above the water, and such as saturated saltwater, salt particle within it and vapour above the liquid which coexists and constitutes a binary tri-phase system (Shi, 1994).

1.2.2 Phase transformation

Phase transformation is change in number of phases or properties of a phase. A system making of atoms or molecules reaches equilibrium and the system will be partitioned to be one or multi reciprocally distinguished homogeneous regions under function of certain outside restricting condition. Phase transformation is a change in macro state, and reduce of free energy is a gist to judge the direction of the change. Workers in material science and technology pay more attention to change of crystalline structure regulated by atoms space alignment, and change in chemical composition, long range order, etc.; and physicists concern more about phase transformation involving change in super fluid, superconductor, magnetic order and energy grade order, etc. According to characteristic of some properties phase transformation can be divided roughly to: first-order phase transformation and second-order phase transformation as per thermodynamic function. Phase transformation can also be divided to diffusion phase transformation, non-diffusion phase transformation, and phase transformation of qualitative treating controlled or heat treating controlled as per atoms transfer during formation of new phase (Shi, 1994).

Material state can be divided to liquid state, solid state and gas state, transfor-mation among them is called as phase transformation or change of state.

1.2.3 Alloy

Chemical combination among metals forms alloy, and it is difficult to mark off between “alloy” and “metal”. To research alloy is to use alloy as material and it is usually to substitute alloy and metal material reciprocally.


1.2.4 Material

Material is a substance to be used to manufacture useful parts of construction, apparatus and goods. According to its usage material can be divided to energy source material, construction material, and electronic material and aviation mate-rial; and as per material property characteristic it can be divided to configurable material and functional material; or can be divided to metal material and non-metallic material as per chemical component and structural features.

Material is the substance base for production and lives of human, is an important sustain of civilization of human being. Progress of material depends on progress of science and technology and social productivity, whereas progress of material also promotes progress of science and technology and development of social economy.

1.2.5 Alloy phase

In alloy a phase that forms uniform crystal structure and properties through inter-action among component elements, has parts with homogeneous composition and is disjoined by specific interfaces is called alloy phase (Xiao, 2004). The alloy phases can be divided to four types: solid solution, pure monomeric component substance, order solid solution and metallic compound.

1.2.6 Solid solution

Solid material can more or less dissolve other elements without changing its struc-ture and thus forms solid solution. Features of solid solution will change along with change of its composition. Solid solution is divided to gap solid solution and re-placement solid solution as per location of solute atoms in structure of solid solution.

1.2.7 Exsolution precipitation

The maximum dissolubility of some component in solid solution is limited, the dissolubility decreases with decreasing of temperature thus produces precipitate with solute atoms as the main component. The process to form new phase struc-ture is called Exsolution precipitation.

1.2.8 Thermodynamic bases for phase transformation and classification

Phase transformation proceeds along a path of minimum resistance and develops towards direction of energy lowing. Magnitude of free energy of the system varies


as G H T S� � � � � , if G� <0, it means the system is instable thermodynami-cally, then diffusing and transfer occur spontaneously. Thus it can be seen that thermodynamics is the driving force for diffusion and structure transfer.

If considering one type of substance system this substance takes which struc-ture at a certain temperature and pressure depending on its free energy to be high or low, i.e., to realize state of minimum Gibbs free energy, Gmin. Free energy G = H – TS (T is temperature; S is entropy), enthalpy H =E + pV, where E is internal energy; p, V are pressure and volume, which in solid phase transformation of metal have little change; using the Helmholts free energy to substitute Gibbs free energy may be more convenient, free energy of anthracite can be expressed as: F = E – TS, if temperature and pressure are changed and other structure of this sub-stance has a lower free energy then its structure will change and thus transform to be other phase, thus phase transfer or phase transformation occurs. Fig. 1.5 shows that free energy of two phase changes along with change of temperature and free energy of each phase can be written as:

� � �

� � �

G H TSG H TS

� ��

Fig. 1.5 Relationship of phase transformation and system free energy

System free energy at high temperature depends on the entropy; a phase with large entropy has low free energy and is stable. The temperature at which free energy of two phases are equal is called phase transfer point or phase intersectant point T0, when temperature is higher than T0, then S� <S H�<H . At T0, G�=G

and L= H��H =T (S �S�). In process changing from high temperature to low tem-perature the structure will transform from � phase to � phase, and discharges la-tent heat L; in process changing from low temperature to high temperature the structure will transform from � phase to � phase, and absorbs heat L.

At melting point of solid material

l s m 1 2( )L H H T S S� � � �

where Hl, Hs, S1, and S2 represent enthalpy and entropy of liquid phase and solid phase, respectively, Tm represents temperature of melting point.

At phase transformation point T0, � 0 0 � 0 0( , ) ( , )G p T G p T� , when pressure of


system changes as p0 p0 �p, phase transformation temperature changes as 0

0 � .Free energy G expands by Thaler progression at (p0 T0):

22

0 0 0 0 2

2 22

2

1( , ) ( , ) d d d2

1 d d d2

pT T

p

G G GG p p T T G p T p T pp T p

G GT p Tp TT

� ��

� � � ��

� (1.1)

Considering Maxwell relation formula:

T

GVp

� ��

, p

GST

��

Thus above formula can be expressed as:

� � � � 20 0 0 0

2

1, , d d d2

1 d d d2

T

p p

VG p p T T G p T V p S T pp

S VT p TT T

� ��

� ��

(1.2)

Defined by homothermous compressible rate and isobaric expansible rate: 1

T

VxV p� ��

� � � ��

1

p

VV T

� ��

Integral of thermal capacity ratio c by considering total sum of system entropy

at various temperatures is:

0d

T cS TT

� � , and the thermal capacity ratio at a con-

stant pressure is pp

Sc TT��

, thus

� � � �0 0 0 0

2 2

, , d d

1 1d d d d2 2

p

G p p T T G p T V p S Tc

xV p T V p TT

�

� � � � � � � �

� � ��

(1.3)

Thus free energy of � and � phases can be expanded in progression, and then obtains following formula by subtracting each other:

� � � � � �

� ��

2� � � � � � � �

2� � � �

1d d d d2

1 d d d2

p p

G V V p S S T x V x V p

c cT V V p T

T� � �

� � � � � � �

��

(1.4)

From point of view of dynamics phase transformation is usually defined as per variation of system free energy, such as the nth-order differential coefficient


dd

n

nG

T and d

d

n

nG

p of free energy vary at phase transformation point, which is de-

fined as nth-order phase transformation (n-1-order coefficient series). Moore free energy function for first-order phase transformation system vs.

first-order derivative function of some variables appears discontinuous phase transformation at phase transformation point. Classified guideline of this phase transformation is suggested by Ehrenfest that let Moore free energy vs. restrictive variable derivated by stages, and its derivation function appears discontinuously the lowest-order which represents the order of phase transformation. If only con-sidering unit coefficients of two restrictive variable of temperature and intensity of pressure (p), Moore free energy (G) of each possible phase is a curve face of T-p-G space. An intersectant curve of two curve faces gives restrictive condition for equilibrium of the two phases. Moore free energy function to first-order partial derivative is (dG/dT)p = �S of temperature at a constant pressure, and Moore free energy to first-order partial derivative is (d d )TG p V� at a constant temperature, where S and V represent Moore entropy and Moore volume, respectively. When first-order phase transformation occurs variation of Moore entropy of system is discontinuous, and S H T� � � , thus when first-order phase translation occurs there is latent heat effect and variation of Moore volume is discontinuous; then there is volume effect along with first-order phase transformation which can be measured by means of dilatometer. First-order phase transformation usually goes along with larger heat stagnant, such as phase transformation from �-Fe to �-Fe (Shi, 2004).

First-order phase transformation is occurred when two one-order differential items in Eq.1.4 have transformed (limited change), in Eq.1.4 one-order differen-tial item has:

� �T

G V Vp

� ��

� �p

G S ST

��

That is, first-order differential coefficient of free enthalpy at the point of phase transformation, in the other word, in the system there is enthalpy change (absorb-ing or releasing of latent heat), and also there is volume change, such as phase transformation phenomenon like melting, vaporizing, sublimating, etc.

Above formula preserves one-order differential item as follows:

� �

� �

dd

S SpT V V

��

� (1.5)

According to definition of latent heat it can be expressed as Claypeyron equa-tion which gives relationship between temperature and intensity of pressure of a


first-order phase transformation:

� ��

dd

p LT T V V

��

(1.6)

If one of the compressible rate x, heat capacity ratio cp and heat expansible rate � changes it will be second-order phase transformation, that is, the second-order phase transformation has no change of entropy and volume at constant tempera-ture and constant pressure, but free enthalpy and second-order derivative of sys-tem represent discontinuous variation. Preserving quadric item from developed formula of Thaler progression of free energy (the primary item is 0):

2

2 2pcGTT

��

� (1.7)

2

2G xV

p�

��

(1.8)

2G V

p T��

��

(1.9)

Ehrenfest equation is derived to give relation of temperature and intensity of pressure of second-order phase transformation point:

� � � �

� � � �

dd ( )

c cpT x x TV

� ��

� ��

� � (1.10)

This kind of phase transformation mainly has: transformation of for ferromagnetic material from ferromagnet to paramagnet at Curie temperature, transfer from conductor to superconductor at fixed temperature, transfer of some alloy material from in-order to out-of-order, etc.

Free energy function of second-order transformation system to second-order derivative function of restrictive variables is discontinuous at phase transforma-tion point. When phase transformation occurs the corresponding physical quantity of this second-order derivative function appears discontinuous vitiated phase transformation. If discussing only the unit coefficient of two restrictive variables of temperature (T) and pressure (p) there may appears the 1st phase with its own free energy function (G) in the system, and this function is a curve face among (G-p-T). The geometrical mode of intersectant two curve faces is differ from the intersectant mode of first-order phase transformation, and its free energy curve representing two phases does not intersect sharply, but is tangent along a curve. At this phase transformation point the free energy of two phases is equal and the first-order derivative of the free energy function is also equal, but the second-order derivative varies discontinuously. The second-order partial derivative of

free energy function to temperature at constant pressure is 2 2( ) pG T� � �


r( )S T C T� �� , that is, the heat ratio at the phase transformation point is not

equal. At constant temperature the second-order partial derivative of free energy

function to intensity of pressure is 2 2( )G p� � � � , where � is volume com-

pressible coefficient, that is, the compressible coefficients of two phases at phase transformation point is unequal. There are some examples of second-order phase transformation that its cp does not jump limitedly at the phase transformation point but tend to infinite, and Cr exhibits entering state along with change of tem-perature and this kind of transfer is called entering phase transformation. A typical method to process second-order phase transformation is average field theory, such as Weiss theory to analyze ferromagnet transferring to paramagnet, and Dragg-Williams theory to analyze long range order. Landau suggested that all first-order phase transformation can be described by using a generalized order referent amount (�), and � will change with change of temperature and becomes zero when temperature rises to T0 (phase equilibrium point) and keeps zero above T0; once phase transformation occurs the � will decrease discontinuously to zero after reaching T0. Thus first-order phase transformation is also called as discontinuous phase transformation whereas second-order phase transformation is called as con-tinuous phase transformation (Shi, 1994).

The second-order phase transformation is differing completely from first-order phase transformation. Its characteristic is not to have variation of entropy and volume at phase transformation point at constant temperature and constant pres-sure, but its second-order derivative of free enthalpy has discontinuous variation. That is called second-order phase transformation. When the second-order phase transformation takes place variation of free enthalpy, entropy volume and heat capacity is shown in Fig. 1.6, where cp is the heat capacity at constant pressure, � is the expansion coefficient at constant pressure, and � is volume compressing coefficient at constant temperature, thus the second-order derivative of G is:

2

2p

p pp

cG G ST T T TT

� �� (1.11)

2

2T T

G G V Vp p pp

� ��

� � � �� (1.12)

2

pT p

G G V VT p T p T

��

(1.13)

Ehrenfest equation describes the relationship of temperature and intensity of pressure at this kind of phase transformation point:


( ) ( )

f i

dd ( )

f ip pc cp

T TV d d�

��

(1.14)

f i

f i

dd

d dpT

��

� (1.15)

Fig. 1.6 Variation of free enthalpy, entropy, volume and heat capacity when

second-order phase transformation occurs

When second-order phase transformation occurs chemical potentials and first-order partial derivatives of two phases are equal, but their second-order partial derivative is unequal. The other type expressing formulae are as follows (Shi, 1994; Wang, 1980):

� �

� �

� �

2 � 2 �

2 2

p p

T T

p p

T T

T T

T T

� � � � � � � �

�� ! � �

�� "� � � �� #

(1.16)

2 � 2 �

2 2

2 � 2 �T Tp p

T p T p

�� " � � � ��

!� � � ��

"� � � � � � � �� #

(1.17)

According to thermodynamics function formulae are:

2

2

2

2

2

p

pp

TT

pp

cST TT

V V VBV pp

V V VAT p V T

��

� � �� ! � �� #

(1.18)


where cp is invariable heat capacity; 1

p

VAV T

�� is expanding coefficient of

material; 1

T

VBV p� ��

� � � �� is compressing coefficient of material.

Substitute above three formulae into Eq.1.16 and Eq.1.17 and then obtain: � � � �,V V S S� �

� � � � � �, ,p pc c � �" " "

It can be seen that when second-order phase transformation occur heat effect and volume have no actual variation whereas volume of the phase and entropy has continuous variation. There are only expanding coefficient, heat capacity and compressing coefficient have discontinuous variation. But as 0, 0,pc � " � "

0,�� " as shown in Fig. 1.6. The second-order phase transformations being found mainly are: transfer to su-

perconductor state, i.e., conductor transfer to superconductor at fixed transfer tem-perature under non-magnetic field; some alloys transfer from in-order to out-of-order at the critical point; ferromagnet transfers to paramagnet at Curie tempera-ture; transferring of liquid nitrogen, that is, He liquid transfers to He liquid at entering point of temperature and pressure (such as 2.19K and 5,147.31 Pa); turn-ing of +

4NH in crystal at fixed transfer temperature. Classified in point of view of thermodynamic different thermodynamic relation formulae can be derived, which is only applicable to unitary system. Whether are third-order phase trans-formation or higher-order phase transformation found? Phase transformation of third-order or above is infrequent.

1.2.9 Single crystal

Single crystal is a crystal grown from a single nucleus. Its characteristic is that its crystallography orientation can keep consistency in all of its interior. Whereas its profile can be a regular polyhedron or an irregular arbitrary shape. Single crystal can reveal fixed anisotropic features of the crystal, which is used widely in high technology fields. Crystal material mostly means single crystal, such as photoelectric crystal, laser crystal, acousto-optic crystal, piezoelectricity crystal, magnetic crystal film, functional crystal, etc. (Shi, 2004).

1.2.10 Single crystal superalloy

Single crystal superalloy is developed on bases of amorphous in production of casting of superalloy and oriented crystallizing superalloy. Feature of the alloy


components is not to add boundary strengthening elements, has a higher melting temperature in comparison with casting and oriented crystallizing superalloy, and has a simple form. Single crystal is manufactured by special precision casting method-selective crystallization method and seed crystal method and by using a unilateral heat flow oriented crystallization furnace. Single crystal superalloy may appear as dendrite structure, cystiform structure and plane solidification structure along with different temperature gradient and crystallization speed, this kind of alloys can effectively decrease segregation and specially have homogenous ele-ments distribution in plane solidification structure. Use temperature of single crystal superalloy can be enhanced significantly in comparison with casting and oriented crystallized superalloy (Shi, 2004).

1.2.11 Enthalpy

The enthalpy of an object system is defined as

H U pV$ �

where U, p and V represent internal energy, pressure and volume, respectively. According to the 1st law of thermodynamics for a system with only acting ex-

pansion work heat Qp absorbed from outside in process of constant pressure equals to increment of enthalpy: pQ V p V H� � � � � � . Where p�V represents expansion and compressing work of system. If the system acts other works be-sides expansion work in process of status change, such as external magnetic field act magnetizing work to system, external force caused deformation work, etc., then a generalized enthalpy can be introduced:

i iH V y x$ �%

where y is intensity, such as pressure (p), magnetic inducing intensity (B), tensile stress (F), surface tensility (�), etc.; whereas x corresponds a widely extended amount, adding volume (V) total magnetic torque (VM, M) as magnetization in-tensity, displacement (L), area (A), etc. In process of constant intensity Y system absorbs heat from external equals to increment of the generalized system enthalpy in an uniform process.

d d d d d di i Yi

H V y x V W Q� � � � �%

where d di ii

W y x� % is the primary work in the process.

1.2.12 Entropy

The entropy represents a status function of a system, and is usually represented by


a common symbol S. Change of entropy of a system is defined as heat absorbed of the system from external dividing temperature T of external heat source in process of a particulate passing, that is, d dS Q T� . For a reversible process there is only an infinite small amount in difference between temperature of the system and temperature of external heat source, thus T in the formula can be sub-stituted by system temperature, whereas TdS represents quantity of heat in re-versible process. Unit of entropy is J/K. The system absorbs quantity of heat in a process is proportional to amount of substance containing in the system, thus entropy is an extensive parameter of additivity. It can be testified in statistical physics that entropy of system at one of macro-status has relation to number � of various possible micro status corresponding to that macro status, which is S=Kln�. In the formula K=1.381×10�23J/K, which is called as Boltzmann constant. The entropy represents a physical quantity of out-of-order of system state.

1.2.13 Latent heat of phase transformation

Latent heat of phase transformation is that phase transformation from the matrix phase (�) to a new phase (�) of 1 mol substance absorbs or releases heat (J/mol) within a system at constant temperature (T) and constant pressure (p). The equi-librium condition of two phases is that chemical potentials of two phases are equal.

� � � �� , ,T p T p �

Then latent heat of phase transformation can be derived as:

� �� L T S S h h h� � � � � �

where S and h represent mol entropy and mol heat enthalpy, respectively (Shi, 1994).

1.2.14 Driving force of phase transformation

New phase (� phase) of unit volume generated from the matrix phase causes de-crease of Gibbs free energy G at constant pressure and constant temperature. That is called the driving force of phase transformation. In condition that degree of super-cooling or superheating mT T T� � � is not big, then ��G� �

� �mL T T T�� , where Tm is equilibrium temperature of two phases, T is thermo-

dynamic temperature when phase transformation takes place; L is latent heat of phase transformation of unit volume. Process of phase transformation is always in direction to decrease Gibbs free energy, the more negative the ��G� is the big-ger the driving force for phase transformation will be (Shi, 1994).


1.2.15 Rule of growing up of new phase

The rules of growing up of new phase are: 1. Growing up of new phase is carried out through transfer along new phase in-

terface in solid phase transformation. 2. Growth of new phase is classified to various types, there are cooperating

type transfer, non-cooperating type transfer, continuous growing up and stage mechanism growing up, diffusion control and non-interface control.

3. Non-diffusion phase transformation has unchanged composition; cooperated transfer relies on displacement movement of interface.

4. Gliding interface constituted by Shockley displacement, stowing sequence in dense alignment lattice is ABCABC on fcc lattice of dense alignment face, and on dense alignment face of hcp lattice stowing sequence is ABABAB. Actually there is stowed layer displacement in crystals and displacement exists at the edge of layer displacement, as shown in Fig. 1.7 as follows.

Fig. 1.7 Gliding interface constituted by Shockley displacement

It can be seen from Fig. 1.7 that there is a Shockley displacement on every

other dense alignment face and a series of Shockley displacement constitutes the interface. The left side of the interface is fcc lattice and the right side is hcp lattice. The interface constituted by a series of Shockley displacement will transfer along with gliding of displacement because Shockley displacement may glide along y face (111) in direction of y (112). This type of interface is called possibly gliding interface. This result will lead to growing up of one phase and reducing of the other phase (Liu, Ren, Song, 2003).

5. Non-cooperative type transfer and growing up without change of composition. When interface between new phase and the matrix phase is a non-lattice inter-

face the interface of matrix phase is prone to contain the atoms from the matrix phase so that growing up can be carried out continuously, which is called con-tinuous growing up. And because there is no composition change in growing up of new phase thus the growth only needs short distance diffusion of atoms near the interface, thus this transfer is only controlled by interface course.


Moving rate of interface becomes slowly along with decrease of temperature and will increase along with superheated degree in heating transfer (Liu, Ren, Song, 2003).

6. Non-cooperative transfer and growing up with change of composition. Volume diffusion controlled growing up and global new phase growing up by

heat activation: growing up speed of new phase is proportional to diffusion coef-ficient of solute in the matrix phase and concentration gradient, but the bigger the difference of composition between new and old phases is the smaller the growing up speed is. When degree of super-cooling increases concentration gradient will be raised. Diffusion coefficient takes on exponential decrease along with decrease of temperature. Decrease of temperature has two contrary function, that is, en-hancing driving force and lowering diffusion ability, which leads to the maximum value on the curve of growing speed (Liu, Ren, Song, 2003).

Growing up in side of sheet new phase (Qi, 2003): growing up is not in con-stant speed, and growing up speed is related to diffusion coefficient of atom and the time that the bigger the diffusion coefficient is the faster the growing up speed is. But they do not appear a linear relationship. Growing up speed is slowly along with prolonging of time, the longer the time is the smaller the growing up speed is. That may be due to that the diffusion distance becomes far and far along with passing of time and interface controlling is bigger; if the interface of new and old phases is a coherency or semi-coherency interface and interface containing factor is very small and is hard to be moved the transfer can only be realized by step moving. Transfer of step needs long distance diffusion of solute atoms because different composition of the new phase and old phase. Difficulty for mechanism of step rests with source of step. One comes from heat activation which forms a two-dimension crystalline nucleus, the other comes from outcrop of screw dis-placement existed in new phase on the interface.

1.3 Research Methods of the Magnetic Properties of Rare Earth Permanent Magnets

There are two methods to increase the coercivity of permanent-magnet alloys. One is to seek the relationship between coercivity and physical or geometrical parameters of the alloy phase, based on certain physical model and followed by mathematic treatment. The other is to directly observe the relationship between coercivity and geometrical parameters of microstructure or physical parameters of phase. Both are supplement and complement each other. For rare earth perma-nent-magnet alloys, the coercivity is very sensitive to the microstructure of the alloys. Therefore, the observation of microstructure and the study of the relation-ship and interaction between the microstructure and magnetic domain are basis of investigating coercivity mechanism (Li, Zhu, 2003; Pan, Xiao, 1989).


After it is determine that the magnetic properties of the rare earth permanent-magnet alloys are studied based on the observation of the microstructure and its interaction with domain structure, it is necessary to consider how to achieve the objectives, using TEM, SEM or EPMA? In experiments, some or all of them can be used coordinately. The development of electron, vacuum, and computation technologies has the development and application SEM, TEM and EPMA redou-bled. The apparatus with high resolution and sensitivity lay solid foundations for studies of materials.

The author discovered in experiments that the magnetic properties of rare earth permanent-magnet alloys are influenced by many factors. For Co and Al substi-tuted NdFeB alloys, the maximum coercivity occurs at Al concentration of 4%, as Co concentration is 10%, while the maximum coercivity occurs at Al concentra-tion of 2%, as Co concentration is 16%. The substitution of Al can increase the coercivity, while the substitution of Co can elevate the Curie temperature. The combined substitution of Co and Al can increase both the coercivity and the Curie temperature of NdFeB alloy. If an appropriate amount of Nb is added, the coerciv-ity of NdFeB alloys can be greatly increased, while the Curie temperature is also elevated.

Based on the in situ and dynamic observation to SmCo5 using JEM-1000 kV super-high voltage TEM, the crystal nuclei of Sm2Co17 phase segregated from the parent phase SmCo5 at 600-750 grow, collide each other and get together under the phase transition driving force, and thus slowing down the rate of phase transi-tion. At the same time, the crystal nuclei in other regions still grow. This is the “isothermal phase transition”. After above process is finished, the new phase formed at the given temperature cannot continuously grow. It is necessary to in-crease temperatures to provide more phase transition driving force. The new nu-cleation will occur and grow in the thermal excitation. This is the “varying tem-perature phase transition”. Besides the Sm2Co17 phase, Sm2Co7 phase is segre-gated from the parent phase SmCo5 at 600-750 . There exists the interaction in growth of the Sm2Co17 and Sm2Co7 phases. As temperature is lowed, the phase transition is reversible.

The in situ observation at 400-750 to Sm(Co, Cu, Fe, Zr)7.4 permanent-magnet alloys with large coercivity indicated that the Zr-rich phase with long plate in shape appears, while the cellular microstructure (main phase Sm(Co, Cu, Fe, Zr)7.4 with the size of 50 nm is surrounded by Sm(Co, Cu, Fe, Zr)5 with the thickness of 10 nm), which has decisive effect on coercivity of the alloys, is growing. Due to different wall energy between 1:5 and 2:17 phases, the domain wall of the 2:17 phase is pinned by the 1:5 phase in the reversibly magnetized process, thus acquiring large coercivity. The coercivity of the alloys is determined by the pinning field. As the applied magnetic field is larger than pinning field, the alloys are magnetized to saturation state. The coercivity of the alloys is not de-


termined by the crystal size, but is determined by the structure of Sm(Co, Cu, Fe, Zr)7.4 with the size of 50 nm and Sm(Co, Cu, Fe, Zr)5 with the size of 10nm (Pan, Xiao, 1989; Pan, Ma, Li, 1993; Pan, Liu, Luo, 1990).

1. Application of the “dynamic crossing” in manufacturing the rare earth per-manent magnetic alloys:

The good results are obtained from high vacuum in sintering, based on the principle of high activity and easy oxidation of rare earths. For example, in order to obtain high vacuum, a small amount of highly active La and Ce is located in furnace. As La and Ce are easily compounded with oxygen, oxygen in air is de-creased.

Add Zr-Al oxygen absorber into sintering furnace to seize oxygen in air; and then add a chemically active element to reduce oxygen in the furnace so that to increase sintering efficiency and to raise product yield.

Develop dynamically adjust liquid amount for means of cross-handle of liq-uid, solid and air phases in hydro-milling to reduce the amount of air phase as more as possible, so that to achieve the target to lower oxygen content in the alloy. That had gotten good result.

In vacuum melting use mechanic pump to draw to low vacuum, and then add argon into furnace in order to drive out the remained oxygen faster; rapid the above two steps for several times. This was testified as an effective deoxydation process.

Design a process for “reproducing rejected product because of unqualified performance of rare earth magnetic material due to oxidation” to make use of waste goods and to improve environment. The surface of rare earth ferrous per-manent magnetic material is likable to be oxidized in manufacturing process, which may produce Nd2O3 and lead to degrade magnetic performance. New proc-ess, designed under the principle of “dynamical crossing and supplement benefi-cials”, is to add element which is inadequate.

Property of the rare earth metals are active, that is easy to be oxidized to be Nd2O3 which make the formation of Nd2Fe14B become difficult; supplement neo-dymium can promote the formation of Nd2Fe14B. This process is a three-dimension melting, one is heat activation, another is draw vacuum and the third is magnetic stirring. Afterwards, by passing through corresponding milling and sin-tering process the magnetism can be re-produced and the magnetic performance can be recovered. By this way the waste becomes qualified product.

2. Application of the “dynamic crossing” in developing the rare earth perma-nent magnetic alloys and raising performance of the alloys:

There are many milling techniques in product process of the rare earth per-manent magnetic alloys, among them by hydrogenated crush and then mill to 3- 4m is an effective method to reduce oxygen content of the alloys. However, a severe problem occurred in the hydrogenated milling process: crystal size of the


powder by hydrogenated milling after sintering was found to be extremely big in test, that the granule size of crystal of permanent magnetic alloy, sintered at 1000 , is as large to 400nm, which will degrade the coercivity. To resolve this problem a new project was designed-add element dysprosium to make granule of alloy fined. Element Dy is mainly distributed crystal boundary of alloy microtexture so that increase of Dy can make granule become finer. Analysis by X-ray indicates that when Dy content is raised the lattice constant a, c are decreased. For Nd15Fe77B8 alloy its lattice constant a was reduced from 0.883nm to 0.870nm, and c was re-duced from 1.226nm to 1.211nm by adding 1% (at.) Dy replaces a few Nd.

Studies indicated that when add dysprosium in 0.4% (at.) the average size of granule was 7.0 m; and when add Dy in 2% (at.) the average size of granule was 5.1m; when add Dy to 3%(at.) the average size of granule reduced to 3.0m. The result of magnetic measurement showed that the coercivity of the rear earth permanent magnetic alloy was enhanced obviously.

Analyze functions of elements for the rear earth ferrous permanent magnetic alloys of ternary system or above, and analyze problem existed in material prepa-ration process in order to improve performance of material.

If the relation between the coercivity and the manufacturing process for the rear earth permanent magnetic alloy is described in functional formulate, or say, the coercivity is the function of (element component, melting, milling, forming, sin-tering, and magnetization). Thus to enhance coercivity it is necessary to opti-mize every variables in manufacturing process. For example, alloy elements from ternary change to seven, the melting process and heat treatment technique of the alloy should be changed.

For example, texture solidification status of NdFeB permanent magnetic alloy in melting process, including factors such as columnar crystals, degree of orienta-tion, size of crystal granule, etc., has important influence on the magnetism of the alloy. A desired crystal texture of equiaxial and fine crystal granule is interrelated to melting temperature, molding temperature, cooling speed, design and prepara-tion quality of crystallization apparatus, etc.

What worth to be discussed is the effect of change in melted alloy composition on above mentioned problems and eliminate �-Fe crystals. In fact the columnar crystals can be seen directly without using microscope which is mainly quadrilat-eral crystals of Nd2Fe14B, and there is Nd-rich phase among the columnar crystals. Crystals of casted ingot grow fast along with directions of (411) and (410), that can be certified by intensity of diffract peaks on (410) and (411) planes in X-ray diffraction pattern. There were a lot of equiaxial crystals’ texture in alloy ingot, copper was severely segregation, copper was poor within crystals, and dissocia-tive copper enriched among crystals, which is not good for improving magnetism of the alloy, because of unreasonable melting procedures in melting Sm(Co, Cu, Fe, M)7.4 (M=Zr, Mn).


By enlightened of “dynamic crossing” the aforementioned alloy ingot was put into solid solution treatment. The elements were fully “dynamic crossed” when solid solution treatment, that copper enriched a lot on crystal boundary and dis-solved dispersed copper and part of 1:5 phase of copper existed on wall of cell to act as pinning function on domain wall, that would act to rise coercivity of the alloy. The equipment to do that requires only a crystallizing apparatus with excel-lent heat conductive coefficient and a good coolant structure.

3. Application of the “dynamic crossing” in development of subject selection of the rear earth permanent magnetic alloy:

Search after a element to be add into nanocrystal compound exchange cou-pling permanent magnetic alloy to raise magnet performance of material.

Improve process to enhance property of the nanocrystal compound exchange coupling permanent magnetic alloy.

Increase alloy types under the principle that compound hard magnetic phase and soft phase within the range of nanocrystals.

The coercivity is the function of alloy composition and manufacturing proc-ess. Thus change alloy composition and improving manufacturing process can enhance the coercivity of SmCo5, Sm-Co 2:17 type, rear earth ferrous base and the nanocrystal compound exchange coupling permanent magnetic alloys, SmFeN, etc., and increase stability of the alloys.

Using allying method to raise integrative performance. For the last 50 years, driven by demand, China has gradually formed a

unique high temperature alloy system. However, the technology and performance is not stable, which need to be solved based on a more solid foundation. And one of the key to further development is to develop high temperature alloys used in civilian industry market (Shi, Zhong, 2010).

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Chapter 2 The First Generation Rare Earth Permanent-magnet Alloys

The purpose of this chapter is to discourse the relationship between phase trans-formation and the magnetic energy product of the 1st generation rare earth per-manent magnetic alloy at both room temperature, and from 25 to 950 . The in situ and dynamic observation and analysis were reported systematically and integrally about high temperature phase transformation of the SmCo5 filmy specimen under JEM-1000 HVEM (High Voltage Electronic Microscope). The observations revealed the driving force, resistance, change of free energy and free enthalpy, nucleation power, nucleation velocity, nonhomogeneous nuclea-tion, and new phase growth rule in phase transformation of SmCo5, and was rose to solid phase transforming theory. Research was conducted on the mecha-nism of the nonlinear variations in coercivity of SmCo5 permanent magnetic alloy from room temperature to 950 after annealing. The observations showed that the number of faults clearly increased at 750 on the Sm2Co17 phase which precipitated from the SmCo5 alloy. The measurements indicated that the SmCo5 shows the maximum amount of irreversible magnetic loss at 750 . In addition, some multi-faults area with very low anisotropy existed in the Sm2Co17 phase. These faults became anti-magnetization nucleation centers, which resulted in degradation of the coercivity of SmCo5 alloy to the lowest point, following annealing at 750 . Through the phase transformation process of SmCo5 permanent magnetic alloy, additional information and new theoretical opinions can be revealed. Further discussions were carried out on the prepara-tion principle and process, composition and process innovation, and crucial ma-terial preparation technique of the alloy and their influence on performance and microtexture of SmCo5 permanent magnetic alloy. The magnetism of SmCo5 permanent magnetic alloy was researched at 1.5K.


2.1 High temperature Phase Transition and Magnetic Proper-ties of SmCo5 Permanent-magnet Alloys

For magnetic scientists who are engaged in the investigation and development of magnetic materials, it is very interested in the relationship between the high tem-perature phase transition and coercivity for SmCo5 permanent-magnet alloys. Therefore, investigation of the project can accelerate the improvement of mag-netic properties and the development of coercivity and micromagnetic theories.

The coercivity of SmCo5 permanent-magnet alloys cannot be explained by the traditional coercivity theory. It is not consistent with the results predicted by the impurity or single-domain theories. The coercivity of SmCo5 alloys regularly ex-hibits an abrupt minimum at 750 as the alloys are annealed from 25 to 900 . In addition, the coercivity acquired resolved up to now (Pan, Jin, 1990).

A lot of studies about eutectic decomposition, precipitation phase, crystal de-feat, phase transition and microstructure have been made for SmCo5 alloy, since it was discovered in 1967 (Pan, Jin, 1990; Zhou, 1990; Xie, et al, 1980; Strnat, Ray, 1974; Wan, 1977; Fidler, Kronmüller, 1980; Li, Dai, et al, 1982). As the coerciv-ity of SmCo5 is very sensitive to the crystal defeat, precipitation phase and aging treatment, these problems have attracted much attention. Table 2.1 lists experi-

Table2.1 Opinions of researchers on decline of Hc of SmCo5 at 750 from 1972 to 1995

Year Author Main conclusion

1972 Den Broeder SmCo5 was eutectically decomposed into two phases: Sm2Co7 and Sm2Co17 (Den, Buschow, 1972)

1973 Rao After annealed at 750 , Sm2Co17 phase was separated from SmCo5, based on the observation of optical microscope and TEM. The coercivity decrease due to Sm2Co17 phase (Rao, et al, 1973)

1974 Smeggil Support above viewpoint (Smeggil, et al, 1974)

1974 Martin

The coercivity of SmCo5 greatly decreased, after annealed at 750 for a few minutes. This fact cannot be interpreted by the coercivity theory based on eutectic decomposition, since the eutectic decomposition cannot occur in so short during (Martin, et al, 1974)

1974 F.J.A Den Broe-der, et al

The coercivity of SmCo5 decreased at 750 because of the easy nucleation of reverse magnetization (Den Broeder F.J.A., et al, 1974)

1975 Pfeiffers There was no eutectic decomposition in sintered SmCo5, based on electron optics method (Pfeiffer, Metallkunde, 1975)

1976 Den Broeder

For heat treatment about 700 , Sm2O3 could be formed if oxygen at crystal boundary is enough. Not only this would weaken the domain wall pinning, but also promote the reversely magnetized nucleation, thus leading to a decrease in coercivity (Den, Zijstra, 1976)

1977 Kütterer

Based on calculation, the coercivity of SmCo5 decreased at 750 due to the appearance and growth of reversely magnetized nuclea-tion in each grain. The nucleation was large in size and inside the anisotropy was small (Kütterer, Hilzinger, Kronmiiller, 1977)


Continued Table 2.1

Year Author Main conclusion

1978 Kronmüller Support above viewpoint (Kronmüller, 1978)

1979 Nishio A decrease in coercivity at 750 occurred in short time, while the separation of Sm2Co17 phase was just observed after anneal for relatively long time (Nishio, et al, 1979)

1980 Fidler

The transition region between Sm2Co17 and parent phase SmCo5 was observed. The transition phase can have different composi-tions to Sm2Co17, and can form the reversely magnetized nucleus (Fidler, Kronmüller, 1980)

1982 Fidler Some of Sm2Co17 phase grew in anneal at 710

1983 Pan Shuming, Jin Hanmin

Based on in situ observation at high temperature using TEM with 1000V high-voltage and positron-annihilation technology, the decrease in coercivity of SmCo5 at 750 is due to factors as follows: the separated Sm2Co17 itself is not the center of reversely magnetized nucleation. Some defect region is the center due to the low magnetic anisotropy, and thus leading to the decrease in coer-civity (Pan, Jin, Tian, 1983)

1990 Pan Shuming, Jin Hanmin

In anneal from 25 to 1000 , the minimum coercivity and the maximum irreversible magnetic loss were found at 750 . The distribution of Sm3+ in the reference (Pan, Li, Ma, 1990)

1995 Zhou Shouzeng

As SmCo5 alternated the heat treatment between 750 and 950 , reversible change in coercivity was related to the formation and disappearance of inhomogeneous solid solution. The inhomogene-ous solid solution of Sm ions and Co or O ions must be formed due to narrow homogeneous regions, as SmCo5 was heated from 700 to 750 . The rich region of Sm or Co ions inside SmCo5 had low coercivity and formed the center of reversely magnet-ized nucleation, which leads to low coercivity. As heated at 950 , the inhomogeneous region disappeared and the coercivity regained (Zhou, et al, 1990)

mental results and main conclusions reported by some researcher during the re-cent thirty years.

Table 2.1 showed that some common conclusions had been reached. 1. SmCo5 is stable for annealing above 800 . It is sub-state below 800 ;

eutectic decomposition occurs in annealing, from SmCo5� Sm2Co7+Sm2Co17. 2. The different ratio between Sm and Co leads to different separated phases.

Sm2Co17 is precipitated from Co-rich SmCo5, but Sm2Co7 is precipitated from Sm-rich SmCo5 after eutectic decomposition.

3. The coercivity of SmCo5 is determined by the nucleation field of reversal magnetization.

4. The rotation theory that high magnetocrystalline anisotropy leads to high co-ercivity cannot explain the coercivity of SmCo5. However, large domain wall en-ergy (about AK ) and narrow wall width (about /A K ) produced by the large magnetocrystalline anisotropy energy derives the wall displacement theory of large coercivity (Li, Dai, et al, 1982), where A and K are exchange and magne-tocrystalline anisotropy energies, respectively.

5. The narrow and large anisotropy field is easily influenced by various defects.


6. As Sm is easily oxidized to Sm2O3, the coercivity is very sensitive to heating treatment technology.

7. The reversible magnetization process of SmCo5 is controlled by reversally magnetized nucleation. However, the reversible magnetization of the whole mag-net is avoided due to pinning at crystal boundary (Pan, Jin, 1990).

8. Some effects of defeats on the nucleation and pinning have been confirmed by TEM.

The new rule, original theory and new explanation obtained from studies of ba-sic theories and applications for the mechanism of SmCo5 permanent-magnet al-loy are introduced as follows:

1. The experiments, performed by Zhou, et al in 1983, showed that the coerciv-ity of SmCo5 significantly decreases after annealing at 700-750 for short time. The coercivity of the sintered SmCo5 can be regained after it is annealed at 730 for less than 250 min and followed by heating at 950 . However, after it is an-nealed at 730 for more than 400 min and followed by heating at 950 , only can the coercivity be partly restored. X-ray diffraction showed that the precipita-tion phase was not observed for SmCo5 alloys annealed at 730 for 50 h. As SmCo5 alloys alternated the heat treatment between 750 and 950 , the re-versible change of coercivity is related to the formation and disappearance of inhomogeneous solid solution. SmCo5 is the homogeneous solid solution alloy, as it is heated at 950 and followed by fast cooling. However, as it is heated at 700-750 , the Sm-rich or Co-rich region (or atomic cluster) has to be formed due to relatively narrow homogeneous regions. The anisotropy of the rich regions is rela-tively small, and thus forming nucleation centers of reversely magnetized do-mains, which leads to a decrease in coercivity. The inhomogeneous solid solution of Sm ions and Co or O ions must be formed due to narrow homogeneous regions, as SmCo5 was heated from 700 to 750 . The rich region of Sm or Co ions inside SmCo5 had low coercivity and formed the center of reversely magnetized nucleation, which leads to low coercivity. As heated at 950 , the inhomogeneous region disappeared and the coercivity recovered (Zhou, et al, 1990).

2. New viewpoints discovered by the author from experiments are: (1) The coercivity mechanism and segregated phase in annealing were investi-

gated for SmCo5 permanent-magnet alloys in 1982 using positron-annihilation, super-high voltage TEM and magnetic measurement. The results showed that the Sm2Co17 phase itself is not the center of reversibly magnetized nucleation. Some defects in the separated phase have low magnetic anisotropy and, therefore form the center of reversibly magnetized nucleation, which leads to low intrinsic coer-civity.

As SmCo5 is annealed from 25 to 900 , the maximum irreversible loss was found at 750 (Pan, Jin, 1990).

(2) In 1982, new high temperature phase transition, the separation of Sm2Co17


phase from SmCo5 and the growth of the phase, has been observed using super-high voltage TEM with 1,000 kV for SmCo5 permanent-magnet film with the thickness of 90 nm. At room temperature, no separated phase is observed. As temperature is raised to 300 , the fine precipitates are separated. At 350 , growth of the precipitates grows can be observed. At 400 , the precipitation and growth continue. After 420 for 30 min, the precipitates greatly grow. At 550 , the precipitates continue growing, but the growth rate is slowing and the stripes of Sm2Co7 appear on the original precipitates Sm2Co17. The precipitates grow slowly at 600 and are merged at 680 . The precipitates are confirmed to be Sm2Co17 by electronic diffraction analysis (Pan, Zhao, 1989).

(3) In 1982, the maximum irreversible loss occurs at 750 , as SmCo5 perma-nent-magnetic alloys are annealed from 25 to 900 . The mechanism was dis-cussed (Pan, Jin, 1990).

(4) In 1982, it was observed from TEM that separated Sm2Co17 is not perfect and are stripe and circle in shape. The separated phases at 650 are random in the direction. At 750 , the separated phases turn into particles (Pan, Zhao, 1989).

(5) In 1983, it was found that no peak of oxygen was observed in photo-electronic spectra for the sintered SmCo5 permanent-magnet alloys followed by annealing at 750 . Base on the results measured at room temperature, about the Sm and Co elements segregated at 900 , the concentration of oxygen does not observably change after annealing for various temperatures (Pan, 1992).

(6) In 1992 observation using 1000kV HVEM found that there was transit area (or called as interface of phases) between the matrix phase of SmCo5 and the pre-cipitated phase (Sm2Co17 or Sm2Co7) at 750 , such a transit area had low mag-netic anisotropy, and in defect areas in the Sm2Co17 phase the magnetic anisotropy was very low, which lead to degradation of the coercivity. When temperature was raised from 750 to 900 the observation discovered that the aforementioned transit area and defect areas in Sm2Co17 phase disappeared; it was found from photoelectron energy spectrum that the phenomenon of segregation in Sm and Co elements reduced gradually and at last disappeared, and the segregation area was testified as SmCo5 phase by X-ray diffraction analysis. That was the reason that the coercivity of SmCo5 alloy declined to the lowest after annealing at 750 , but it then turned up when temperature was raised to from 750 to 900 .

(7) In year 1994 to 2003 the dynamic observation on SmCo5 film sample by us-ing HVEM at increasing temperature found new phenomenon and new informa-tion about phase transformation at high temperature in experiment as follows:

Nucleation speed and coarsening of the new phase (such as Sm2Co17) pre-cipitated from the matrix phase change along with heating and holding times, solid state phase transformation and new phase coarsening process change with the time.

Part of Sm2Co17 was found to be resolved into the matrix phase again and


disappeared when heated to 600-750 in dynamic observation. The difference in structure between precipitated Sm2Co17 phase and the matrix phase leads to form new interface which increases system free energy. Thus, extremum of free energy

variation can be deduced as: �Gz = 3 32

az2 hz �, where az, hz are critical size of

small phase point of Sm2Co17 corresponding to different temperature; � is the in-terface energy between new phase and the matrix phase.

The growth rate of new phase of Sm2Co17 precipitated from the matrix phase of SmCo5 was also found to be different in the in situ and dynamic observation.

Sm2Co17 phase can precipitated from the position of Sm2Co17. That SmCo5 eutectoid dissolved out phases of Sm2Co17 and Sm2Co7 needs a

process and such a process is slowly in heat activation, the precipitated Sm2Co17 and Sm2Co7 phases exist synchronously, but correspondingly existent time of Sm2Co7 phase is shorter than that of Sm2Co17 phase. The minimum nucleation field for eutectoid decompounding of Sm2Co17 phase from SmCo5 phase is: Hn � (1/h +1/a)& /Ms.

The following orders were found by the in situ and dynamic observation: SmCo5 samples were different in coercivity and in variation extent of the coerciv-ity at temperature from 25 to 1000 , and generally the SmCo5 samples with higher coercivity the degradation of their coercivity would be smaller at 750 , and their irreversible magnetic loss also be smaller; and the height of coercivity will effect on nucleation rate as well.

Phase transformation is reversible, that is to say, Sm2Co17 and Sm2Co7 phases may dissolve into the matrix phase of SmCo5 again, once the sample is put on eletronic diffraction repeatedly in the in situ and dynamic observation, then its reversible property will become irreversible.

3. Experiment made by Hanmin Jin, et al. in 1983 indicated that SmCo5 alloy was annealing at low temperature (50-400 ) the coercivity and longevity of posi-tive electron had complicated change obviously along with last of time. When annealing at 710 its coercivity and nucleation field did degraded obviously, but the oblivion longevity of its positive electron was not changed basically. SmCo5 and rhombic Sm2Co17 have the central Sm3+ in conditions as more as 8 types with magnetisable axis along plane c or near plane c among 11 most possibly defect near neighbor Sm3+ distributions. Therefore, there are low magnetic anisotropy areas in the defect areas, such areas may become reverse magnetization nuclea-tion centers.

4. Experiment carried out by Hongzu Xie, et al. in 1978 showed: when heated at 750 the electric resistant rate of SmCo5 alloy increased with last of time.

5. Experiment conducted by Wanji Zhao, Hongru Zhai, et al. in 1983 showed: there was no magnetic split when adding tin alone into SmCo5; but adding iron into SmCo5 could result in magnetic split, that addition of iron was the root caus-


ing tinny state magnetic split. Using the Mössbauer Effect is a near practice and good method to observe and study the function of adding iron and tin into SmCo5.

6. Experiment by Houding Song in 1978 indicated: addition of iron and tin into SmCo5 can make for a stable manufacturing process, the SmCo5 with addition of iron and tin has a high intrinsic coercivity as the same as that before the addition.

7. Experiment by Xianchun Yang, Jinhua Wu, et al. in 1982 indicated: coerciv-ity of (Sm, Gd, Er)Co4.5 made by adding gadolinium and erbium was related to precipitation of 2:7 phase, and am appropriate heat treatment condition can make the magnet sample with a proper amount of 2:7 phase, symmetrical and integrated granule of crystals, magnetisable, complete domain orientation, hard to be de-magnetized and a high coercivity.

8. Experiment by Mingjie Cheng, Hongzu Xie, et al. in 1982 showed: there were only two phase existed in edge precipitation process, besides SmCo5 phase

the 2nd phase for Sm-rich sample was Sm2Co7 phase only, and the 2nd phase for Co-rich sample was Sm2Co17 phase only.

9. Experiment by D. L. Martin, et al. in 1974-1975 indicated: SmCo5 phase could not be formed at temperature below 800 and CeCo5 phase could not be formed either. RECo5 is instable and it will be eutectoid decompounded at certain temperature.

10. Experiment and calculation by J. Fidler, et al. in 1980 indicated: the place of dislocation had big stress field and strain energy (Fidler, Kronmüller, 1980). Nucleation in precipitated phase caused change of system free energy which is expressed by �G, thus

e ep p p m p p p p m1/ ( )d d 1/ 2 dD D

D y yG V G G G G V w n V' (� � � � � � �� (2.1)

where Gm, Gp are the chemical energy of a single cell of the matrix phase and precipitated phase, respectively; Vp is the volume of a single cell;

epG is elastic

aberrance energy;

epDG is interactional energy between dislocation and precipi-

tated phase; wp is the surface energy density of precipitated phase; n is unit vector of surface normal of the precipitated phase; D

y' is the stress at place of dislocation; Dy( is the strain of dislocation.

Symbol minus (�) in the formula is the driving force for precipitation process, plus (+) is resisting force for the precipitation process.

Assuming that the elastic property of the precipitated phase is the same as that of base body, and they are anisotropy. When the precipitated phase is a flat ellip-soid, thus the elastic aberrance energy can be obtained as:

e:2 7G = (0.043 + 0.174 Y/Z)

e:217G = (�0.908 + 0.1736 Y/Z) (2.2)

where, Z and Y are the short radius and long radius of the ellipsoid precipitated phase, respectively.


Precipitation energy of 2:17 phase is favorable by looking at elastic aberrance energy alone. This is one of the reasons that the 2:17 phase precipitates at first in SmCo5 alloy, which had been testified be many experiments.

Interaction energy of precipitation and dislocation

epDG is expressed by formula

as follows:

ep p

D DD y yG V' (� (2.3)

where Dy' is the stress of dislocation; D

y( is the elastic strain of precipitation; pV is the volume of the precipitated phase.

According to Eq. 2.3 the interaction energy between an atom of the precipitated phase and the dislocation stress field can be estimated. Then a conclusion was obtained from above calculation: multiples dislocation is like to be nucleation center of Sm2Co17 phase (only a few screw dislocation will be favored to be nu-cleation center of Sm2Co7 phase), thus precipitation of phase Sm2Co17 is easier than that of Sm2Co7 phase. Though rhombic columnar dislocation has a strong effect on domain wall, its contribution to coercivity is not very big. Fidler con-cluded on study in theory and practice that sheet precipitated phase in some dislo-cation has strong pinning effect on domain wall; but even precipitated phase al-most has no effect on domain wall, unevenly distributed precipitated phase on interspersed dislocation has strong pinning effect on the domain wall.

11. Kronmüller derived micro-magnetic theory to calculate in 1976 the reverse magnetization nucleation field HA being formed by the sheet precipitated phase in SmCo5, the coercivity formula was obtained as follows:

mHc = 2kp/Mp – 2�Mp + 2Ap�2/ (MpD2) (2.4)

where subscript p represents physical amount of the precipitated phase, by substi-tuting parameters of the precipitated phase in to Eq. 2.4 the nucleation power can be estimated as 7,960 kA/m (100 kOe). This value is far higher than the coercivity obtained in practice. The problem was that the calculation neglected interface (transit area) between the matrix and precipitated phase, and the structure and magnetism of the transit area might act important effect on reverse magnetization domain nucleation. When thickness D of the precipitated phase is very small, Z0 = 0, zr0�D (r0 is the transit area of phase interface, and Z and Z0 refer to the figure), reverse magnetization domain may nucleate on uneven phase interface. In this condition the nucleation field can be calculated as per continuity of magnetic the-ory as below.

1/ 2n n p 02( � ) /( )H H A K M r)� �

Hereby introduce thickness of the domain wall �B, assuming �K�K1, and mak-ing use of the following formula:

�n = � p/A K � � /A K�


Thus the formula above can be written as:

n n p B p 02 /(� )H H K M r�)� � When thickness D of the precipitated phase is very small thus D�2r0, and when

phase interface transit area r0 is also small (r0��0), the formula of nucleation field can be expressed as:

maxn n 1 B p 02 /(� )H H K M r�)� �

Actually, nucleus is very small, thus the exchange integral constant of the pre-cipitated phase and magnetization intensity can not be take as constants any more because D�2r0 and r0��B. In the condition of a very thin precipitated phase the anisotropy constant of the precipitated phase Kp will be lowered considerably; if K/K1 is lowered to 10�3~10�2, the Hn

� would become meaningless, thus Hn = (2K1�B)/(�Mpr0). Assuming �n = 2.6 nm, r0=26 nm, K1 = 15×107 erg/cm3, Mp = 955 emu, and Hn can be calculated as Hn = 796×103 kA/m, that is coincident with the experiments and can explain the reason that the coercivity of SmCo5 is lower than its theoretical value is caused by phase interface of the precipitated phase. The coercivity of SmCo5 mHc is determined by nucleation field of reverse domain (Kütterer, et al, 1977).

Nucleation of reverse magnetization domain is a rotational process of magnetic torque, and that the expanding process of reverse magnetization domain is a proc-ess of displacement of domain wall; the growth process of reverse magnetization domain nucleus needs to overcome energy of demagnetization field and resistance by increased domain wall energy because of increased surface area of magnetiza-tion domain. In this instance coercivity of alloy is mainly determined by critical field of reverse magnetization domain expansion. Kronmüller educed after ana-lyzing and calculation that the critical field of reverse magnetization domain ex-pansion of SmCo5 was far smaller than nucleation field.

2.2 The in Situ and Dynamic Observation on High Temper-ature Phase Transformation of SmCo5 Permanent Magnetic Alloy at 25-750 *

SmCo5 permanent magnetic alloy is provided with very high saturated magnetic intensity, Curie temperature and magnetic crystal anisotropy. It has been devel-oped speedily and applied widely since it was born. Some problems have ap-peared along with its application and thereby brought forward a series of new subjects which are worthwhile to study in-depth.

Cooperators of this study include: Fengzuo Tian, Chengzhou Yu, Shikuan Ren, General Re-

search Institute for Nonferrous Metals; Hanmin Jin, Jilin University; Guokun Li, Yongshu He, Research Institute of High-Energy Physics, the Chinese Academy of Science.


2.2.1 Magnetic measurement

Early in 1970 Wenstendorp discovered in experiment that the coercivity of SmCo5 permanent magnetic alloy varied nonlinearly along with rising of annealing tem-perature, as shown in Fig. 2.1.

Fig.2.1 The relation curve of coercivity of SmCo5 permanent magnet alloy at room

temperature vs annealing temperature 1 Coercivity vs simtered temperature; 2 Coercivity vs annealing temperature

It is observed from Fig. 2.1 that the coercivity of SmCo5 was degraded to mini-

mum after annealing at 750 . Afterwards the coercivity was raised along with rising of annealing temperature over 750 . This regularity that the coercivity of SmCo5 varies with change of annealing temperature was verified by experiments. Author, et al. worked out experimental curve in 1981, as shown in Fig. 2.2. It can be seen that the regularity for both intrinsic coercivity and technique coercivity was consistent.

It can be seen from Fig. 2.2 that the intrinsic coercivity appeared a small peak value by annealing at 600 or below, but generally varied a little. After exceeding 650 the coercivity degraded and reached minimum at 750 . Annealing at 750–850 the coercivity began to rise until 850 at which the coercivity reached the maximum, and degraded after 850 . This phenomenon attracted great interesting of people who engaged with study and application of SmCo5. Thus a great lot of experimental studies were carried out, and as the significance the study results en-riched and developed the coercivity theories, and so that to provide experimental basis in practical application and to promote development of magnetic theory.

Since 1970 up to now many scholars presented their thesis related to annealing


Fig. 2.2 The relation curves of coercivity of SmCo5 alloy vs annealing temperature

(The above one of two curves shows the relation of the intrinsic coercivity vs temperature from 0 to 1000 . The one below shows the relation of the technique coercivity of SmCo5 samples vs tempera-

ture from 0 to 1000 . Both curves indicate that the coercivity goes down at 750 and goes up at 800 )

effect on coercivity of SmCo5. Those can be concluded to 5 sorts in accordant with their opinions.

1. The annealing effect on coercivity of SmCo5 is caused by eutectoid decomposition at 750 that the SmCo5 decomposed in to two phases of Sm2Co7 and Sm2Co17. Early in 1974, K. H. J. Buschow presented his thesis in J. less-Common Met. in which he believed that eutectoid decomposition of SmCo5 would occur at below 800 ,i.e., SmCo5 � Sm2Co7 + Sm2Co17. And in the same year F. A. Den Brocedan, et al also supported the opinion mentioned above. They consid-ered that the Sm2Co7 and Sm2Co17 phases eutectoid decomposed from SmCo5 had the magnetic crystal anisotropy was much lower than that of SmCo5. These two phases may become nucleation center on antimagnetic domain and so that re-sulted in degradation of coercivity. In 1975 I. Pfeiffer, et al. and D. L. Uartin et al. considered that the SmCo5 alloy did not form SmCo5 phase at 800 or below.

2. The annealing effect on coercivity of SmCo5 was due to precipitation of Sm2Co17 phase in which the magnetic crystal anisotropy was smaller in a quantita-tive level than that of SmCo5. In 1978 H. Kronmüller believed that the precipi-tated Sm2Co17 phase would become antimagnetic nucleation centers to degrade the coercivity.

3. The annealing effect on coercivity of SmCo5 was due to spinodal decomposi-tion reaction. Y. C. Chuang, et al suggested in 1982 that RE2Co17, RE5Co19 and RE2Co7 derived from structure of SmCo5 phase: 3RECo5 �*RE + 2Co � RE2Co17, or 4 RECo5 �*Co + RE � RE5Co19. That indicated that decomposition of SmCo5 did not need fluctuation of components and formation of new phase which may lead to decline of system free energy.

4. The annealing effect on coercivity of SmCo5 was due to more oxygen being


dissolvd in SmCo5 at 750 . F. J. Den Broeder, et al. presented above opinion in 1976. Many people believed that when SmCo5 was annealing at 700 – 800 the precipitated Sm2Co17 phase was formed due to absorption of oxygen, which re-sulted in formation of Sm2O3 by oxidation of samarium and appearance of cobalt enriched area, then so that to promote precipitation of Sm2Co17 phase.

5. The annealing effect on coercivity of SmCo5 was caused by anisotropy from heat expansion of SmCo5. K. Kumax, et al. suggested in 1978 that heat expansion among crystal granules naturally caused different heat stress and the difference in cooling speed also caused different in heat stress.

Study on the subject of annealing effect of SmCo5 and make it clear are of sig-nificance both in practical application and theory. For this purpose Shuming Pan, Hanmin Jin, et al. studied this subject from 1980 to 1983 by using photoelectron spectrum, antielectron oblivion technology, X-ray diffraction and 1000 kV HVEM dynamic observation technique, and combining magnetic measurement and com-puting magnetic anisotropy of Sm3+ in imperfect structure of Sm2Co17 precipitate phase. As the result they presented a new viewpoint different with above men-tioned 5 opinions and formed new theoretic opinion.

2.2.2 Sample preparation and experiment method

2.2.2.1 Preparation process

The SmCo5 sample for experiment was confected by using alloy melting method, i.e., melted in arc furnace under argon protection atmosphere. The ingot was crushed to about 5 m by vibrant boll grinder (protected from oxidation in organic medium) and orientation formed under 1.5 T magnetic fields. The formed sample was sintered at 1,120 for 1 hour and cooled to 900 in a speed of 0.8 /min and held. Afterwards the sample was quenched to room temperature. The specimen after heat demagnetization was machined to required specification for observation.

2.2.2.2 Method for antielectron oblivion experiment

The antielectron oblivion longevity () and Doppler widened linear parameter (S) of oblivion radicalization were measured for above mentioned specimen at room temperature.

The longevity spectrum instrument is a quick and slow coincident system, and its time resolving power (FWHM) for 60Co radiate source is 260 ps. Each speci-men was measured 2-4 times repeatedly. Total countered number for each meas-urement was about 8×103. Measured data was treated by index fitting method and analyzed 1 and 2 two longevity components. Among them the long longevity group 2 has a comparative intensity was very small, that was considered to be caused by source effect. Thus the short longevity group (1) was taken as the aver-age longevity of antielectron obliterated in integral crystal lattice and defect of


specimen. When new vacancy defects occurred, the longevity of antielectron would be increased because density of electron decreased in defect.

The Ge( Li ) spectrum instrument for measurement and S parameter has a resolving power (FWHM) of 1.45 keV. Each specimen was measured 3 – 5 times repeatedly and total countered number for each measurement was about 7×105. S parameters were computed directly by online computer. In this experiment the S parameter was defined as:

S = A / (B + C) where A is counter number of 24 paths in center of the energy spectrum of oblit-erated radicalization; B and C are counter number taken from 25 paths in two wings of the spectrum. As shown in Fig. 2.3 that energy corresponding to each path was 75eV. The more defects the specimen has the smaller the average elec-tron momentum will be in the defect location. As the result the counter path in center will be increased and the counter path in wings be decreased so that the S parameter will be big.

In the measurement the non-carrier 22Nd clamped between polyester films were taken as antielectron source, and the source was lamped between two same specimens and its strength was 5 C.

Fig. 2.3 The diagram of the definition of S-factor of Doppler broadening

2.2.3 Influence of annealing treated specimen on coercivity

All specimens were SmCo5 magnet with samarium enriched equivalent amount of component quenched from 900 besides those specially explained. Magnetism measurement was carried out at room temperature. The magnet was cut to 15.5×7.5×0.95 lamellas with its easy magnetisable axis paralleling with 7.5 mm rhombohedra. After being polished using metallographic paper each of two specimens were annealed at 710 for 4 min, 8 min, ,60 min and 80 min at pro-tection of argon atmosphere and quenched, and then measured the demagnetized curve. These loop-lines were makeup by superposition on a big loop-line of high coercivity with small loop-lines of low coercivity. The latter disappeared after surface electrolytic polishing by dozens microns. It is obviously that those small loop-lines come from surface cobalt enriched layer formed by preference oxida-tion of samarium (Pan, et al, 1983; Fidler, et al, 1980).


Fig. 2.4 indicates intrinsic coercivity after electrolytic polishing and the relation between the comparative volume �Br / (Br+�Br) of cobalt enriched surface layer phase and aging time, where �Br and Br are remanence of small loop-line and big loop-line before polishing. The intrinsic coercivity is not very sensitive to exis-tence of cobalt enriched layer, which explains that the domain wall is pinned strongly by crystal boundary. The degradation of coercivity is caused by decrease in nucleation field of antimagnetic nucleation in crystals.

Fig. 2.4 Annealing time dependence of iHc, S and �Br/(Br+�Br), in which annealing temperature is 750

The antielectron oblivion longevity () and Doppler line widening parameter S

were measured for this specimen group (Fig. 2.4), among them value of the latter indicated widening extent of curve. That the and S did not varied with passing of time is not consistent with the result measured at room tempera-ture. This phenomenon indicated that there was no vacant defect appeared in process of aging at 710 .

Experiments for influence by ag-ing on coercivity were made for mass specimens at different tem-perature within 25-1000 and ag-ing 1 hour. That obtained the similar results. The intrinsic coercivity dras-tically degraded by annealing at 750 for short time and aroused at 900 , seeing Fig. 2.5.

Fig. 2.5 The relation of the intrinsic coerciv-

ity iHc vs annealing temperature (The composition of the specimen is 37.66%Sm,

62.34%Co. It is showed that the intrinsic coercivity has the lowest value

at 750 )


2.2.4 The in situ and dynamic observation by 1000 kV HVEM under heating condition

The specimen with constituents of Sm 37.66% (wt.), O2 1.2% (wt.) was sliced into lamellae in direction perpendicular to magnetisable axis and then was ionic thinned to film for observation by using electronic microscope. The used microscope was JEM-1000 type with 1000 kV, the vacuum was 1.6×10�5Pa (1.2×10�7 Torr) and after addition of liquid nitrogen the vacuum was 9.3 × 10�6Pa (0.7×10�7 Torr).

2.2.4.1 Observation by JEM-1000 HVEM at room temperature

The observation result and analysis at room temperature: the Fig. 2.6(a) shown electronic micrograph of SmCo5 alloy at room temperature that most of area in the electron micrograph did not have precipitate but once in a while may saw a few defect such as layer dislocation and in the same area also observed precipi-tated phase. Fig. 2.6(b) and Fig. 2.6(c) were electronic diffraction pattern and in-dexes of the matrix which was determined as SmCo5.

Fig. 2.6 The microstructure of SmCo5 permanent magnet alloy at room temperature (a) The picture of the microstructure under electron microscope of SmCo5 permanent magnet alloy at

25 ; (b) Indexing of electronic diffraction pattern of SmCo5 permanent magnet alloy at selected area under 25 ; (c) The electronic diffraction pattern of SmCo5 permanent

magnet alloy at selected area under 25


2.2.4.2 The in situ and dynamic observation on phase transformation at high temperature by JEM-1000 HVEM

This section revealed the whole process of phase transformation of SmCo5 alloy from room temperature to 950 , including driving force, transformation resis-tance, homogeneous and heterogeneous nucleation, growth process, etc.

A. Observation on experiment result

The specimen of SmCo5 was heated in the side inserting type heating dais. The dynamic observation found that there is no precipitate in the selected area through tilting the rotation dais in many times. The SmCo5 specimen did not appeared any precipitates under heat activation before 350 . After raised to 420 rapidly there was highly dispersed small precipitate points which was not clear even after being magnified to 20000 times. When the temperature reached 420 and held for 2 min the highly dispersed precipitates could be observed clearly and fulfilled all of the vision in the place originally had not any precipitate. And being held for 40 min continually the precipitate phase would grow up to dozens nanometers, which was Sm2Co17 by electronic diffraction analysis (Fig. 2.7(a), (b), (c)) . The

Fig. 2.7 The microstructure of SmCo5 specimen under JEM-1000

ultra-high voltage electron microscope (Dynamic in situ observation of the high temperature phase transition, e.g. precipitation of

Sm2Co17-phase, upgrowth, coarsening, merging while heating temperature reaches 750 from 25 . Multi-defect area in Sm2Co17-phase appears as (a) to (i))


Fig. 2.8 indicated the holding time and curve of heating temperature along with change of time. The variation status of precipitate phase point while the specimen was held for 8 min as follows: some Sm2Co17 precipitate grown was not fast, but the Sm2Co17 precipitate phase point identified as spot Q grown 1.5 times, and the R precipitate spot grown twice.

Fig. 2.8 The heating-up speed curve of the 1000kV HVEM observation specimen

B. Analysis on nucleation speed

SmCo5 has three SmCo5 cells and when in one of the cell the rare earth atoms (RE) was replaced by a Co atoms the RE2Co17 will be form:

3 RECo5 �RE + 2Co +,RE2Co17

3 SmCo5 � Sm +2Co +, Sm2Co17

The Sm2Co17 precipitated from SmCo5 matrix phase nucleated at first depend-ing on heat activation process because of the growth of activated atoms in the matrix phase. The above mentioned Q and R points growing from 0.5 nm to sev-eral nanometers may exceed critical nucleus under a certain degree of supper-cooling and becomes nucleus of new phase. The number of nucleus of new phase increased along with time prolonging and rising of temperature from about 350 until highly dispersed precipitates fulfilled the all vision. It can be seen that the nucleation speed was very fast but was variable. The nucleation speed depends on atoms number jumped to critical nucleus by heat activation and nucleus number of critical size in unit time (Wang, 1982). That may be formulated as:

v Cexp [��G / (KT)] (dn/dt)

where v is nucleation speed; �G is the power for critical nucleation; T is tempera-ture; dn/d t is the atoms number jumping to crystal nucleus by heat activation in unit time; and C is a constant coefficient.

It can be seen from the formula that the nucleation velocity is higher at a higher temperature; the more the atoms number jumped to crystal nucleus by heat activa-


tion in unit time, the higher the nucleation velocity is.

C. Variation of free energy in phase transformation and power of nucleation

The new phase of Sm2Co17 precipitated from SmCo5 matrix phase grown 1-2 times in 8 min. What is the decisive factor for growth of the new phase? Transfer process of phase boundary is the growth process of new phase. Thus the transfer speed of phase boundary influenced the growth speed of new phase. On the boundary there is difference in concentration between the matrix phase and the new phase, or says, there is an equilibrium concentration. Difference in concen-tration existed in the matrix phase where the diffusion is caused by the concentra-tion difference. The growth of new phase is aim to restore equilibrium concentra-tion in the boundary, and that the atoms diffusion speed determines growth speed of the new phase.

It was observed dynamically that precipitation of Sm2Co17 phase from SmCo5

matrix phase was comparatively slow but after holding at 420 the speed of phase transformation became the highest. This was because that when the new phase of Sm2Co17 precipitated at beginning the nucleus was very small, volume increment of phase transformation was not big by linear growth, and the grown up of the nucleus caused volume increase in three-dimension. The dynamic variation process of (f ), (g), (h) and (i) in Fig. 2.7 is the new phase gradually grown up to the terminal stage. That mutual attraction, mutual contact, aggrega-tion and phase transformation existing among new phase points again make the speed lowered gradually.

Nucleation speed of new phase varies along with time and interface of phase varies along with time in nucleation process of solid phase transformation thus nucleation speed also varies with time, seeing Table 2.2. It can be seen from this table that nucleation speed and embryo growth and accretion in all varies along with passing through of time. And growth speed (accretion speed) slowed gradu-ally along with grown up of the new phase of Sm2Co17. This reason is that solute atoms needed for continually growth of new phase of Sm2Co17 (or called contin-ual growing coarse) need to be diffused from farther area (long distance diffusion) so that will need longer diffusion time (Feng, Shi, Liu, 2002). It was observed also that the SmCo5 alloy formed embryo of new phase of Sm2Co17 as the begin-ning because of heat activated concentration undulating and fluctuate, then the

Table 2.2 Nucleation rate of new phase changes with time

No. of Sm2Co17 new phase points (precipi-tated from SmCo5 ma-trix phase) (Fig. 2.7)

Volume increase from

(d) to (e)

From (e) to (g) crystal coarsen-ing and aggre-

gating

From (g) to (h)550�600

From (h) to (i) volume increase inhomogeneous area appeared in

new phase Nucleation velocity:

Used time t 3'36" 54" 3'36" 5'24"


component atoms coming from SmCo5 matrix phase continual diffused to sur-face of the new phase of Sm2Co17 to make small precipitate points (embryo) grow up. Not all of small phase points (or embryos) can grow up (or coarse) independently to massive phase point of Sm2Co17 and survived because the SmCo5 alloy is metastable at below 850 . And that part of small phase points (or embryos) was too small to be re-dissolved into SmCo5 matrix phase and disappeared because the smaller phase points have a bigger solubility so that they can not grown to be nucleus of new phase. In above mentioned process that the SmCo5 matrix phase in metastable status transferred consequentially to a stable new phase of Sm2Co17 and as the result to lowered the unit free energy of system, and the new boundary is formed because the difference in structures between newly formed new phase of Sm2Co17 (in above mentioned process grown up nucleus of the Sm2Co17 phase) and the matrix phase, that caused in-crease in free energy. The above mentioned process is formulated as:

�G = 3 32

a 2 hn�H (�T/T 0 )3 3 a 2 +6ahn� (2 .5)

where �G is variation of free energy; n is number of new phase point (embryo); �H is the value for phase transformation of enthalpy; �T is degree of supper-cooling; T0 is equilibrium point of phase transformation; � is energy on the boundary face between new phase and the matrix phase; a is height of the hexan-gular prism.

It can be seen from Eq. 2.5 that increase in a, h, n and � will increase the free energy for phase transformation also increase, and there is a functional relation between them; the easiness extend of phase transformation will determine the power required for nucleation. The power required for nucleation can be induced as follows: derives at first the extreme value in variation of the free energy using Eq. 2.5, the extreme value in variation of the free energy is:

�G z = 3 32

a2z hz � (2.6)

where az and hz is the critical size of small phase points of Sm2Co17 at different temperature.

Power for nucleation is the energy vallation necessary to get across for the critical nucleus (Feng, Shi, Liu, 2002), and is the critical value for variation of the free energy (�Gz). The higher the power of nucleation is the more difficult the phase transformation will be, i.e., the less the small phase points of Sm2Co17 precipitated from SmCo5 matrix phase will grow up. Only the small phase points of Sm2Co17 is bigger than the critical size of phase point those small phase points may grow up. In contrary, the smaller phase points will be automatically melted into the other phase of SmCo5.


D. Analysis on inhomogeneous nucleation

We have changed the view to find the area with many dislocations and layer dis-locations in the matrix phase through observation by transmission microscope. These small areas with many dislocations and layer dislocations have comparatively more impurities so that the phenomenon of new phase points of Sm2Co17 precipi-tated from SmCo5 matrix phase occurs in advance and more phase points (embryo) are produced. Those indicate that the nucleation will be easily. This is due to that energy at interface between different phases of new phase points (embryos) of Sm2Co17 and impurities is low (compared with non impurity phase) and the nuclea-tion is inhomogeneous and needs a smaller driving force for phase transformation. This explained that in dynamic observation under heating environment the small phase point of teens A size moves quickly, and the phenomenon of two new phase point of Sm2Co17 gathered together along with prolonging of time was observed, seeing Fig. 2.7 (e), (f) and (g) for two new phase spots of R and Q of Sm2Co17.

E. “Dynamic cross” and eutectoid decomposition

The precipitation of new phase of Sm2Co17 from the matrix phase of SmCo5 is long distance diffusion because Sm2Co17 and SmCo5 are constituted of different chemical compositions. Surely, condition for the diffusion is the step that atoms of matrix phase traverses the phase boundary under heat activation; the matrix phase has 3 crystal cells of SmCo5, and the samarium atoms in one of cells wan instituted by a transit cobalt atoms to form the new phase of Sm2Co17. This type of short distance diffusion that atoms transit the boundary from the matrix phase to new phase of Sm2Co17 and long distance diffusion mentioned above processed continually; and for growth rate of new phase the long distance diffusion acts the decisive function. The above process is the dynamic cross diffusion of Sm and Co under heat activation that can be formulated as 3SmCo5 – Sm + 2Co�Sm2Co17, in detail the diffusion formula is 3SmCo5 – Co + Sm � 2Sm2Co7, 3RECo5 – RE + 2 Co � RE2Co17, and 3SmCo5 – Sm + 2Co�Sm2Co17. The metastable status at below 805 is actually the temperature from 650 to 750 , at which the ma-trix phase processes the following eutectoid decomposition, i.e., SmCo5 � Sm2Co7 + Sm2Co17.

F. Observation and analysis on defects in Sm2Co17 phase precipitated from matrix phase of SmCo5

In dynamic observation on SmCo5 in heating condition it was seen clearly that defects existed in the new phase of Sm2Co17 precipitated from matrix phase of SmCo5, seeing Fig. 2.7. The new phase precipitated from (g) to (i) had some de-fects which possess very low magnetic anisotropy. And amount of the defects in the phase of Sm2Co17 reached maximum at 750 . We believe that this is the rea-son why the coercivity degrades by annealing at 750 .


G. Observation and analysis on phase of Sm2Co7 precipitated from the ma-trix phase of SmCo5

In the phase of Sm2Co17 fulfilled the vision there were big piece of hexagon, small piece of round D and Sm2Co7 phase of striation E, seeing Fig. 2.9 (a), (b). And it was observed specially that the phase of Sm2Co7 located among big pieces of hexagon spots of Sm2Co17 transformed to be Sm2Co17 phase (from the Sm2Co7


Fig. 2.9 Electron micrographs and electron diffraction patterns of the precipitation process (a-c) The precipitation of Sm2Co17; (a,b) The SmCo5 matrix decomposes eutectoidally into Sm2Co17

and Sm2Co17 striations; (c) The Sm2Co17 striations change to Sm2Co17 particles

(For the annealing conditions, see Fig. 2.7)

phase of striation transformed to be Sm2Co7 phase of small round spots), seeing Fig. 2.9 (c). The mutual effected energy between atoms and dislocation force field existed in desolvation precipitation of Sm2Co7 and Sm2Co17 from SmCo5. Is what kind of dislocation prone to form nucleation center for Sm2Co7 phase or Sm2Co17 phase? The Sm2Co7 phase is a hexagonal structure at high temperature and a rhombic structure at low temperature. No diffusion phase transformation from hexagonal structure to rhombic structure occurs in annealing. Generally, some screw dislocation avails to formation of nucleation center of Sm2Co7 phase, but many type of dislocations all avail to formation of nucleation center of Sm2Co17 phase (Zhou, et al, 1983).

SmCo5 phase and Sm2Co17 phase are coherent. Thus aberrance of crystal lattice and strain energy occurred in phase transformation at high temperature. For eutectoid decomposition of phases of Sm2Co7 and Sm2Co17 from SmCo5 matrix phase there is the formula as below:

3RECo5 – Co + RE +, 2 RE2Co7 3SmCo5 – Co + Sm +, 2 Sm2Co7 3RECo5 – RE + 2Co +,RE2Co17

3SmCo5 – Sm + 2Co +, Sm2Co17 It was observed when temperature was raised to 650 that the new phase of

Sm2Co17 appeared as integral hexagon and Sm2Co7 was in striation. The new phase appeared clearer but the striation microstructure appeared to be better grown. When temperature was raised to 750 the serried spots of Sm2Co17 new


phase transfer from area of Sm2Co7 phase, seeing Fig. 2.9 (c). Distribution of the precipitated spots of Sm2Co17 phase was littery, but between original precipitated big spots of Sm2Co17 phase there was a large serried row newly precipitated spots of Sm2Co7 phase.

H. Driving force and resistance for phase transformation of SmCo5

SmCo5 belongs alloy phase transformation of solid state, is a phase transforma-tion of nucleation – growth of nucleus. Its phase transformation driving force is the difference of free enthalpy between new and old phases (Liu, Ren, Song, 2003): �G driving force for phase transformation = G2 –G1 < 0 (2.7)

where G1 represents old phase, i.e., free enthalpy of SmCo5; G2 represents new phase, i.e., free enthalpy of Sm2Co17; �G represents driving force for phase trans-formation.

�G driving force for phase transformation = G free enthalpy of SmCo5 –G free enthalpy of Sm2Co17 < 0

Necessary proceeding precondition for phase transformation is that the free en-thalpy of Sm2Co17 new phase must be larger than that of the SmCo5 old phase. Phase transformation occurs spontaneously. Supposing n is atoms number, thus that each atom transforms from SmCo5 matrix phase to Sm2Co17 new phase will cause variation in free enthalpy as formulated below: �G = (G2 –G1) / n (2.8)

According to Eq. 2.5 and Eq. 2.6 as long as the inter-phase energy being worked out the �G can be calculated, inter-phase energy between SmCo5 matrix phase and Sm2Co17 new phase consist of structural inter-phase energy and chemi-cal inter-phase energy. The structural inter-phase energy is formed due to cutting off or weakening of atoms bond linkage which caused heightening of potential energy (Liu, Ren, Song, 2003). SmCo5 and Sm2Co17 are coherent. When Sm2Co7 and Sm2Co17 precipitated in SmCo5 crystals the atoms on the inter-phase will has a certain extent of mismatch because of their different lattice constants. When precipitate phase of Sm2Co7 exists in SmCo5 the degree of mismatch will be 0.8% in direction of a axis of lattice, and 2.3% in direction of c axis of the lattice; when precipitate phase of Sm2Co17 exists in SmCo5 the degree of mismatch will be �3.2% in direction of a axis of lattice and 2.7% � 7.4% in direction of c axis of the lattice. While the degree of mismatch increased to a certain extent the prop-erty of the inter-phase will be changed and the fully coherent inter-phase will be broken that dislocation appears in the inter-phase to form partly coherent. Its nu-cleation power is smaller than that for non-coherence and larger for fully coher-ence. The total energy of interphase is less than 0.5 J/m2. The geometrical item is lowered from high to very low and chemical item enhanced from low to compara-


tive high (Fidler, et al, 1980; Liu, Ren, Song, 2003; Zhou, et al, 1995). Configura-tion dislocation on interphase may reduce energy of interphase; inhomogeneous nucleation center for Sm2Co17 is dislocation; and inhomogeneous nucleation cen-ter for Sm2Co7 is crystal boundary. The structure of crystal lattice of SmCo5 is more complicated than that of normal 3d metals and some complicated disloca-tions, such as rhombic mast dislocation, dipole rhombic mast type dislocation, etc., may appeared in SmCo5 permanent magnetic alloy. The thickness of domain wall of SmCo5 is around 2.6 nm and the thickness of domain wall for cobalt is 14.7 nm, but the average action force of dislocation on domain wall in SmCo5 is large than that in cobalt. All of SmCo5, Sm2Co7 and Sm2Co17 belong to rare earth intermetallic compound, that the surface energy will be enhanced by chemical factors (chemical interphase energy is plus) in phase transformation or results in decrease in surface energy (chemical surface energy is minus). The difference between atoms bond linkage of interphase and interior atoms linkage of two phases causes enhancement in interphase energy. The bigger the concentration difference between SmCo5 phase and Sm2Co17 phase is the bigger the rising of interphase energy is caused to be. The nucleation of Sm2Co17 new phase may cause elastic aberrance and form stress field and elastic aberrance energy in lim-ited range around the nucleus. The desolution precipitation of SmCo5 matrix phase is more avail Sm2Co17 phase referring to elastic aberrance energy. The Sm2Co7 phase of striation type, which was appeared at 650-750 and observed by using 1000 kV HVEM, is actually the sheet phase precipitated along (0001) plane of matrix phase and perpendicular to c axis (Fidler, et al, 1980).

In conclusion, the above in situ and dynamic observation on SmCo5 is carried out by using transmission electronic microscope and in heating environment that provides driving force for phase transformation in thermodynamics. In the view-point of equilibrium state thermodynamics, the change in outside condition makes above mentioned system reach phase transformation point and causes phase trans-formation to form new phase. The proceeding of this process is the result that the difference of enthalpy between new phases (Sm2Co17, etc.) and old phase (SmCo5) is less than zero but their absolute value is bigger than the resistance.

2.2.5 Discussion

SmCo5 permanent magnetic alloy possesses the maximum single axis anisotropy but the SmCo5 is unstable at below 805 and then eutectoid decomposition oc-curs promptly. It is testify by experiment that SmCo5 is eutectoid decomposed to be Sm2Co7 and Sm2Co17 phases by annealing at temperature ranges 300-750 . Also the magnetic crystal anisotropy of Sm2Co7 and Sm2Co17 is lower that that of SmCo5, thus scholars considered that minimum coercivity of SmCo5 appeared by annealing at 750 came down to Sm2Co7 and Sm2Co17 phases produced by eu-


tectoid decomposition. The experiment discovered that eutectoid decomposition of SmCo5 by annealing at 300-600 may cause precipitation of Sm2Co17 phase, but the coercivity of the alloy was not degraded remarkably and the coercivity of some specimen even was enhanced. Therefore, we took investigation emphases on eutectoid decomposition of Sm2Co17 phase itself for why SmCo5 appeared the minimum coercivity by annealing at 750 , and carried out theoretical calculation on precipitated Sm2Co17 phase and the in situ and dynamic observation on high temperature phase transformation with transmission microscope. We preceded the in situ and dynamic observation using 1000 kV HVEM and videotaped the obser-vation process. It was found through repeatedly studies that there were multi-defect areas in eutectoid decomposed Sm2Co17 phase at the temperature range of 650-750 and the amount of multi-defect areas reached the maximum at 750 . The magnetic anisotropy was the lowest in those multi-defect areas of Sm2Co17 phase through calculation. That is why the coercivity degraded to the lowest by annealing at 750 . The observed Sm2Co17 eutectoid decomposed from SmCo5

appeared as hexagon (actually three dimension hexagonal prism) and can be cal-culated by making use of the nucleation field Hn theory. Thus theoretic minimum of nucleation field may be calculated by formulae as below:

2M sH n3 32

a 2h = ( 3 3 a 2 + 6ah)& (2 .9)

The volume of regular hexagonal prism:

V = 3 32

a 2h � 2.5981a 2h

The surface area: S = 3 3 a 2 + 6ah � 5.1962 a 2 + 6ah

Substitute into Eq. 2.9 it is derived: 2.5981 a 2h×2M sH n =(5.1962 a 2 + 6ah )&

The result is: H n � (1/h + 1/a )v /M s (2.10) where & is energy density of domain wall; Ms is saturated magnetization intensity; h is the height of regular hexagonal prism; a is length of one side.

The &/M s of SmCo5 is treble of that of Sm2Co17 through calculation, thus Sm2Co17 is easily nucleated at given a, h. Supposing radius of antimagnetic nu-cleus is R it is calculated that Hn = 2,388 kA/m while R = 10 and Hn = 238.8 kA/m when R = 100 (Buschow, et al, 1980). The above calculation approximated the

length of diagonal of the regular hexagonal prism d = 2 24 4h a� to be the ra-

dius of a sphere so that the sphere radius will be R = d /2 =2 24

2h a� . It was


observed by electronic microscope that a considerable part of Sm2Co17 phase eu-tectoid decomposed from SmCo5 is also in sphere type so that the above supposi-tion is acceptable. It can be seen from the calculation that sphere radium of anti-magnetic nucleation increases by ten times the nucleation field H n will be de-crease to one of tenth. Defects in Sm2Co17 precipitated from SmCo5 took the most in amount when annealing at 750 and expanded continually. Thus R is bigger and H n is smaller. This indicated that the coercivity was degraded to the mini-mum by annealing at 750 because of the production and growth of antimag-netic nucleation centers in crystal granules. The nucleation centers were compara-tively very big, the defects increased and grown, and the interior magnetic anisot-ropy is very small so that based on micro magnetics the calculated coercivity of nucleation center in an integral Sm2Co17 precipitate phase should be one level over the actual measured coercivity. We may induce very reasonably that the Sm2Co17 phase itself is not the center of nucleation. It can be seen from the calcu-lation that the precipitate is more close to the Sm2Co17 structure the chance is big for being minus magnetic anisotropy. Again to consider the magnetic anisotropy of Sm2Co17 is one quantitative level lower than that of SmCo5 and in precipitated Sm2Co17 phase the cobalt sublattice possesses weak magnetizable direction in magnetic anisotropy of c plane. Therefore, the areas of considerable volume ex-tent with very low magnetic anisotropy may exist in some multi-defect Sm2Co17 phase, and these areas may become the centers of nucleation. The behavior of iH c in Fig.2.4 and Fig.2.5 can be explained as that the size of nucleation centers, with following the precipitated Sm2Co17 phase, grown dramatically in annealing stage and afterwards became slowness.

Referring to above mentioned it is summarized as follows: 1. It was discovered in the in situ and dynamic observation on SmCo5 specimen

in heating condition through 1000kV HVEM that the Sm2Co17 phase itself, precipitated from matrix phase of SmCo5, is not the antimagnetic nucleation center. The multi-defect areas appeared in the Sm2Co17 phase at temperature from 620-750 and reached the maximum amount at 750 . These multi-defect areas in the precipitated Sm2Co17 phase possessed very low magnetic anisotropy and acted as antimagnetic nucleation centers under effect of antimagnetic field so that resulted in degradation of coercivity.

2. SmCo5 permanent magnetic alloy was annealed at temperature from 200 to 1000 and then was quenched to room temperature. The coercivity of the al-loy did not degrade visibly and the coercivity in some specimens was even en-hanced slightly in measurement of the coercivity. It was found in dynamic obser-vation that a mass of Sm2Co17 phase precipitated at all temperature from 300 to 600 , and in meantime a small quantity of Sm2Co7 appeared. This indicated that the eutectoid decomposition of SmCo5�Sm2Co7 + Sm2Co17 occurred but the co-


ercivity did not varied remarkably. Thus it can be concluded that when annealing the alloy at temperature ranges from room temperature to 1000 the coercivity degraded drastically at 710-750 and reached the minimum value at 750 . That phenomenon did not attribute to eutectoid decomposed Sm2Co17 phase of low magnetic-crystal anisotropy.

3. The antielectron oblivion longevity of SmCo5, measured by antielectron oblivion technique, was a flat curve and basically unaltered. There was small pre-cipitate in a few minutes before, which indicated not to be vacant defects but is the other type of defect, i.e., substituent type defect. The observed defect with transmission electronic microscope at 650-750 was substituent type defect. The Sm3+ was substituted by Co due to incomplete precipitated Sm2Co17:

3SmCo5 – Sm + Co +, Sm2Co17 The and widened parameter S of Doppler curve did not vary with annealing

time, that there was not visibly vacant defect. 4. Experiments testified that the coercivity is not sensitive to enriched cobalt.

The coercivity is not influenced by surface but is influenced by pinning in interior crystal boundary. The coercivity did not varied basically on long time annealing time.

5. Magnetic anisotropy of various Sm3+ neighboring near the distribution center of Sm3+ varies along with changing of temperature. The Sm3+ is more close to the Sm2Co17 structure it would be more possibly to be a minus magnetic anisotropy.

6. It was observed at 650-750 that the location of striation Sm2Co7 phase transformed to very fine granular Sm2Co17 phase.

7. It was found in the in situ and dynamic observation on phase transformation of SmCo5 alloy at high temperature that nucleation speed increased with rising of temperature at temperature of 650-750 and the nucleation speed varied along with passing of time.

8. It was found in the in situ and dynamic observation on phase transformation of SmCo5 alloy at high temperature that part of the new phase spots of Sm2Co17 precipitated from matrix phase of SmCo5 grown up and coarsen at high tempera-ture ranges 650-750 but the other part of them re-dissolved into the matrix phase of SmCo5.

9. It was found in the dynamic observation on phase transformation of SmCo5 alloy at high temperature that the eutectoid decomposition of Sm2Co7 and Sm2Co17 from SmCo5 was a slow process and the above precipitated two phase may exist in the same time but the precipitation of Sm2Co17 is easier and its exis-tence time is longer in comparison with that those of Sm2Co7.

10. Phase transformation of SmCo5 is reversible at high temperature and that was testified by the in situ and dynamic observation experiments.

11. The following regularities are concluded through coercivity variation data


of 11 specimens of SmCo5 measured at 475-1000 , as listed Table 2.11: degrada-tion of the specimen with a higher coercivity will be smaller than the degradation of the specimen with lower coercivity at 750 after annealing. The specimen with higher coercivity will have less irreversible loss. The nucleation speed is related to the coercivity of specimen. The difference in performance of the alloy is related to difference in composition and manufacturing process. The difference in performance of specimen also influenced on variation of unit free energy of sys-tem and power for nucleation in phase transformation at high temperature.

2.3 Magnetism and the in Situ and Dynamic Observation of Permanent Magnetic Alloy of SmCo5 by Annealing at 600-1000 *

Investigators undertaking in study of magnetics and magnetic materials have at-tached importance to mechanism of the coercivity of SmCo5 permanent magnetic alloy. Their reason to do so is that the theoretical coercivity of the SmCo5 perma-nent magnetic alloy is 31,840 kA/m but in practice generally the coercivity of the SmCo5 permanent magnetic alloy produced in factory is only 1360-2036 kA/m, about seventeenth of the theoretical value. The SmCo5 produced in library has the maximum magnetic performance of 4776 kA/m (Zhou, et al, 1995), merely being sixth to seventh of the theoretical value. In addition mechanism of the coercivity can not be explained by conventional coercivity theory. Granule size of the SmCo5 is about D = 5-20 m and the size of its single domain is about D = 0.3-1.6 m, that the former is about seventeen times of the latter. According to theory of the single domain the coercivity should be proportional with the anisotropy constant K1 of magnetic crystals and the K1 of the SmCo5 should be lowered with lowering of temperature so that the mechanism of coercivity is not consistent with the theory of the single domain. Moreover, the coercivity valley appeared in SmCo5 permanent magnetic alloy by annealing at 750 , i.e., degraded to the minimum, but the coercivity rose at temperature from 750-950 (Pan, Ma, Li, 1993).

To study above mentioned problem clearly will enrich the coercivity theory and magnetics theory and promote development of the coercivity theory and magnet-ics theory.

The difference in this study from the method of previous studies is that the ob-servation was carried out dynamically using transmission electronic microscope under the condition of raising temperature from room temperature to 960 . This innovative method may systematically and precisely observe the relationship

Cooperators of this study include: Fengzuo Tian, Shikuan Ren, General Research Institute for Nonferrous Metal.


among the precipitation of new phase, phase transformation and the coercivity, so that it is better in continuity and systematism than the previous study in one or two temperature points using a massive specimen.

2.3.1 Specimen preparation and experimental method

SmCo5 sample was melted by alloy melting method. The sample was made of the nominal composition and melted in an arc furnace at protection atmosphere of argon. The alloy obtained was grinded by vibration ball miller and pulverized into powder of around 5 m. The powder was orientation formed under 1.5 T mag-netic field and then sintered. The sintered sample was sliced into 0.3 mm lamellae by a linear cutter and then thinned to films of about 100 nm by ionic thinning de-vice. The filmy specimen was placed into the side insertion heating dais of 1000 kV and observed dynamically by the transmission electronic microscope in heat-ing condition. In the meantime the observation was videotaped. The accelerating voltage was 1000 kV (Pan, Ma, Li, 1993).

2.3.2 Analysis on chemical composition of the SmCo5 permanent magnetic alloy

The chemical composition of SmCo5 Permanent Magnetic Alloy is shown in Ta-ble 2.3.

Table 2.3 Chemical composition of SmCo5 permanent magnetic alloy Composition/%

Alloy type Sm Co

Ratio of atoms Oxygen content/%

SmCo5-1 37 63 4.5 1.21 SmCo5-2 36.5 63.5 4.8 1.01

2.3.3 Magnetic measurement

The measurement was carried out by using CL6-1 magnetic parameter measuring instrument. The specimen was prepared as per composition of permanent mag-netic alloy listed in Table 2.4 and 4 samples were selected for measurement. The result of measurement of magnetic performance is shown as Table 2.4. The specimen was annealed in a heat treatment furnace at temperature from room temperature to 1000 , annealing at selected temperature of 250 , 420 , 500 , 600 , 750 , 850 , 900 , 950 , 1000 for one hour and quenching to room temperature, respectively. The specimen was measured in magnetic parameter measurement instrument. The result of the measurement is described as variation curve as shown in Fig. 2.10. It can be seen from the figure that the coercivity of SmCo5 permanent magnetic alloy appeared linear variation after annealing and


Table 2.4 Magnetic performance of SmCo5 permanent magnetic alloy

4Ms Br bHc mHc (BH)max Density Speci-men T kGs T kGs kA/m kOe kA/m kOe kJ/m3 MGOe g/cm3

1 1.03 10.3 0.87 8.7 532.4 6.63 1081 13.6 129.3 16.2 7.87

2 1.04 10.4 0.92 9.2 596.2 7.41 1096 13.7 159.1 20.0 7.57

3 1.05 10.5 0.95 9.5 673.1 8.32 1592 20.0 165.9 21.1 7.96

4 1.04 10.4 0.93 9.3 645.6 7.98 1085 13.8 161.2 20.5 7.76

Fig. 2.10 The coercivity of SmCo5 specimen after 1h annealing at

different annealing temperatures

reached the minimum at 750 . Afterwards the coercivity was restored until 900 .

2.3.4 Structure of magnetic domain

Using method of Kerr magnetic-optical effect to observe structure of magnetic domain obtained a figure of magnetic domain as shown in Fig. 2.11 to Fig. 2.13. The figures are photographs of structure of magnetic domain being magnified by 600 times and observed by Kerr magnetic-optical effect at room temperature.

The photograph shown in Fig. 2.11 is the pattern of planar magnetic domain in vertical to easily magnetisable axis. It can be seen from Fig. 2.11 that the volume of the positive domain equals to the negative domain and has no magnetism.

Fig. 2.11 Domain structure (600�) at thermal demagnetization


Fig. 2.12 Domain structure (600�) after magnetization with 3T magnetic field strength

Fig. 2.13 Domain structure (600�) of the optimized state after magnetization

with 3T magnetic field strength

The Fig. 2.12 shows photographs of specimen of permanent magnetic alloy af-ter magnetization under 3.0 T gauss magnetic field. The magnetization condition for Fig. 2.13 is the same as that of the Fig. 2.12. It can be seen from Fig. 2.13 that the volume of the positive domain decreased remarkably but the negative domain exhibited two traits: the first is the increase in number of the negative domain, the second is the widening of the negative domain. Pattern of the domain is like a labyrinthine.

2.3.5 Irreversible loss of SmCo5 permanent magnetic alloy after annealing at 25-1000

Experiment method was to hold the permanent magnetic alloy specimen for one hour at temperature 100 , 200 , 400 , 500 , 650 , 700 , 750 , 800 , 850 , 900 , 950 , 1000 , respectively, and quench the specimen to room temperature. Afterwards, their magnetism was measured and the irreversible loss was calculated. The measurement result was drawn to the variation curve as shown in Fig. 2.14 (comparison between magnetic performance before and after the annealing) (Pan, Ma, Li, 1993; Pan, Ren, Tan, 1983).

It can be seen from Fig.2.14 that its irreversible loss is very small when the an-


nealing temperature was below 500 . The magnetic irreversible loss becomes the biggest when the specimen was annealing at 750 for one hour. By annealing at 900-1000 for one hour the magnetic irreversible loss again increases gradu-ally and by annealing at 750-900 the magnetic irreversible loss decreases gradually.

Fig. 2.14 The curve of the magnetic irreversible loss of SmCo5 alloy

after 1h heating at 25-1000

2.3.6 Electronic energy spectrum experiment and analysis of SmCo5 permanent magnetic alloy

The electronic energy spectrum of SmCo5 alloy was measured at room tempera-ture. Conditions for measurement of AES was: energy of incident electronic beam 3 keV, the beam current 1 A, test voltage 60 V, multiplying voltage 1200 V, time constant 0.03 s, magnifying multiple 40 times and vacuum degree (2.66-3.99) × 10�5 Pa ((2-3)×10�7 Torr). Main measurement condition of XPS was: using radia-tion of magnesium target as light source, voltage 8 kV, electric current 30 mA and flux energy 50 eV.

Fig. 2.15 shown the experimental result of optical electron energy spectrum measured for SmCo5 specimen after annealing at different temperature from 25 to 900 as per above mentioned conditions. It can be seen from the figure that the proportion of samarium and cobalt changed along with rising of temperature so that the samarium atoms enriched area or cobalt atoms enriched area appeared in interior of the alloy. The element segregating of samarium and cobalt atoms caused the alloy to be SmCo5 phase partly and the other phases (Sm2Co7 and Sm2Co17) in the other part. No matter the Sm2Co7 or Sm2Co17 their magnetic crys-tal anisotropy are all lower than that of SmCo5. Inhomogeneous of the solid solu-tion magnetic crystal had a much lower anisotropy. This inhomogeneous area re-


duced gradually and the SmCo5 phase increased when the temperature was above 750 and as the result the coercivity measured after quenched was enhanced. When this inhomogeneous was reduced to zero the homogenous solid solution was quenched to room temperature as per proportion of SmCo5 phase and could hold the 1 : 5 phase. Then the coercivity could still be restored because the second phase could not be formed by element segregation but under an important premise that the oxygen should not be too high in the alloy. It can be seen from Fig. 2.15 that the oxygen did not increase along with rising of temperature so that the Sm2O3 was not formed within the alloy. An optic-electron energy spectrum ex-periment was designed to test the variation status of the elements of samarium, cobalt and oxygen. The specimen was annealing at high temperature of 750 (heating under protection gas) for one hour and then quenched to room tempera-ture. Afterwards, the optic-electron energy spectrum experiment was carried out. The distribution curve of elements of samarium, cobalt and oxygen of SmCo5

after annealing at 750 for one hour was described as shown in Fig. 2.16 in ac-cordance with peeling measured distribution of elements of samarium, cobalt and oxygen. The abscissa is the peeling time (min) and the ordinate is the atomic frac-tion. It can be seen from the figure that there was no peak value of oxygen.

Fig. 2.15 The study result of the surface of SmCo5 which contains Sm of 36.5%(wt.) by

means of photoelectron energy spectrum (It gives the distribution of Sm, Co, O, etc. The annealing temperature ranges from 25 to 900 )

To explain the problem of oxygen further two specimens of the samarium en-

riched SmCo5 alloy (containing high samarium content) were annealed at 750 for one hour. The optic-electron energy spectrum experiment of distribution of elements of samarium, cobalt and oxygen and the measured curve is shown in Fig. 2.16. The optic-electron energy spectrum experiment of distribution of elements of samar-ium, cobalt and oxygen and the measured curve is shown in Fig. 2.17 for speci-mens annealed at different temperature from 25 to 900 for one hour.

The rare earth elements are easy to combine with oxygen to form rare earth ox-ides (such as Sm2O3) because of activity of the rare earth elements. Therefore, to


Fig. 2.16 The study result of the surface of SmCo5 which contains Sm of 36.5%(wt.) by

means of photoelectron energy spectrum

Fig. 2.17 The study result of the surface of SmCo5 which contains Sm of 37%(wt.) by

means of photoelectron energy spectrum (It gives the distribution of Sm, Co, O, etc. The annealing temperature ranges from 25 to 900 )

make a judge is to check whether appearing of the peak value. It can be seen from Fig. 2.17 and Fig. 2.18 that there is no peak value of the oxygen.

Fig. 2.18 The study result of the surface of SmCo5 which contains Sm of 37%(wt.) by

means of photoelectron energy spectrum (The annealing treatment is 750 for 1 h)


2.3.7 The in situ and dynamic observation on eutectoid decom-position of SmCo5 by electronic microscope

The above mentioned SmCo5 experimental specimen was sliced into lamellae of 0.25 mm in direction perpendicular to easy magnetisable axis. Then the lamel-lae were thinned to about 100 nm by electrolysis polishing and ionic thinning. The specimen film was placed into the side inserting and heating dais of JEM-1000 and heated under condition of a vacuum degree of 266×10�7Pa (2×10�7 Torr) and an accelerating voltage of 1000 kV. Then the dynamic observation was carried out. The Fig. 2.19 was the electronic micrograph (the graph of bright field) and the electronic diffraction pattern in P area of electronic micro-graph being taken directly in thermal state inside of the electronic microscope when temperature of the filmy specimen was 500 . Through diffraction and calculation of the precipitated tilted striate texture and approximate round black piece, respectively, it was testified: the precipitated tilted striate texture is Sm2Co7; the approximate round black piece texture is Sm2Co17. This result indi-cated that the eutectoid decomposition existed in SmCo5 permanent magnetic alloy at 500 . The experiment also observed that the SmCo5 permanent mag-netic alloy eutectoid decomposed remarkably into phases of Sm2Co7 and Sm2Co17 at 400 , 600 and 750 , even appeared composition of the other phase.

Fig. 2.19 The electron micrograph (the left figure) of SmCo5 specimen at 500 (near round black area is Sm2Co17 and diagonal striation around is Sm2Co17 ) and

P area diffraction pattern (right figure)

0.1m


2.3.8 The in situ and dynamic observation of SmCo5 in thermal state using transmission electronic microscope

The phase precipitation was observed in the specimen of SmCo5 after holding at 750 for 10 min. This precipitate phase was found as Sm2Co17, Sm2Co7 and other phase texture with different composition through electronic diffraction. The obvious defects in the newly precipitated Sm2Co17 phase could be observed clearly after holding for 20min. The Fig. 2.20 (a), (b) shown these defects by magnified 100 thousands times in the texture of precipitated Sm2Co17 phase.

Fig. 2.20 The defect in Sm2Co17 precipitated from SmCo5 after 20min annealing at 750 (electron micrograph)

It was observed that the precipitated phase varied, dissolved, grown and aggre-

gated, drastically after holding 30 min. The Fig. 2.21 (a), (b) shown the elements segregation appeared in the precipitated Sm2Co17 phase by holding at 750 for 50 min. The black piece in the photograph was the precipitated Sm2Co17 phase. The element fluctuation and segregation of the precipitated Sm2Co17 phase were shown in Fig. 2.22 when holding for 60 min within the electronic microscope. It


can be seen from Fig. 2.22 that defects increased until to become a hole in the middle of the precipitated 2:17 phase. That some striation appeared in the figure is the other form of defect in the Sm2Co17 precipitated phase, and that to cause incomplete of the 2:17 phase and to enlarge segregation extent of phase composi-tion. It was found in the electronic microscope that something like fog form ap-peared when holding for 80 min, but this fog like things disappeared along with increase of holding time and rising of temperature.

Fig. 2.21 The electron micrograph of SmCo5 permanent magnet alloy after 50min

annealing at 750 (light field graph)

Fig. 2.22 The electron micrograph of SmCo5 permanent magnet alloy after 60min

annealing at 750 (light field graph)

2.3.9 The in situ and dynamic observation on SmCo5 in thermal condition of 750-960 by Transmission Electronic Micro-scope

The in situ and dynamic observation on SmCo5 specimen was carried out by hold-ing for 1h at 750 . Afterwards, the temperature was raised inside of electronic microscope from 750 to 960 to observe variation of the specimen directly in


1000 kV HVEM, seeing Fig. 2.23.

Fig. 2.23 The result of in situ and dynamic observation of the film of SmCo5 alloy (cut perpendicular to c axis) under 1000kV ultra high voltage electron microscope (heating at

950 for 1h) electron micrograph

The phase transformation phenomenon is as follows:

The fog state thing occurred at 750 disappeared; All vision of the microtexture pattern became very clear; The striation type Sm2Co7 phase disappeared and the precipitated phase did

not appeared after its disappearance; The inhomogeneous area in texture of the precipitated phase of Sm2Co17 dis-

appeared at 750 and became an homogenous precipitate phase, and its variation was in the following 5 procedures, seeing Fig. 2.24;

Fig. 2.24 The diagram of precipitation of Sm2Co17 from SmCo5 alloy annealing at 750-960

(a) inhomogeneous area in precipitated Sm2Co17; (b) merging of the inhomogeneous area in precipitated

Sm2Co17; (c) white point appears in precipitated Sm2Co17; (d) merging and partial broadening in precipitated

Sm2Co17; (e) merging in precipitated Sm2Co17 and outline left

Could not observe rhombus pillared dislocation and other typical disloca-tions which are easily to form the precipitated phase of Sm2Co17;

The interphase boundary (intergradations area) between the matrix phase and


the precipitated phase of Sm2Co17 disappeared at 750 in observation; The whitened area in the precipitated phase of Sm2Co17 and the white area in

the matrix phase were analyzed to be pure SmCo5 phase, and the X-ray diffraction peak of the alloy being quenched to room temperature as shown in Fig. 2.25.

Fig. 2.25 The peak of X-ray diffraction of SmCo5 alloy heating to 950 and then quick

cool down to room temperature under 1000kV ultra high voltage electron microscope

2.3.10 Discussion

1. The coercivity of permanent magnetic alloy is to indicate the magnetic density that magnetic induction intensity is zero on its static saturation Magnetic hystere-sis loop. The coercivity is the important parameter of permanent magnetic alloy and is the character for studying mechanism of magnetization reversal thus the study on coercivity may make various structures and nonstructural factors clearly instead. To analyze what factors which controls the coercivity of permanent mag-netic alloy thereby to control consciously this factors in manufacturing process so that to enhance the coercivity and improve the magnetic performance of the per-manent magnetic alloy (Li, Dai, 1982). Thus the scholars in study on the perma-nent magnetic alloy have paid great attention on the coercivity of SmCo5 and 2:17 type Sm-Co permanent magnetic alloys because the nonlinear variation of the SmCo5 alloy along with rising of temperature, that the coercivity of the alloy de-graded to the minimum at 750 when the alloy was annealed at temperature from 600-950 but appeared a peak value at 750-950 . This phenomenon could not be explained using conventional coercivity theory. Therefore, this study made use of advanced modern instruments and advanced methods and selected 1000 kV HVEM to observe the microtexture and structure of the phase precipitation and phase transformation. By using the HVEM we could observe the specimen of 100nm without any damage to the texture of specimen; could raise temperature from room temperature to 1000 in the electronic microscope and so that to en-able observation on phase transformation in the state of heat activation. It could observe directly in the electronic microscope that the Sm2Co17 phase, precipitated from the matrix phase of SmCo5, had some defects at 750 . The defects in this type of nanometer crystalline phase texture possesses two fold characters, the big


size defects is avail of nucleation and so that to degrade the coercivity; the small size defects is avail of acting as a pinning on the magnetic domain wall so that the defects may act as pinning point to enhance the coercivity. If supposing R is the radius of magnetization reversal nucleus, is the energy density of magnetic do-main, and Ms is saturation magnetization intensity, in condition of given magneti-zation reversal nucleus radius the theoretical critical value for nucleation field Hn can be calculated by formula as follows:

2Ms Hn (4/3)�R3 = 4�R2 The nucleation field:

Hn = (3/2) / (Ms R) For nucleation center of the Sm2Co17 phase if R = 10nm and 100nm the Hn can

be calculated by above formula as Hn = 30kOe and 3kOe, respectively. It can be seen from the calculation that the severe decrease of Hn is attributed to the size of 100 nm of nucleation center. The nucleation locus mostly occurred in the area with weak anisotropy or biggish demagnetization of the magnetic crystal. A little decomposition flakes of RE2Co17 in SmCo5 and cobalt enriched SmCo5 samples on surface of crystalline granule due to some reasons, all of those cobalt precipi-tated locus are nucleation point of magnetization reverse. If nucleation in defects can reduce energy of magnetic domain wall and increase demagnetization thus in order to obtain high coercivity the first thing is to control nucleation to avoid big size defects, especially to prevent cobalt couple from forming and prevent oxida-tion and aberrance. We found by 1000 kV HVEM in the in situ and dynamic ob-servation under state of heat activation that defects was largened obviously in the phase of Sm2Co17 precipitated from SmCo5 at 750 . The measurement and cal-culation indicated that samples appeared the biggest irreversible loss at 750 and some multi-defective areas with very low magnetocrystalline anisotropy existed in Sm2Co17 phase which became magnetization reverse nucleation centers so that the coercivity of SmCo5 permanent magnetic alloy was the minimum by anneal-ing at 750 .

2. For the reason of enhanced coercivity at temperature from 750 to 850 experiment fact provided this section explained as follows: the defects in Sm2Co17 phase precipitated at 750 and inhomogeneous area disappeared gradually along with rising temperature from 800 and 850 , that was dynamic observed with transmission electronic microscope. The disappearing process is shown in Fig. 2.22. The variation process of hexagonal precipitated Sm2Co17 phase in d area of Fig. 2.23 is shown as (a)-(e) in Fig. 2.24. Homogenous Sm2Co17 phase (without multi-defect area in the precipitated phase) was precipitated from the matrix phase of SmCo5 without degradation in coercivity but even appeared small peak value (shown in Fig. 2.10). It can be seen from Fig. 2.21 that there is a certain interphases between the matrix phase and the precipitated phase (transition area)


and the structure of this interphase or say transition area acts important function on magnetization reverse domain (thickness of interphase close to 0.5-1.0 m). As the composition and structure vary in the transition area the magnetism also varies so that this interphase causes big degradation in coercivity of the alloy. It can be seen from the figure that circumambience of the approximated round Sm2Co17 precipitated phase (the transition area between the matrix phase and Sm2Co17) disappeared so that the coercivity was enhanced. It can be seen from the optic-electron energy spectrum described in Fig. 2.15 to Fig. 2.18 that the short path segregation of samarium and cobalt elements appeared segregation junction at 750 and this segregation junction of samarium and cobalt elements reduced at temperature from 750 to 800 being close to the proportion of the samarium and cobalt elements gradually the at 600 (Pan, Ren, Tan, 1983). Thus this ex-periment explained why the coercivity enhanced again at 750 to 800 .

3. Why did the coercivity enhance again after the SmCo5 permanent magnetic alloy annealing from 750 to 800 but still be lower than the coercivity level at 600 ? This is because that the magnetic irreversible loss occurred in annealing, and from 750 to above 800 the process occurred that the inhomogeneous solid solution reduced gradually and homogenous solid solution increased gradu-ally. This is consistent with the equilibrium diagram of SmCo5. This type of proc-ess from inhomogeneous to homogenous is caused by heat activation and the SmCo5 phase becomes homogenous solid solution. This state was held by quench-ing, i.e., heat demagnetization sample for measurement. Its phase analysis by X-ray diffraction is shown as in Fig. 2.25 that the sample is basically pure SmCo5.

2.3.11 Conclusions

It concluded through above study: 1. The coercivity of SmCo5 permanent magnetic alloy showed a nonlinear

variation after annealing. The minimum of the coercivity was attributed to precipitation of Sm2Co17 phase with multi-defect area and the interphase (transition area) between the matrix phase and the precipitated phase with a very low magnetic anisotropy which then became the anti-magnetic nucleation center.

2. The enhancement in coercivity of the alloy at temperature from 750 to above 800 was because of the function of heat activation by which the multi-defect areas in Sm2Co17 phase precipitated from SmCo5 alloy and on interphase between the matrix phase and the Sm2Co17 phase disappeared, the solid solution with segregation of samarium and cobalt elements and inhomogeneous segrega-tion junction of elements became homogenous solid solution so that the high magnetocrystauine anisotropy of SmCo5 was restored.

3. The irreversible loss of SmCo5 permanent magnetic alloy measured after an-nealing was biggest at 750 in the process from 25 to 1000 . This is because


that the percent of homogenous solid solution in the mixture as per nominal com-position formed by thermal state was lowered to the minimum at 750 and the percent of inhomogeneous solid solution by segregation of samarium and cobalt elements became the maximum.

2.4 Analysis on Variation of the Coercivity and Phase Trans-formation *

Coercivity of SmCo5 permanent magnet presented nonlinear variation through annealing at different temperature. The outstanding specialty of this nonlinear variation is that the coercivity degraded drastically by short time annealing (heat treatment) at 700-750 and afterwards the coercivity was restored by heating at 900 , which was shown as the curve in Fig. 2.2 (Pan, Jin, 1990). The above spe-cialties attracted great attention of scholars.

Broeder and Smeggi, et al. believed that the degradation in coercivity of SmCo5 magnet by heating at 700-750 in short time was related to the eutectoid decom-position; but Kumar, et al. considered that heating at 700-750 in short time never could cause the eutectoid decomposition. The linear variation of the coer-civity was caused by the other reasons, in fact this nonlinear variation already appeared at below 600 but the variation was not so more remarkable as that in 700-750 so that it usually be neglected. This section is mainly to introduce the peak value of coercivity of SmCo5 permanent magnet measured in annealing at below 600 (Broeder, et al, 1972; Smeggil, et al, 1973; Kamar, Das, Wettstein, 1978; Pan, Tan, Jin, Yu, 1983).


All specimens were prepared by alloy melting method. The alloy was melted by arc furnace with the composition as per the proportion of metallic Sm:Co as 1:5 and among the row materials 11% was in liquid phase with the composition of samarium 60% and cobalt 40% in mass percentage. The melted alloy was crushed roughly and then pulverized into about 5m. The powder was oriented under a strong magnetic field and molded. Then the sample was ready for measurement after sintering.

The measurement instrument used in this experiment was magnetic parameter measurer. Measurement of the coercivity was conducted at room temperature. The experiments used JEM-1000 HVEM to heat and dynamic observe the speci-men. The specimen for electronic microscope was prepared as follows: sliced the

The people who have attend this test and discussion include: Shikuan Ren, Fengzuo Tian, Chengzhou Yu, Jiuyu Liu, General Reasearch Institute for Nonferrous Metal.


sample into lamellae of 0.03mm and afterwards being ionic thinned to about 100 nm; the shinned specimen was placed in enclosed into the side inserting type heating dais of JEM-1000 for heating and dynamic observation; the accelerating voltage was 1000 kV and electron beam current was 10 A.

2.4.2 Experimental result and discussion

The coercivity variation curve is showed as Fig. 2.2 for specimen after annealing at different temperature. The coercivity peak at below 600 for different speci-mens is showed as Fig. 2.26.

Fig. 2.26 The variation of coercivity of SmCo5 specimen annealing below 600

Magnetic loss is constituted with reversible loss and irreversible loss. Magnetic

performance can be restored through magnetizing (magnetization). Magnetic irre-versible loss can be divided into two parts, i.e., initial loss and odditional loss. Initial loss is the magnetic loss caused by short term aging but odditional loss is the magnetic loss caused by long term aging. In process of long term aging most of magnetic irreversible loss can be restored after re-magnetization so that only a small part of magnetic irreversible loss can not be restored.

Authors found in the experiment that the intrinsic coercivity was 2706– 3024kA/m but the magnetic irreversible loss of SmCo5 specimen after long term aging at 180 was only 1.4%-1.9%. And that 1.4% of magnetic irreversible loss is goods sample of the high coercivity and good quadrate degree (quadrate degree) of demagnetization curve, and 1.9% magnetic irreversible loss is a specimen of the coercivity of 2706kA/m and bad ridgy degree of the demagnetization curve. All SmCo5 used in the environment above 180 should pass this aging experi-


ment (stabilization treatment). It can be seen from above analysis that the peak value of coercivity is related to ridgy degree in the demagnetization curve of the intrinsic coercivity of specimens.

It can be seen from Fig. 2.26 that the measured coercivity of specimens ap-peared peak value at below 600 and appeared concave at 420 . After the coer-civity peak value of the magnet was found in above mentioned condition different specimens was measured as well and the result is shown in Fig. 2.26. How to ex-plain this phenomenon? The coercivity was varies very sensitively along with variation in microtexture and structure caused by phase transformation, to inves-tigate specialty of phase transformation and microtexture is the base to study co-ercivity mechanism of SmCo5 alloy. Thus using advanced microanalysis instru-ment enables directly observation in condition of heating under high vacuum en-vironment and that resulted in phase growing, extending and dissolving of phase at different temperature and mutual relation and interaction between geometrical shapes study of the phase and its structure. It is necessary to understand physical hypostasis of coercivity and above mentioned experiment phenomenon.

It is known from equilibrium diagram of Sm-Co (Fig. 2.27) and relative curve of Sm-Co composition versus free energy that the SmCo5 is comparatively stable at 800 but is metastable at below 800 .

Fig. 2.27 The phase diagram of Sm-Co

The eutectoid decomposition occurred in SmCo5 magnet at below 750 , i.e.,

SmCo5 � Sm2Co7 + Sm2Co17. There is big divarication for this problem as men-


tioned in the introduction, but the dynamic process of eutectoid decomposition in a special condition, phase growing, extending and dissolving in thermal state and geometric shape were revealed by dynamic observation experiment using and 1000kV HVEM with 0.204nm resolving power in General Research Institute for Nonferrous metals (Pan, Tan, Jin, Yu, 1983). At room temperature most of areas of SmCo5 specimen does not have precipitated phase but once in a while the de-fect like stowage layer dislocation and very small size precipitated phase could be observed, as shown in Fig. 2.28. The specimen was samarium enriched SmCo5

sinter (including 37.66% (wt.) Sm) being sintered at 1150 and quenched from 900 . This type of specimen not only had defects but even also had pinhole in some small areas, as shown in Fig. 2.29. Fig. 2.29 showed photograph of mag-netic domain structure of specimen at room temperature observed by Korr optic-magnetic effect (600×). It can be seen from Fig.2.29 that pattern of the magnetic domain located in middle of the pinhole.

Fig. 2.28 The electron micrograph of SmCo5 at room temperature

Fig. 2.29 The domain pattern of the specimen magnetized at room temperature (600�)

After the specimen was heated to 420 in the 1000 kV transmission electronic

microscope and held for 2 min the specimen in the area originally without any precipitate could see clearly highly dispersed precipitate phase fulfilled the all


vision. After holding 10 min continually the precipitated phase grown to the size of teens nanometers and some precipitate phase already aggregated, as shown in Fig. 2.30 (the black precipitate was analyzed to be Sm2Co17 by electronic diffract- tion, R and Q in the Fig. 2.30(a) already aggregated together in the Fig. 2.30(b)). The electronic diffraction photograph, electronic diffraction pattern in selected area of precipitate phase and indexation are shown in Fig. 2.31 for SmCo5 speci-

Fig. 2.30 The upgrowth, coarsening, gathering and joining of the precipitation of

Sm2Co17 in SmCo5 at 420

Fig. 2.31 The precipitation of Sm2Co17 and Sm2Co7 in SmCo5

(a)The electron micrograph of SmCo5 at 420 ; (b) The selected-area electron diffraction pattern of Q area (precipitated phase) ; (c) Indexing of the diffraction pattern in (b)


men at 420 . The analysis of electronic diffraction of SmCo5 specimen at 420 indicated that the precipitate phase was Sm2Co7 and Sm2Co17.

For which one of Sm2Co7 and Sm2Co17 had ever precipitated at first Fidler had made systemic investigation on interaction between dislocation in SmCo5 and the second phase. He supposed at first that the base substance had the same elasticity as that of the precipitate phase and both were isotropic when precipitate phase appeared as oblate elliptical spherical the elastic aberrance energy of a cell in the precipitate phase could be derived.

The elastic aberrance energy of Sm2Co7: e1G = (0.043 + 0.174 Y / Z )

The elastic aberrance energy of Sm2Co17: e2G = (�0.908 +0.1736 Y / Z )

The Z and Y respectively indicate the short and long radius of the oblate ellipti-cal ball precipitate phase in the formulae. In comparison between the e

1G and e2G it can be seen that e

1G > e2G . Thus regarding the elastic aberrance energy the

precipitation of Sm2Co17 is more favorable so that Sm2Co17 phase precipitated firstly in SmCo5.

When the temperature was heightened to 500 many precipitate phase were still in stable state. At above 600 the imperfection in the precipitate phase could be observed obviously, which in Sm2Co17 may be caused by concentration fluctu-ating and short path clustering of the matrix phase. This section does not intend to treat of the relation between phase precipitation and coercivity but emphases on discussion the peak value of coercivity measured at room temperature after an-nealing at below 600 .

We can find without difficulties from the above mentioned observation by elec-tronic microscope that the defect, pinholes the specimen in thermal state due to precipitation, growing up, , and quenching of Sm2Co7 and Sm2Co17 phases are not the same as those of specimen without annealing in size and distribution. Then does what king of the influence of the newly formed defects impose on the coercivity. The original defects disappeared thus to influence on pinning of crystal boundary on the wall of magnetic domain. If the newly formed defect size is big-ger than or equal to the wide of magnetic domain wall, the nucleation center will be nearby the defect and the defect will only be localized there so that this kind of defect will restrict expanding of the anti-magnetization nucleus. This kind of pin-ning may enhance the coercivity so that the peak value appeared in measurement. The newly formed defect when the specimen was heated again and quenched to room temperature would be different from the defect before. If this kind of defect is avail of restricting e expansion of anti-magnetization nucleus then the coerciv-ity would be enhanced further and the peak value be increased (Pan, Tian, Jin, Yu,


1983; Yang, 1979). If just in opposition then the coercivity will be degraded so that become a concavity. The different specimen appeared different defects after quenching from 900 due to technical diversity thus as the result not all of the appearance time of the coercivity peak is the same in annealing treatment. Thus annealing at a special temperature (such as just select the temperature at which mH c peak value appears) is avail of enhancement of mH c .

2.4.3 Conclusions

Experiments indicated that samarium enriched SmCo5 magnet sinters in all had a peak value in the variation curve of coercivity versus temperature being measured at room temperature after annealing at different temperature below 600 . It was ob-served directly the matrix phase of SmCo5 eutectoid decomposed into Sm2Co7 and Sm2Co17 phases at 400-600 in heating dynamic observation with 1000 kV HVEM.

2.5 The Optic-electronic Spectrum Study on SmCo5 Perma-nent Magnetic Alloy

SmCo5 permanent magnetic alloy not only possesses high magnetocrystauine ani-sotropy but also has comparatively high saturated magnetization intensity and Curie temperature. Thus it has been developed to a very widely applied perma-nent magnetic material since the first batch SmCo5 permanent magnetic powder material was made by Strnat et al. in 1967.

In recent year applications asked continuously for higher requirements for magnetic performance of SmCo5. Oxidized layer on surface of SmCo5 magnet also become the problem the users concerned. They hope that the problem can be overcome in material research and facture.

In order to obtain high coercivity in material research and facture firstly is to con-trol nucleation, to control nucleation then need to avoid defect, and to avoid defect thus need to pay attention to prevent the surface from aberrance and oxidation.

Since many years ago people have researched the oxidation problem of SmCo5 permanent magnetic alloy in manufacturing process by different aspects and achieved important outcomes. P.J. Jorgenson pointed out: oxidation was very quickly in interior of SmCo5 so that the magnetic performance of the alloy was damaged in the oxidized area and the structural variation extent caused by this oxidation might be used to predict the irreversible loss of SmCo5 after long term aging (Jorgensen, 1981). B. Labal, et al. pointed out in reference that the interior oxidation of the rare earth cobalt alloy already became the object to be researched by metallography, electronic probe and other technique; the interior oxidation of CeCo5-x-Cux compound (0 x 2) was optional and as the result it was to form CeO2 structure and a Co-rich phase (Labulle, Petipas, 1980).


In order to find out further the oxidized layer formed on sintered SmCo5 oxi-dized specimen some researches were carried out to study the oxidized layer on SmCo5 surface using surface analysis instrument, that was purposed to study the state of elements and atomic concentration variation of elements of samarium, cobalt and oxygen in different depth of the oxidized layer which affected mag-netic performance of the material. And in what chemical state of oxygen was and in what state the samarium and cobalt existed in the oxidized layer.

2.5.1 Specimen preparation technique and experimental condi-tion of optic-electronic energy spectrum

Specimen used in experiment was made using alloy melting method: the alloy composition was prepared per atomic proportion of 1:5 for samarium and cobalt metals with purity 99.5%; melting was curried out using arc furnace. In melting drawn out vacuum at first and then fulfilled argon. To ensure alloy composition to be homogenized the alloy was melted thrice continually. About 11% of the main phase alloy was added into the melted liquid alloy with a composition of Sm 60%, Co 40%. The alloy was crushed coarsely and then pulverized to around 5m in vibrant ball mill and then the powder was oriented under 1.5T magnetic fields and molded. Afterwards, the alloy was sintered in low vacuum condition. Sintered alloy appeared obviously in brown-yellow color on the moment out of furnace. The color turned to black after a period of time. After a long time the alloy ap-peared a white oxidized layer, of which the thickness varied along with oxidation extent. Experimental specimen contained Sm 35.3% and its magnetic perform-ance was measured to be Br = 0.56T, bHc = 191.36 kA/m. It could be seen that magnetic performance was severely damaged due to oxidation (Pan, Jin, 1990; Pan, Liu, Luo, 1990).

Surface analysis instruments used in experiment were Auger spectrum appara-tus (AES) and optic-electronic energy apparatus (XPS). Condition for AES meas-urement: energy of incident electron beam was 3 keV, beam current was 1A, test voltage was 60 V, amplifying voltage was 1200 V; time constant was 0.03 s; mag-nifying multiple was 40 times; vacuum degree was (266-399)×10�7 Pa((2-3)×10�7

Torr). Main condition for XPS measurement was: with magnesium target radia-tion as light source, voltage 8 kV, current 30 mA and flux energy 50 eV (Pan, Ma, Li, 1993; Pan, et al, 1987).

2.5.2 Investigation on surface composition of SmCo5

Surface composition measurement used Auger electronic energy spectrum and the experimental condition as mentioned before. Fig. 2.32 shows AES spectrum be-fore and after peeling by hydrogen ion at room temperature: curve a indicates the


AES spectrum before peeling by hydrogen ion; curve b indicates the AES spec-trum after hydrogen ion peeling for 1min; curve c indicates the AES spectrum after hydrogen ion peeling for 2min.

Fig. 2.32 The photoelectron energy spectrum (AES) of SmCo5 at room temperature

It can be seen from Fig. 2.32 that: impurities on surface of specimen mainly in-clude carbon, oxygen, etc., that peaks of samarium and cobalt comparatively be-ing weakened may be due to contaminated surface by carbon; after peeling 1min to remove carbon the peaks of samarium and cobalt reveal their original form. After peeling for 2min the surface carbon was removed completely and oxygen peak was reduced comparatively by its peak was still relative strong. This indi-cates that the oxygen exists on surface but considerable part of it exists as com-pound in deep of surface layer.

2.5.3 Atoms concentration variation of elements of samarium, cobalt and oxygen from surface to depth

Atoms concentration variation of elements of samarium, cobalt and oxygen from surface to depth is showed in Table 2.5.

Table 2.5 Concentration change of elements Sm, Co and O from surface to depth

Atom concentration fraction/% Ratio of atom concentration of elements Peeling time/min

Sm Co O Sm/Co Sm/O Co/O

4 41.0 3.7 55.3 11.08 0.74 0.07

11 42.0 5.6 52.4 7.78 0.80 0.11

13 42.0 6.7 51.3 6.27 0.82 0.13

Data in Table 2.5 indicates: cobalt atoms concentration fraction increased from surface layer to interior of surface oxidized layer but oxygen atoms concentration


fraction decreased; although samarium concentration fraction did not varied big, the rate of Sm/Co declined yet while Sm/O correspondingly increased.

2.5.4 Surface compound

Surface compound was measured using XPS. Its measurement condition as de-scribed before in experimental method. Five lines in each spectrum map are listed from low to high in emission intensity as the spectrums from surface, peeling for 2min, 4min and 13min, respectively.

Comparing the actual measured XPS spectrum map with standard XPS spec-trum map that the actually measured distance between samarium (3d 5/2) and (3d 3/2) is 28.0 eV and the standard is 27.2 eV thus the result is corresponding. This standard samarium compound is Sm2O3. This indicated that Sm2O3 formed by samarium and oxygen is a stable tervalent state thus samarium on specimen sur-face is in oxidized state. The standard binding energy of oxygen (1s ) is 531.6 eV and its compound is M2O3; and that the actual measured binding energy is 531.8 eV. This result indicates that oxygen on specimen surface chemically combined with metals to form compound in �2 valences; the standard combination energy of cobalt (2p 3/2) is 777.9 eV and the actually measured value of the specimen was as 777.8 eV. Cobalt is in null valence state. The distance between cobalt (p 3/2) and cobalt (2p 2/1) is 15.05 eV, the actually measured distance between peaks is 15.0eV so that the result corresponds with each other. Thus cobalt of the specimen in this experiment is null valence, i.e., is a metal cobalt, at measurement spot.

2.5.5 Conclusions

1. Oxidized layer of the oxidized SmCo5 specimen in sintering process is Sm2O3. 2. It can be seen from concentration percentage of elements of samarium, co-

balt and oxygen from surface to depth of the oxidized layer of the specimen that cobalt atomic concentration fraction correspondingly increased but oxygen atomic concentration fraction correspondingly decreased; Sm/Co reduced correspond-ingly and Sm/O increase correspondingly.

3. Cobalt of this specimen in the measuring spot is null valence.

2.6 Analysis on Magnetic Hysteresis Loop of SmCo5 Perma-nent Magnetic Alloy *

Observation on magnetic hysteresis loop of material and study the reason for

The people who attended this test include: Shikuan Ren, Fengzuo Tian of General Research Institute for Nonferrous Metal and Fuming Yang of Institute of Physics Chinese Academy of Sciences.


various shape of the loop is a way to analyze difference of material and seek a approach to improve performance of materials. However, to obtain an exact mag-netic hysteresis loop normally need a magnetic measurement apparatus with mag-netic field intensity over 100,000 Gs. This apparatus was lacked and some had bad exactness before in continent of China. Therefore, we used the newly devel-oped strong magnetic field apparatus with magnetic field intensity near 200,000 Gs in physics institute to measure the magnetic hysteresis loop and the curve of the loop varying with temperature for three types of specimens (Yang, Zhao, Li, et al, 1983). This section emphasizes to introduce the measurement result and dis-cuss the result with combination of technical condition and observation result by electronic microscope.

2.6.1 Specimen preparation technique, magnetic measurement and transmission microscope condition and experimental method

All specimens used in experiment were prepared using alloy melting method (Pan, Jin, 1990): the alloy composition was prepared per atomic proportion of 1:5 for samarium and cobalt metals; melting was curried out using arc furnace. The melted liquid alloy with a composition of Sm 60%, Co 40% was used. The adding amount was 11% of total amount of the alloy. The melted alloy was crushed coarsely and then pulverized to around 5m. The powder was oriented under a strong magnetic field and molded. After sintering the specimen was ready for measurement.

Instrument used for measurement was a strong magnetic field apparatus in Institute of Physics Chinese Academy of Sciences with a pulse magnetic intensity up to 200,000 Gs. The magnetic field wave was close to a half sine wave and its durative time was 15.9ms and was switch-controlled. The magnetic field could be reversed to adapt the measurement of the magnetic hysteresis loop. Magnetic field H and magnetization intensity 4�M was measured using induction method. Measurement scope was 77-550K.

Experiment used JEM-1000 HVEM for dynamic observation on and heating specimen. Preparation method of specimen was sliced the alloy sample into 0.3 mm lamellae using a linear cutter; and then thinned the sample to about 100nm using ion; its accelerate voltage was 1000 kV and its current was 10A.

2.6.2 Analysis on chemical composition of three kinds of speci-mens

Fig. 2.33 to Fig. 2.35 show the magnetic hysteresis loops being measured of three types of different specimens (classified as three types, i.e., , and ).


Fig. 2.33 The magnetic hysteresis loop of SmCo5 specimen



In order to explore difference among three magnetic hysteresis loops the fol-

lowing experimental analysis and discussion was made, seeing Table 2.6.

Table 2.6 Analysis on chemical composition of three types of different samples Specimens Chemical composition/%

Type No. Sm Co Ratio of atom Oxygen content/%

1 35.80 64.20 4.57 1.08

2 37.92 62.08 6.17 1.85

3 38.13 61.86 4.14 0.67


It can be seen from comparison between Fig. 2.33 and Fig. 2.34 that the mag-netic hysteresis loop of the type (No.2) had a wasp waist shape and a bad quadrate degree. It can be seen compared with Table 2.6 that oxygen content of the specimen No. 2 was the highest.

2.6.3 Analysis on preparation technique

In preparation technique the type (No.3) specimen had a higher oriented field thus its quadrate degree of the loop was better than that of the type .

2.6.4 Curve of magnetic performance and analysis at 77-550K

The magnetic hysteresis loop was measured using 200,000 Gs magnetic field measurement apparatus at different temperature from 77K to 550K. The meas-urement obtained a variation curve of jH c vs T and 4�Mr vs T, as shown in Fig. 2.36 and Fig. 2.37.

Fig. 2.36 The variation of jHc of SmCo5 vs temperature

Fig. 2.37 The variation of 4�Mr of SmCo5 vs temperature


It can be seen from Fig. 2.36 and Fig. 2.37 that 4�Mr did not vary considerably along with rising of temperature. The coercivity variation speed of specimen type

(No.2) was bigger than that of type (No.1) and type (No.3).

2.6.5 Observation and analysis on specimen using TEM

Observation using HVEM was carried out on microtexture of three types of specimens. Fig. 2.38 shows texture, dislocation network and precipitate particles of the specimen with a worst quadrate degree of magnetic hysteresis loop at room temperature. Fig. 2.39 is the microtexture, precipitate phase and dislocations of the specimen of type (No.3) with a better quadrate degree of magnetic hystere-sis loop at room temperature. It is known from the observation by the microscope

Fig. 2.38 Dislocation and precipitation particle in microstructure of SmCo5 specimen

(No.2) at room temperature

Fig. 2.39 Microstructure of SmCo5 specimen (No.3) at room temperature,

dislocation and precipitation phase


that the specimen of type (No.2) had more precipitate phase than that in the specimen of type (No.3), i.e., had more precipitate which are not 1:5 phase (the specimen of type (No.3) was the 1:5 phase being testified through elec-tronic diffraction analysis) (Pan, Jin, 1990; Pan, Zhao, 1989).

2.6.6 Conclusions

Through measurement of magnetic hysteresis loop and magnetic parameters, and observation using 1000 kV HVEM it is concluded: to obtain material with good quadrate degree of magnetic hysteresis loop and high magnetic performance it is necessary to decrease oxygen content, increase magnetic field intensity in powder molding and strictly control sintering process to avoid precipitation of none 1:5 phase besides paying attention to components of raw material.

2.7 Magnetism of SmCo5 Permanent Alloy at 1.5-523 K

Magnetism measurement of rare earth permanent magnetic alloy was to place the specimen firstly under magnetic field being above the 4�Ms of the specimen and then the specimen was magnetized to saturation. Afterwards the magnetized specimen was placed into electromagnetic iron of magnetic measurer to measure demagnetization curve using magnetic parameter measurement apparatus (Zhang, Cheng, Zhang, 1982). The demerit of this magnetism measurement method is unable to measure intrinsic magnetism of the specimen with a magnetic field in-tensity above 2.1 T because normal electromagnetic iron can only produce a mag-netic field intensity about 2.0 T. Therefore, for a rare earth permanent magnetic alloy with an intrinsic coercivity above 2.1 T this measurement can not describe its whole intrinsic magnetic measurement curve. In order to study magnetic per-formance of rare earth permanent magnetic alloy at low temperature add to widen application scope of the rare earth permanent magnetic alloy, we adopted ad-vanced instrument with ultra strong magnetic field produced by superconductive magnet which is vibrant specimen magnetic intensity apparatus to measure and study the magnetic speciality. This section provides magnetism of rare earth per-manent magnetic alloy at 1.5-523 K and one type of magnetism measurement and study method to study magnetism of magnetic materials from high temperature to deep low temperature and to avail application of the material.

2.7.1 Specimen preparation technique, magnetic measurement apparatus and experimental method

This experiment prepared a few types of specimens with the specific composition


(in mass fraction) were 36.9% Sm + 63.1% Co (for specimen No.1) and 37.5% Sm + 62.5% Co (for specimen No.2). Manufacturing technique was as follows: enclosed the prepared material into arc furnace for melting; the ingot after melting was milled from coarse, middle and fine to about 5m; the powder was molded under magnetic field and sintered at 1150 for 1.5h; afterwards the temperature was lowered to 950 and held for 1h, and quenched to room temperature; then the specimen was available for measurement (Pan, Jin, 1990; Pan, Ma, Li, 1993; Pan, 1996).

Magnetic measurement instrument and experimental method: vibrant specimen magnetization intensity apparatus for measurement at 1.5K, 40K; magnetism pa-rameter measurement apparatus of rare earth magnet made in Baotao Rare Earth Research Institute for measurement at temperature from �196 to 250 , meas-urement condition 50Hz single cycle pulse magnetic field.

2.7.2 Magnetism measurement and curve of SmCo5 permanent magnetic alloy at 1.5 K and 40 K

1. Measured curve using magnetic parameter measurement instrument at room temperature (25 ) is shown in Fig. 2.40. The magnetic performance is as follows: specimen of 36.9% Sm (wt.) + 63.1% Co (wt.) (specimen No.1), measurement data are Br = 0.92 T, mHc = 1592.00 kA/m, bHc = 716.40 kA/m, (BH)max = 167.16 kJ/m3.

Fig. 2.40 Demagnetization curve of SmCo5 permanent magnet alloy (25 )

2. Measuring result in ultra conductive magnetic field using vibrant specimen magnetic intensity apparatus: measured M = 19.36 emu/g at temperature of 1.49 K in center of apparatus; intrinsic coercivity mHc = 30,009 Oe (2388.72 kA/m), its curve refers to Fig. 2.41, its performance refers to Table 2.7.


Fig. 2.41 Demagnetization curve of SmCo5 permanent magnet alloy at

temperature of 1.5K

Table 2.7 Magnetism of SmCo5 at 1.5K (1.49-1.55K)

Havant/Oe M Mmoy/emu�g�1 Tbas/K T/K H/Oe

100 �77.757 �77.757 1.66 1.55 99

395 �78.043 �78.055 1.66 1.55 395

701 �77.929 �77.929 1.66 1.55 701

1,985 �77.400 �77.400 1.65 1.54 1,983

3,998 �76.466 �76.471 1.64 1.53 3,996

5,983 �75.519 �75.528 1.64 1.53 5,981

8,003 �74.548 �74.548 1.64 1.53 8,001

9,985 �73.073 �73.073 1.63 1.52 9,986

11,990 �70.994 �70.993 1.63 1.51 11,994

14,003 �67.449 �67.445 1.62 1.51 14,007

15,994 �63.276 �63.280 1.62 1.50 15,998

18,001 �57.608 �57.601 1.62 1.50 18,005

22,011 �43.903 �43.903 1.61 1.49 22,016

23,999 �33.195 �33.195 1.61 1.49 24,000

26,011 �17.195 �17.195 1.61 1.49 26,016

28,013 1.700 1.700 1.61 1.49 28,009

30,011 19.360 19.360 1.60 1.49 30,009

3. Measurement data for specimen No.2 with composition of 37.5% Sm (wt.)

and 62.5% Co (wt.) at temperature 1.5K: temperature in center of instrument Tcent = 1.48 K, M = 33.483 emu/g, measured intrinsic coercivity was 37,975 Oe (3,022.81 kA/m); measurement curve refers to Fig. 2.42; performance refers to Table 2.8.


Fig. 2.42 Measuring curve of SmCo5 specimen No.2 at temperature of 1.5K and 40K

Table 2.8 Magnetism of SmCo5 at 1.5K (1.49-1.51K)

Havant/Oe M Mmoy/emu�g�1 Tbas/K T/K H/Oe

105 �79.113 �79.113 1.69 1.51 105

402 �79.448 �79.448 1.70 1.51 401

702 �79.288 �79.303 1.69 1.51 701

1,982 �78.729 �78.715 1.69 1.51 1,981

3,981 �78.251 �78.251 1.69 1.51 3,980

5,996 �77.973 �77.973 1.69 1.51 5,996

7,987 �77.715 �77.710 1.69 1.51 7,986

9,986 �77.357 �77.357 1.67 1.51 9,989

12,000 �76.713 �76.717 1.67 1.50 12,004

13,980 �75.741 �75.741 1.68 1.50 13,984

15,998 �74.021 �74.032 1.62 1.50 16,003

18,000 �70.868 �70.868 1.66 1.50 18,004

20,000 �65.846 �65.855 1.64 1.50 20,005

21,987 �59.039 �59.046 1.64 1.50 21,991

23,996 �50,701 �50.701 1.66 1.50 23,996

26,012 �40.643 �40.630 1.62 1.49 26,015

28,006 �29.210 �29.210 1.62 1.50 28,002

30,012 �16.569 �16.569 1.66 1.49 30,010

32,009 �3.246 �3.246 1.60 1.49 32,004

33,975 9.855 9.855 1.60 1.49 33,974

35,978 22.350 22.350 1.61 1.49 35,974

37,977 33.483 33.483 1.60 1.49 37,975

2.7.3 Measurement of demagnetization curve and value of mag-netic parameter at �196-250 by magnetic parameter measurement apparatus

Measurement sequence:


25 � 100 � 150 � 200 � 250 � 25 � �60 � �196 � 25

Fig. 2.43 showed demagnetization curve at 7 different temperatures. The meas-urement was carried out as per sequence from room temperature (25 ) to high temperature (250 ), return to room temperature (25 ), and then to low tempera-ture (�196 ) and again return to room temperature (25 ). The measured de-magnetization curves can be superposable in all, which indicated that there was not any magnetic irreversible loss caused by variation of temperature in above mentioned temperature condition.

Fig. 2.43 Demagnetization curves of SmCo5 permanent magnet alloy at �196-250 Magnetism of three SmCo5 specimens at �196-250 refer to Table 2.9.

Table 2.9 Magnetic performance of SmCo5 at temperature range from �196 to 250

Values Specimens Magnetic

parameter �196 �60 25 100 150 200 250

Br/T 0.86 0.85 0.82 0.8 0.78 0.75 0.73

bHc/kA�m�1 668.6 652.72 612.92 581.08 549.24 533.32 477.6

mHc/kA�m�1 1830.8 1544.24 1273.6 1082.56 923.36 780.08 668.64 SmCo5

1

(BH)max/kJ�m�3 147.26 140.41 133.73 127.36 121.07 108.97 103.48

Br/T 0.97 0.94 0.93 0.92 0.86 0.83 0.73

bHc/kA�m�1 708.44 628.84 537.3 477.6 437.8 374.12 302.48

mHc/kA�m�1 859.68 700.48 589.04 493.52 429.84 366.16 310.44 SmCo5

2

(BH)max/kJ�m�3 194.86 178.62 167.16 157.29 143.28 129.9 96.95

Br/T 0.98 0.95 0.91 0.89 0.87 0.86 0.84

bHc/kA�m�1 756.2 732.3 724.4 636.8 597.0 541.2 469.6

mHc/kA�m�1 2547.2 1958.16 1528.32 1289.52 1146.24 971.12 875.6 SmCo5

3

(BH)max/kJ�m�3 191.04 182.28 167.16 159.2 151.24 143.28 135.32


2.7.4 Reversible temperature coefficient of SmCo5 at �196-250

Reversible temperature coefficients of SmCo5 at �196-250 are showed in Table 2.10.

Table 2.10 Reversible temperature coefficient of SmCo5 at �196-250

Reversible temperature coefficient, � / % �1 Specimens Magnetic

parameter �196 �60 100 150 200 250 Br/T 0.022 0.044 0.041 0.039 0.049 0.048

bHc/kA�m�1 0.041 0.076 0.069 0.083 0.074 0.098

mHc/kA�m�1 0.198 0.250 0.200 0.220 0.022 0.211 SmCo5

1

(BH)max/kJ�m�3 0.046 0.059 0.063 0.076 0.106 0.101

Br/T 0.024 0.031 0.011 0.063 0.045 0.047

bHc/kA�m�1 0.018 0.031 0.105 0.084 0.068 0.058

mHc/kA�m�1 0.183 0.238 0.194 0.238 0.170 0.211 SmCo5

2

(BH)max/kJ�m�3 0.047 0.058 0.207 0.160 0.130 0.111

Br/T 0.039 0.015 0.017 0.030 0.029 0.034

bHc/kA�m�1 0.053 0.110 0.018 0.021 0.030 0.060

mHc/kA�m�1 0.217 0.320 0.173 0.191 0.193 0.200 SmCo5

3

(BH)max/kJ�m�3 0.104 0.142 0.050 0.037 0.050 0.060

2.7.5 Coercivity of SmCo5 at 475-1000

Coercivity of SmCo5 was measured using magnetic parameter measurement appa-ratus after annealing at 475 , 500 , 650 , 750 , 850 , 900 and 1000 , the measured coercivity value is listed in Table 2.11. It can be seen that coercivity of all specimens are the lowest at 750 , and enhanced again at 850 , 900 and 1000 ; the coercivity of specimens No.3 and No.6 appeared higher at a given temperature and higher comparatively also at 750 . The experiment data indi-cates that the coercivity is the lowest at 750 at temperature range from 475 to 1000 . The degradation of Hc is related to coercivity of the specimen itself.

Table 2.11 Coercivity value of SmCo5 at 475-1000 Values

Specimens Coercivity 475 500 650 750 850 900 1000

Oe 7,450 7,100 6,250 4,000 5,800 6,500 6,000 bHc

kA/m 593.02 565.16 497.5 318.4 461.68 517.4 477.6

Oe 15,000 16,000 9,900 4,800 19,200 17,500 12,400 SmCo5

1 mHc

kA/m 1194 1273.6 788.04 382.08 1528.32 1393 987.04


Continued Table 2.11

Values Specimens Coercivity

475 500 650 750 850 900 1000

Oe 6,400 6,500 4,300 4,800 6,700 6,400 4,500 bHc

kA/m 509.44 517.4 342.28 382.08 533.32 509.44 358.2

Oe 8,200 8,800 4,800 5,600 11,000 8,600 5,600 SmCo5

2 mHc

kA/m 652.72 700.48 382.08 445.76 875.6 684.56 445.76

Oe 7,000 7,000 5,400 5,000 6,750 6,500 6,000 bHc

kA/m 557.2 557.2 429.84 398 537.3 517.4 477.6

Oe 31,000 31,600 10,800 11,000 24,000 29,500 33,200 SmCo5

3 mHc

kA/m 2467.6 2515.36 859.68 875.6 1910.4 2348.2 2642.72

Oe 7,400 7,400 6,200 1,500 3,600 6,100 6,200 bHc

kA/m 589.04 589.04 493.52 119.4 286.56 485.56 493.52

Oe 16,000 16,600 9,500 1,500 4,400 17,200 13,400 SmCo5

4 mHc

kA/m 1273.6 1321.36 756.2 119.4 350.24 1369.12 1066.64

Oe 7,600 7,300 6,200 1,100 2,800 5,500 4,800 bHc

kA/m 604.96 581.08 493.52 87.56 222.88 437.8 382.08

Oe 19,600 18,200 11,000 900 3,150 14,500 8,000

SmCo5

5 mHc

kA/m 1560.16 1448.72 875.6 71.64 250.74 1154.2 636.8

Oe 7,200 7,100 6,200 5,000 6,200 6,200 6,000 bHc

kA/m 573.12 565.16 493.52 398 493.52 493.52 477.6

Oe 30,000 30,200 18,400 10,000 24,700 28,000 29,000 SmCo5

6 mHc

kA/m 2388 2403.92 1464.64 796 1966.12 2228.8 2308.4

Oe 7,400 7,000 5,400 5,400 6,600 6,600 6,000 bHc

kA/m 589.04 557.2 429.84 429.84 525.36 525.36 477.6

Oe 15,000 15,600 8,300 8,800 17,900 16,400 12,000 SmCo5

7 mHc

kA/m 1194 1241.76 660.68 700.48 1424.84 1305.44 955.2

Oe 8,000 7,400 6,200 4,600 7,900 7,100 6,000 bHc

kA/m 636.8 589.04 493.52 366.16 628.84 565.16 477.6

Oe 18,600 19,000 11,400 5,400 19,900 18,000 12,600 SmCo5

8 mHc

kA/m 1480.56 1512.4 907.44 429.84 1584.04 1432.8 1002.96

Oe 6,300 6,400 3,600 250 2,450 7,000 5,000 bHc

kA/m 501.48 509.44 286.56 19.9 195.02 557.2 398

Oe 7,000 7,000 3,900 300 2,400 9,200 5,800 SmCo5

9 mHc

kA/m 557.2 557.2 310.44 23.88 191.04 732.32 461.68

Oe 8,200 8,000 5,800 700 3,100 7,000 6,000 bHc

kA/m 652.72 636.8 461.68 55.72 246.76 557.2 477.6

Oe 18,200 16,000 8,300 600 3,300 14,500 11,600 SmCo5

10 mHc

kA/m 1448.72 1273.6 660.68 47.76 262.68 1154.2 923.36



Values Specimens Coercivity

475 500 650 750 850 900 1000

Oe 9,000 8,400 6,700 4,100 5,700 7,200 5,600 bHc

kA/m 716.4 668.64 533.32 326.36 453.72 573.12 445.76

Oe 19,800 18,400 9,400 4,600 12,800 16,800 11,000 SmCo5

11 mHc

kA/m 1576.08 1464.64 748.24 366.16 1018.88 1337.28 875.6

2.7.6 Discussion

The following discussion is carried through based on above mentioned study: 1. Magnetic performance of SmCo5 permanent magnetic alloy decreases along

with rising of temperature and increases along with lowering of temperature, and has good reversible character.

2. SmCo5 permanent magnetic alloy does not exist low temperature irreversible loss by measuring at low temperature (Zhang, Cheng, Zhang, 1982; Pan, Jin, 1990; Pan, Ma, Li, 1993; Pan, 1996; Pan, Zhao, 1989).

3. It is obtained from magnetic measurement of multi specimens and study that magnetic irreversible loss is related to the height of coercivity of the alloy and some steps out of required criterion of technique in preparation.

4. SmCo5 permanent magnetic alloy has good magnetic performance at low temperature of 1.5K and 40K and is of good application capability (Chen, Li, Wang, Cao, Ma, Pan, 1997; Pan, et al, 1990; Tang, Feng, Luo, Pan, 1994; Tang, Pan, Luo, 1994).

5. It can be seen from Table 2.11 that degradation of coercivity of the SmCo5 permanent magnetic alloy with a higher intrinsic coercivity is smaller than that of the alloy with a lower intrinsic coercivity after annealing at 750 . This indicates that to improve thermal stability of SmCo5 alloy it needs to improve coercivity in preparation technique (Pan, Zhao, Ma, 1989; Geng, Ma, Pan, et al, 1991; Pan, Xiao, 1989).

6. In comparison between bHc and mHc that the degradation magnitude of bHc is smaller than that of mHc.

7. It can also be seen from Table 2.11 that among 11 specimens only specimens No.3 and No.7 of which the coercivity was higher than that at 475 after coer-civity degradation at 750 , and the coercivity of other specimens was lower than that at 475 after coercivity degradation at 750 . This indicates that for most of specimens their coercivity after annealing at 750 can not be restored to the co-ercivity before annealing at 750 .

8. It can be seen from the demagnetization curve at 1.5K and 40K that the ridgy


degree of demagnetization curve of SmCo5 measured at 1.5K is bigger than that measured at 40K.

9. In 1996, K J Strnat developed the first generation of rare earth permanent magnets-SmCo5 magnets of which magnetic energy product is about 5.1MGs�Oe, and be extended to 20MGs�Oe in 1970. In 1996, the second generation of rare earth permanent magnets- Sm2(Co, Fe)17 magnets were developed by K J Strnat. The magnetic energy product of Sm2(Co, Fe)17 magnets were 35MGs�Oe in 1980. NdFeB magnets, the third generation of rare earth permanent magnets, were de-veloped by GE and Sumitomo in 1983, of which magnetic energy product reached 38MGs�Oe (Xu, Xie, Zhang, Huang, Xu, 2008).

2.7.7 Conclusions

It is concluded through above mentioned study that: 1. SmCo5 permanent magnetic alloy has good magnetic performance at 1.5-

523K and has a comparatively high magnetic performance at this temperature range so that this alloy can be used in the instrument with usage range from 1.5 K to 523 K.

2. The lower the temperature is the higher the magnetic performance will be. The magnetic performance at 1.5 K is higher than that at 40 K and that at 40 K is higher than that at 300 K. The higher the temperature is the lower the magnetic performance will be.

3. SmCo5 permanent magnetic alloy has different magnetic performance be-cause of the difference in samarium content. The higher the samarium content is thus the higher the intrinsic coercivity will be and the higher the cobalt content is the higher the saturation magnetization intensity will be, and the lower the coer-civity will be comparatively. The regulation in low temperature is consistent with that in normal temperature.

4. Magnetic irreversible loss is related to the height in coercivity of the SmCo5 alloy. The intrinsic coercivity of the alloy is of a good thermal stability and that its magnetic irreversible loss will be smaller. The height of the minimum the coerciv-ity degraded to at 750 is related to the height of coercivity of the specimen.

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of SmCo5 permanent alloy-ultra-high voltage electron microscopic observation. Pro-ceeding 7th International REPM Workshop: 1983: 291-297

Pan Shuming, Li Zhengwen, Ma Ruzhang (1990) Domain wall and Mössbauer effect of Nd-(Fe-Co-Al)-B magnets with high Curie temperature and high coercivity. Rare Metals, 9(3): 202

Pan Shuming, Liu Jinfang, Luo Helie (1990) The in situ observation of Nd-Fe-B magnets. Journal of Magnetism and Magnetic Materials, 89: 79

Pan Shuming, Ma Ruzhang, Li Zhengwen (1993) Precipitation and growth of colloidal structure in Sm(Co,Cu,Fe,Zr)7.4 permanent alloy and the effect of Zr. Science China (A), 23(3): 317, 544 (in Chinese)

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magnetostrictive properties of TbFe2 compounds. Journal of Applied Physics, 76(10): 7146

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to measure magnetic hysteresis loop of the rare earth-cobalt permanent magnets. Mate-rials of the 5th National Conference of Magnetics, 1983 (internal materials) (in Chinese)

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their application. Metallurgical Industry Press, Beijing: 242, 248 (in Chinese)

Chapter 3 The Second Generation Rare Earth Permanent-magnet Alloys

This chapter discusses the 2nd rare earth permanent magnetic alloy, the 2:17 type rare earth cobalt permanent magnetic alloy (RE2Co17).

Sm(Co, Cu, Fe, Zr)7.4 alloy of high coercivity is a precipitated hardening per-manent magnetic alloy. This alloy has the highest Curie temperature of tC=850 . Its intrinsic coercivity, magnetic induction intensity and theoretical magnetic en-ergy product are higher than those of RECo5. The structure of this type of alloy: in the center is Sm2(Co, Cu, Fe, Zr)17 of about 50nm. The boundary phase is com-prised of Sm(Co, Cu, Fe, Zr)5 with a thickness about 10nm surrounding the main phase. This is called a cystiform structure. The coercivity of the 2:17 type alloy does not rely on the size of crystal grain, rather on the microstructure between two phases. These two phases are the 2:17 phase and the 1:5 phase. The domain of 2:17 phase is pinned by the 1:5 phase when magnetization and anti-magnetization, due to the difference in domain between the 2:17 phase and 1:5 phase. Thus, the alloy possesses very high coercivity. The key for obtaining high coercivity depends on the heat treatment technique, which generally adopts solid solution treatment, isothermal ageing treatment, etc.

This chapter expatiates in emphasis on observation and research of high coer-civity Sm(Co, Cu, Fe, Zr)7.4 alloy using JEM-1000 HVEM, especially the in situ observation on precipitation, growth and growth velocity of cystiform structure at continual rising temperature from room temperature to 950 . The analysis and deep research were carried out on the driving force and resistance of phase trans-formation process, phase transformation dynamics, coercivity mechanism of cys-tiform structure broken under long term high temperature aging, diffusion type continual phase transformation and new opinion on enhancing metastable transi-tion phase. Research was made on zirconium in Sm(Co, Cu, Fe, Zr)7.4 alloy to enhance coercivity, heighten solubility of Sm in the alloy, improve square degree of demagnetization curve through nuclear physics method. Iron atom was pro-moted to enter into Co3 site from Co1 site in 2:17 phase and thus to raise uniaxial


anisotropy of alloy. These two points benefit the enhancement of the coercivity of this alloy.

This chapter also researches the magnetism of high coercivity Sm(Co, Cu, Fe, Zr)7.4 permanent magnetic alloy from 1.5K to 523K, measured the magnetism of the alloy after spinning and re-orientation at 1.5K and magnetism at 1.5K with changing temperature.

3.1 Phase Precipitation, Phase Transformation at High Tem-perature and Magnetism of High Coercivity Sm(Co, Cu, Fe, Zr)7.4

The Curie temperature and saturation magnetization intensity of Sm2Co17 are both higher than RECo5 among the rare earth Co-intermetallic compounds, and it has been widely used in industry, agriculture, and national defense and transportation areas for more than twenty years. Among the Sm2Co17 type permanent magnetic alloys the most excellent in magnetic performance is Sm-Co-Cu-Fe-M series, where M are Zr, Hf, Ti, Ni, Mn, etc. However, Sm(Co, Cu, Fe, Zr)z (z=7.0-8.5) is the best one in magnetic performance among all the Sm2Co17 permanent magnetic alloys. Sm2Co17 alloys can be divided into three kinds in accordance with their zirconium content. The first kind has a higher magnetic-energy product and low coercivity, and its magnetic performance is: (BH)max=175-251kJ/m3. It has a bet-ter quadrate degree in its demagnetization curve. The second kind has super high coercivity, and its coercivity mHc = 1990-2388 kA/m, (BH)max = 190-223 kJ/m3. The third kind is in between, which is called Sm2Co17 permanent-magnet alloy with high coercivity that mHc = 796-1259 kA/m, bHc = 636-716kA/m, (BH)max = 199-223kJ/m3. Its Curie temperature is 840-870 , magnetic induction tempera-ture coefficient is 0.02%/ . It can be used at the temperature among 60-350 , and is lower in the quantity of samarium and cobalt, so that it may be a potentially ideal permanent magnet alloy, of which the theoretical magnetic-energy product may reach 525.4 kJ/m3 (66 MGs·Oe). From the magnetic performance mentioned above, its actual magnetic-energy product is only a half of the theoretical value, thus it has to sacrifice some coercivity in order to enhance magnetic-energy prod-uct. How to enhance coercivity further at the same time to get high magnetic-energy product? For this purpose it is necessary to study on coercivity theoreti-cally (Zhou, 1990).

Sm(Co, Cu, Fe, Zr)7.4 with high coercivity is a kind of precipitation-hardening permanent magnet alloy with cellular microstructure. The coercivity of the alloy is determined by pinning from cellular structure to domain wall. What is the pin-ning center of domain wall Are there some relationships between coercivity and the size and shape of the cellular structure, incomplete of appearance and growth of cell, and coherency stress state of matrix phase and precipitated phase?

There are many approaches for studying coercivity. One is to study specimens


with different coercivity using nuclear physical method and analyze the result with combining magnetism loop getting from magnetic measurement; another is to build a physical model to research relationships between coercivity and the physical and geometrical parameters of the precipitated phase after mathematical treatment. The third is to observe at first, the microtexture of the alloy using HVEM and to measure magnetism parameters, and research the relationship be-tween coercivity and geometrical parameter of microtexture and physical parame-ters of new phase after measuring coercivity.

Coercivity is a physical quantity being sensitive to microtexture. The relation-ship and interaction of microtexture and domain structure are the foundations for researching coercivity of permanent-magnet alloy. The scholars engaged with this study and scientists working on magnetics and metal physics have made a lot of excellent achievements. But the microtexture obtained by above experiments was observed at room temperature so that it lacks the integral and systemic under-standing on the formation of phase (Zhou, 1990; Pan, Ma, Li, 1993; Pan, Zhao, 1988; Xiao, Zhang, Zhou, et al, 1982; Rothworf, et al, 1982; Fidler, et al, 1982; Yu, 2003).

This study adopted a 1000kV HVEM for dynamic observation on alloy speci-men at thermal state for a more accurate understanding on formation of various microtexture and associated with variation of coercivity in different states so that to judge the influence of microtexture of the alloy on coercivity further for pur-pose of studying coercivity and thus enhancing the coercivity.

Recent research work on (Sm1 xGdx)(Co, Cu, Fe, Zr)z indicates the temperature coefficient reaches zero when x=0.55. And z value affects the temperature coeffi-cient as well. When z=7.87 the alloy has the lowest point at z value vs the tem-perature coefficient curve (Pan, 2011).

3.1.1 Specimen preparation process and experimental method

Powder metallurgy was used to prepare Sm(Co, Cu, Fe, Zr)7.4 in the experiment. The chemical composition (mass fraction) was: 26%Sm, 51.4%Co, 15.8%Fe, 2.4%Zr and 4.4%Cu. The alloy with above composition was melted in vacuum induction furnace and molding in water-cooling copper mould. After the ingot obtained was crushed coarsely and pulverized into powder of about 5m as the average granularity under the protection of toluene. The powder was oriented in magnetic field of 1.5T and formed in press machine. The formed roughcast was sintered at 1210 and solid dissolved at 1155 and then wind cooled to room temperature for use. The aging temperature is 850 and 400 . Then the speci-men was ready for magnetism measurement using CL6-1 magnetic parameter instrument.

The preparation method of specimen for observation by electronic microscopic was as follows. First, solid resolving the alloy for 0.5h and quenching to room


temperature. Then the specimen was sliced into lamellae with a thickness of 0.25mm at the direction perpendicular to c axis. At last, the filmy specimen was thinned and rinsed at ion-thinning machine.

Both static and dynamic observation of filmy specimen was proceed in JEM-1000 HVEM with operation voltage 1000 kV in vacuum of 2.5�10�4Pa and 0.7 �10�5Pa after adding liquid nitrogen. Ionic beam was 10A. The specimen was inserted into the side inserted heating experiment dais of TEM-1000 and was heated-up from room temperature to high temperature, held and videotaped in electronic microscope in dynamic observation. The change in microtexture of permanent-magnet alloy was observed directly in electronic microscope and rele-vant electronic diffraction diagram was made.

3.1.2 Results of specimen magnetic measurement

The specimen was measured magnetically using magnetic parameter measure-ment apparatus. The results of magnetic measurement was as follows: the mag-netic remanence induction intensity Br = 1.12T, intrinsic coercivity iHc = 1078 kA/m, and maximum magnetic energy product (BH)max = 243.6kJ/m3. The de-magnetization curve of the specimen is shown in Fig. 3.1, and the curve at rising temperature is shown in Fig. 3.2.

Fig.3.1 Demagnetization curves of specimen of Sm(Co,Cu,Fe,Zr)7.4 alloy

(alloy composition: 26%Sm, 51.4%Co, 15.8%Fe, 2.4%Zr, 4.4%Cu)

Fig.3.2 Heating curve of the film surface of Sm(Co, Cu, Fe, Zr)7.4 alloy

(alloy composition shown in Fig. 3.1) perpendicular to c axis


3.1.3 Microtexture of the alloy at room temperature

After solid solution treatment and succedent aging treatment at below 850 , mi-crotexture of the alloy appeared in exiguous cellular at the state of high coercivity. The cell wall was 1:5 type phase, the inner-cell was 2:17 type phase, and the two type phases are coherent (Fig. 3.3). The diameter of cell is about 50nm and the thickness of cell wall was 5nm (Rothworf, 1982). The plane was vertical to c axis thus the cells was equal-axial. Fig. 3.4 shows the electronic diffraction pattern. Fig. 3.5 shows the microtexture at room temperature. Fig. 3.6 shows electron mi-crograph at room temperature. Fig. 3.7 shows the relationship between diameter of cell and coercivity of 2:17 type alloy.

Fig.3.3 Cellular microstructure of Sm(Co,Cu,Fe,Zr)7.4 alloy at room temperature

(The thickness of cell wall is 5nm. Inside cell is Sm2Co17 phase which size is 50nm. The cell wall is SmCo5 phase)

Fig.3.4 Electron diffraction pattern of Sm(Co,Cu,Fe,Zr)7.4 alloy at room temperature


Fig.3.5 Microstructure of Sm(Co, Cu, Fe, Zr)7.4 alloy at room temperature (It can be seen that a lot of crystal cells are severed by bottom crystal plane, and martensite plate is

also observed: along the habit plate in matrix phase)

Fig.3.6 The electron micrograph of Sm(Co, Cu, Fe, Zr)7.4 alloy at room temperature

(a) The dislocation circle in a crystal grain of quenched alloy;

(b) The dislocation nets in a crystal grain

Fig.3.7 The relation of coercivity of precipitation hardening 2:17 type alloy

vs the size of the crystal cell


3.1.4 The in situ and dynamic observation on precipitation, growth up and high temperature phase transformation of cellular structure from room temperature to high tempera-ture

The dynamic observation result of this experiment showed that cellular texture was extraordinary stable from 700 to 840 . Integrated cellular structure was observed again after heat preservation for 80min at 900 , which illustrates the reversibility of the phase transformation.

3.1.4.1 The positioning situ and dynamic observation process on the phase transformation

There were some discussions about the relationship of cellular texture and coer-civity being heard before our experiment study, such as the cellular texture began to be formed at 700 (Fidler, Skalicky, 1982; Sun, et al, 1983) and was de-stroyed at constant temperature aging at 800 (Cai, Rong, 2003). Some changes might happen to inner-cell and cell wall at 830 (Xu, 2000; Kronmüller, 1982), cellular texture was formed after aging at 830 for 30min and cell wall varied at 830 . Thus some big cell would swallow the small one. And that the influence of long term aging on coercivity was attributed to the appearance of zirconium enriched long lamellar phase (Table 3.1). These conclusions, surmise and deduc-ing were primarily deduced from observations at room temperature and in ad-vance designed big specimen and aging condition as per the deduction. And thus these were difficult to avoid being unilateral. To overcome this point we adopted filmy specimen being heated from room temperature to 950 serially in trans-mission electronic microscope so that to observe directly the variation of micro-texture at every temperature section (different aging condition). We observed ap-pearance, upgrowth and formation of cellular texture; and appearance, change, etc. of zirconium enriched long lamellar phase. This section concludes above problem based on experiments as follows.

Table 3.1 Relationship between demagnetization temperature and intrinsic coercivity

Annealing temperature/ 300 400 500 600 700 800 900

Intrinsic coercivity, mHc/kA·m�1 45 46 49 160 328 368 356 1. It is found that cellular texture of permanent magnet alloy appeared at 460

and its shape was spherical (see Fig. 3.8). The coercivity of the alloy was very sensitive to diameter and size of cellular texture. Their relationship is shown in Fig. 3.7.

2. Cellular texture started to form at 460 (Fig. 3.9), and grew and developed at 500-600 ; the cellular texture size was different at 500-600 , some were


grew up but some did not form integral cellular texture and some were only in embryo. Integrated cellular texture appeared at 700 , they came to perfect at 790 (Fig. 3.10 to Fig. 3.13).

Fig.3.8 Electron micrograph of the cellular structure of Sm(Co, Cu, Fe, Zr)7.4 heating to 460

(The black point is initial phase point)

Fig.3.9 Electron micrograph of the cellular structure of Sm(Co, Cu, Fe, Zr)7.4

heating to 500

Fig.3.10 Electron micrograph of the cellular structure of Sm(Co, Cu, Fe, Zr)7.4 heating to 700


Fig.3.11 Electron micrograph of strip structure of Sm(Co, Cu, Fe, Zr)7.4 perpendicular to

c axis heating to 780



Fig.3.13 Electron micrograph of strip structure of Sm(Co, Cu, Fe, Zr)7.4 perpendicular to c axis heating to 790


3. Lamellar phase, which is perpendicular to c axis, appeared at 780 , but only partly in visual field. At 785 , it appeared in most of visual field. And at 790 , the lamellar phase being perpendicular to c axis appeared at the foundation of cellular structure (somebody called this phase as Z phase and somebody called as zirconium enriched long lamellar phase).

4. Cellular texture and lamellar structure coexists at 800 (seeing Fig. 3.13 and Fig. 3.14).



5. Lamellar phase Z phase aforementioned upgrowth at 820 . Lamellar phase increases gradually in company with continually rising of aging

temperature. Fig. 3.14 is the lamellar phase growing on original situs at 810 . Fig. 3.15 is the lamellar phase growing on original situs at 820 . Fig. 3.16 is the lamellar phase growing on original situs at 830 . The lamellar phase at 840 and holding for 0.5h. It can be seen from Fig. 3.17 that is the lamellar phase run through all the vision.

Fig.3.15 Electron micrograph of strip structure at 820


Fig.3.16 Electron micrograph of strip structure at 830

Fig.3.17 Electron micrograph of strip structure of Sm(Co,Cu,Fe,Zr)7.4 heating at 840 for 40min

(Strip and banding structures parallel with c axis appear)

Some striation material and strip material being parallel to c axis appeared at 840 and holding for 40min, the composition of these materials are still un-known.

3.1.4.2 Analysis and discuss on magnetic performance and heat treating system

Sm(Co, Cu, Fe, Zr)7.4 permanent magnetic alloy is a kind of precipitation harden-ing permanent magnetic alloy and the production of coercivity comes from the pinning relative to domain wall by Sm(Co, Cu)5 enriched phase precipitated in aging process (Pan, Ma, Li, 1993). Pinning intensity of Sm(Co, Cu)5 on domain wall depends on its shape, amount and composition difference from the base phase. Although Sm(Co, Cu)5 phase precipitated generally from the base phase at 460 but this time these cellular wall phase did not fully connect together and that the copper in the base phase still could not defused fully to cellular wall be-cause of the comparative low temperature, i.e., in this time the different between domain wall and inner cell is not big enough. Thus no matter the embryo shape of the cellular texture already formed at this temperature but did not acted obvious hardening function on the coercivity, which is consistent with the result shown in Fig. 3.18 (Pan, Zhao, 1988).


Fig.3.18 The relation of peak coercivity of Sm(Co, Cu, Fe, Zr)7.4 vs annealing

temperature

When the temperature was over 700 , the cellular texture grew up adequately and the cellular connected together thus the cupper content in inner cell and cell wall so the difference between domain wall energy of inner cell �domain and do-main wall energy of cellular wall was big enough. Therefore mHc increased rap-idly as per mHc = (�domain �wall) / (2Ms�) (Fig. 3.7) (Pan, Ma, Li, 1993).

Alloys with different types coercivity could be obtained by changing heat treatment condition of Sm(Co, Cu, Fe, Zr)7.4 alloy, and its demagnetization curve is shown in curve and of Fig. 3.19 (Xiao, Zhang, Zhou, et al, 1982). Their respective heat treatment system is shown in Fig. 3.20.

Fig.3.19 Demagnetization curves of Sm2(Co, Cu, Fe, Zr)17 and high coercivity Sm2(Co,

Cu, Fe, Zr)17 after different heat treatment


Fig.3.20 Curve of 2:17 type Sm-Co alloy corresponding to annealing process (a) and

curve corresponding to annealing process (b)

3.1.4.3 The formation, growth and growth velocity of cellular texture

Precipitation appeared in specimen at 300-400 in dynamic observation and this decomposed precipitate was coherent with the matrix phase. These tiny spheric precipitates with a size of several nanometers were observed in transmission elec-tronic microscope and these decomposed material made the system with very high interface energy. In this time, overall system was in unstable state under heat acti-vation. To reduce the total interface energy small particles in the decomposed grain dissolved, the large ones grown up and continually dissolved other compo-nents. Fig. 3.21 shows the longitudinal section figure of Sm-Co-13Cu-10Fe (Cu > Fe) in which samarium accounted for 10%-16% (at.). There are four phases in the figure: L is liquid phase, Co phase is cobalt based solid solution, 2: 17 phase (three structures of Th2Ni17, Th2Zn17, and TbCu7), and 1:5 phase. At 1210 the solubility of 2:17 phase relative to samarium reaches up to 13% from 10.5% (both in atomic fraction). In this composition its melting point is lower along with in-crease of samarium content. The alloy with samarium of 11.9%-12.5% (atomic fraction) is of TbCu type structure at 1210 , and the homogenizing treatment in TbCu phase region can obtain a comparative higher coercivity. Structure of 2:17 phase is related to content of iron and cupper contents that the alloy with 2.0% zirconium and 11.5% samarium has TbCu7 type structure at high temperature. The mHc = 2000 kA/m may be reached after solid solution treatment and aging (Zhou, 1990). Generally speaking, there are only two constituents and solubility of one constituent in the other constituent is limited and the solubility will become smaller while the temperature becomes lower, as is shown in Fig. 3.22. One type of constituent dissolving into an � solid solution and then quenching to suppress the precipitation of this constituent to reach saturated solid solution is called as solid solution treatment. The reverse process of solid solution treatment is decomposition. The hardness and intensity of the alloy would be improved in decomposition proc-ess, namely aging hardening. At the anaphase of decomposition equilibrium phase appears by decomposition. During the dynamic observation of 300-600 , bulky grain grown up and coarsened after dissolving with small particles, so concentration gradient exists consequentially between bulky grain and small particles. In


Fig.3.21 The phase diagram of Sm-Co-13Cu-10Fe (Cu>Fe) in lengthwise section

Fig.3.22 The diagram of the solubility of � in � which decreases when the

temperature goes down (Two constituent elements of the alloy in Fig. 3.21)


the dynamic observation it takes 6min from beginning to 460 , and takes 12min from 460 to 600 . It was found in dynamic observed at 600 some cellular textures began to grow up but their size was very different and most of them were in embryo. There was no integrated cellular texture. Temperature was raised in the electronic microscope and taken 6min from 600 to 700 . At this temperature the integrated cellular texture was observed (Fig. 3.10). Fig. 3.23 is the sketch map of formation process of cellular texture and Table 3.2 shows the growth time.

Fig.3.23 The diagram showing the forming process of the cellular structure of 2:17 type

permanent magnet alloy

Table 3.2 List of time needed for precipitation and growth up (coarsening) of cell texture of Sm(Co, Cu, Fe, Zr)7.4

Growth process

From room temperature

to 460

From 460 to 600 , cellular

structure appears and grows at 460

From 600 to 700 , integrated cellular structure was observed at

700

From 700 to 790 , zirconium

enriched long lamellar phase

appears at 790

From 790 to 840 , strip

structure appears at 840

Time/min 6 12 6 6 30

According to diffusion control grown up theory, can we figure out the growth velocity of cellular texture from precipitation to integral one? It was observed that the cellular texture was orbicular, i.e., a sphere shape. Supposing radius is r1, sol-ute quantity absorbed in unit time in growth process of cellular texture is 2

14 r (Cb Cr)v, where Cb and Cr are the interfacial volume concentration of precipitated phase, v is growth rate. Diffusion flux of particle is 4r 2D(dC/dr). The following formulae are come into existence.

2 21 b r

2 21 b r

d4 ( ) 4d

d d( )

Cr C C r Dr

D C rr C C r

� �

��

v

v

Supposing the rC is the concentration of the base phase being equivalent to average granule radius r of granules system corresponding to solubility (Yu, 2003). By integrating r from r1 to �, C from C = rC to C = rC , it can be derived that:


r r b r( ) ( )D C C C C r� � �v So growth velocity v is

r r

b r

C C DC C r

��

�v

3.1.4.4 Analyzing on driving force and resistance of cellular texture and phase transformation

During the phase transformation process of the alloy, no mater nucleation or not, phase transformation always encounters driving force and resistance. Thus for phase transformation energy proceeding spontaneously a necessary premise is that the free enthalpy of new phase less free enthalpy of old phase is minus. If Q represents for the resistance of phase transformation, then Q = �G1+�G2, �G1 represents for interfacial energy and �G2 represents for distortion energy. Interfa-cial energy is dividend into chemical interfacial energy and structural interface energy. There are four kinds of property of interface, and that is ideal coherence, complete coherence, partial coherence and non coherence. Interfacial energy of complete coherence is very small, and mismatch degree of interface atom is zero. The mismatch degree of non coherence is the maximal. The degree mismatch of complete coherence is: � = �a / a < 0.05; mismatch degree of partial coherence is: � = �a/a = 0.05-0.25 (Liu, Ren, Song, 2003). Table 3.3 shows the lattice constant and lattice mismatch caused by precipitation of 1:5 phase in the 2:17 base phase.

Table 3.3 Variation of lattice constant and lattice mismatch in the matrix of 2:17 phase

caused by precipitation of 1:5 phase (Xiao, Zhang, Zhou, et al, 1982)

Alloy phase Structure (space group) a/nm c/nm �a/% �c/%

Sm2Co17 R3m 0.8406 1.2230 0 0

SmCo5 P6/mmm 0.4997 0.3975 0.030 0.030

SmCo4Cu1 P6/mmm 0.5013 0.4001 0.033 0.019

SmCo3Cu2 P6/mmm 0.5029 0.4025 0.036 0.013

SmCo2Cu3 P6/mmm 0.5045 0.4049 0.040 0.007

SmCu5 P6/mmm 0.5074 0.4099 0.046 0.005

The distortion energy is divided into coherence distortion energy and non-

coherence distortion energy. Coherence distortion energy: shear elasticity coher-ence distortion energy is related to the rotation angle � of atomic plane, matrix phase shear modulus G and shape of new phase in phase transformation. Suppos-ing distortion occurs in matrix phase, volume distortion energy is approximate as (3/2)E�2, where E is elastic modulus and � is mismatch degree that � = �a/a. Anyway shear elasticity distortion energy and shape factor are closely related to each other (Liu, Ren, Song, 2003). Fidler J. thought that the coercivity of precipi-


tation hardened the 2:17 type alloy was determined not only by domain wall en-ergy gradient ��, but also by magnetoelastic coupling energy between domain wall magnetostrictive stress and precipitated phase strain, that is to say, the con-tribution lattice mismatch degree of 1:5 phase and 2:17 phase to the coercivity is not neglected. That can be formulated: mHc � |�11�a/(3Ms)| (Zhou, 1990; Fidler, 1982). It was found in the in situ and dynamic observation of experimental specimen that the Sm2(Co, Cu, Fe, M)17 main phase, which took 50min from pre-cipitated cellular embryo coarsening and upgrowth to integrated organization structure, was surrounded by Sm(Co, Cu, M)5 border phase of 10nm in thickness only took 6 min. Table 3.2 and Fig. 3.10 testified the precipitation of above men-tioned Sm(Co, Cu, M)5 phase, and the transformation did not need very big nu-cleation power and incubation period (that the critical crystal nucleus must jump over the energy barrier is the nucleation power) so that it is easy to complete from embryo to integral cell structure of 1:5 phase. When diffusion driving free energy �G<0, or G�= (d 2G /dC 2 ) <0, the very small fluctuation will be formed and this small fluctuation with certain wavelength provides a condition for new phase up-hill diffusing that (d 2G /dC 2 )<0 (Xu, 2000). It is also found from the dynamic observation in this experiment that cellular texture was extraordinary stable from 700 to 840 , seeing Fig. 3.28. From the zirconium enriched lamella phase run-through at 790 to completely drilling through the 1:5 phase it was still stable, this is testified experimentally that the precipitation of 1:5 phase needed to along pyramidal plane of which lattice misfit is zero and a long rhombic cell was formed (Fig. 3.13 and Fig. 3.28). It is also found from dynamic observation of the experimental specimen that the strip, nubby and anomalous geometrical shape appeared at 840 and holding for 40-80min (Fig. 3.24 to Fig. 3.28). As shown in Fig. 3.29, integrated cellular microstructure reappears when cooling back to room temperature from 80min heat preservation at 900 , which illustrates the reversi-bility of the transformation.

Fig. 3.24 Electron micrograph of Sm(Co,Cu,Fe,Zr)7.4 heating at 840 for 40min

(The observation plane is perpendicular to c axis. The cellular structure of 2:17 type alloy becomes un-equiaxial, irregular shape)


Fig. 3.25 Electron micrograph of Sm(Co,Cu,Fe,Zr)7.4 heating at 840 for 50min

(The observation plane is perpendicular to c axis. The cellular structure of 2:17 type alloy becomes un-equiaxial, irregular shape)

Fig. 3.26 Electron micrograph of Sm(Co,Cu,Fe,Zr)7.4 heating at 840 for 60min (The observation plane is perpendicular to c axis. The cellular structure of 2:17 type alloy disappears

and rhabditiform structure appears)

Fig.3.27 Electron micrograph of Sm(Co, Cu, Fe, Zr)7.4 heating at 840 for 65min

(The cellular structure in the observation plane perpendicular to c axis is destroyed)


Fig. 3.28 Electron micrograph of Sm(Co, Cu, Fe, Zr)7.4 heating at 840 for 10min (The cellular structure is observed and complete Z phase appears in bottom crystal plane cutting a lot

of crystal cells and cell walls)

Fig. 3.29 Electron micrograph of Sm(Co,Cu,Fe,Zr)7.4 alloy heating at 840 for 80min then cooling down to room temperature

(The clear cellular structure with different size is obviously observed)

3.1.4.5 Kinetics analysis on phase transformation

In the following to focus on discussing the phase transformation kinetics analysis on the phenomenon in dynamic observation of Sm(Co, Cu, Fe, Zr)7.4 specimen. In viewpoint of phase transformation kinetics: in homogenous single-phase or sev-eral mixed phases appeared variation of phase with different composition or posi-tion of different ions, atoms and electron, and exposure (different structure), char-acter, different configuration, being so called as phase transformation. Thus that the phenomenon of specimen appeared after holding for 1h at 840 belongs to


phase transformation category. During phase transformation, the matrix phase became unstable and decomposed into stable two phases and their free energy was lower than that of the matrix phase. When different concentration fluctuation appeared in the matrix phase the free energy and nucleation power (nucleation driving force) all could be derived from the intersection of tangent and free en-ergy curve (Xu, 2000). When concentration fluctuation appeared and the concen-tration exceeding a given level, i.e., exceeding the critical value of free energy change will actually become the energy barrier which one crystalline nucleus has to exceed that an embryo is smaller than the critical size it will be dissolved back to the matrix phase spontaneously and the size of one embryo is bigger than the critical size it will up growth to be a crystalline nucleus and continually grow up to new phase. Nucleation field could be deduced by specimen in this experiment and its value can be estimated as follow:

Theoretical minimal value of nucleation field Hn can be calculated by following formula:

2 2s n

3 32 (3 3 6 )2

M H a h a ah -� �

Nucleation field can be derived from formula above:

ns

1 1Hh a M

-� ��

where Ms is the saturation magnetization intensity; � is the energy density of do-main wall; a is the length of hexagonal prism one side; h is the height of the hex-agonal prism.

Supposing the reversal magnetization nuclear radius is R and when R = 10nm, it can be calculated that Hn is 2,388 kA/m and similarly when R = 100nm, Hn is 238 kA/m.

Defects appeared in cellular texture after holding about one hour at 840 , that is to say, a, h and R increase in the above calculation and that nucleation field Hn decreases, so that coercivity will degrade remarkably. It was observed from fore photographs from Fig. 3.24 to Fig. 3.27 that defects appeared in new precipitated and thus holding for long time at 840-860 caused degradation of coercivity. That these very low anisotropy areas existed in a considerable volume in new precipitated phase caused degradation of coercivity, as shown in the curve in Fig. 3.18.

3.1.4.6 Coercivity mechanism after destroying of cellular texture by long time aging at high temperature

When long time aging at above 840 and the cellular texture began to be de-stroyed the Sm(Co, Cu, Fe, Zr)7.4 alloy and its coercivity was controlled by nu-


cleation and local pinning. The so-called coercivity controlled by nucleation field is determined by the magnetic field which is necessary for building reversal mag-netization nucleus, such as inhomogeneous of grain boundary, irregular surface, oxidation and other form defects. It is preferential to become reversal magnetiza-tion nucleation center for these regions with low magnetocrystalline anisotropy, and thus to enhance coercivity it is necessary to avoid these defects. If pinning position is just in near nucleation site and to limit extending of the reversal mag-netization nucleus the coercivity will be enhanced. It was found through dynamic observation in this experiment that there were big and integral cell structure and local high density dislocation network, which promote the appearance of decom-position cupper granules so that to become the pinning position. These are suffi-cient condition and key factor for alloy specimen to have a higher coercivity. For this kind of alloy to enhance coercivity should select aging temperature below 840 , for example, 800-830 would be a good choice.

For the alloy of Sm(Co, Cu, Fe, Zr)7.4 its coercivity is determined by pinning field. The pinning center is cellular texture of 2:17 phase 80% surrounded by 1:5 phase 20% with a thickness of 10nm. The 1:5 phase acts as pinning function on domain wall. The coercivity of cellular texture is:

m cs

32

rHM d �

��

where Ms is the saturation magnetization intensity; � is the thickness of magnetic domain wall; d is cell diameter; �r is energy difference of domain wall:

1:5 2:17 1 1 2 2 1 1 2 22 2 2 ( )r r A K A K A K A K� � � � �

Energy difference of domain wall is mainly decided by magnetic anisotropy constant K1 of 1:5 phase and 2:17 phase and exchange integral A. The A and K1 are mainly determined by composition and defects.

It was seen in the in situ and dynamic observation by electronic microscope that cellular texture has formed at 700 , zirconium enriched long lamella phase appeared at 790 . In order to get this kind of alloy, the first grade aging should select 790-830 to set temperature range and to find out the maximal peak of coercivity by experiment because the alloy of 1:5 phase is stable at 790-830 .

3.1.4.7 Diffusion type continuous phase transformation

Above phase transformation of Sm(Co, Cu, Fe, Zr)7.4 permanent magnetic alloy belongs to diffusion continuous phase transformation. Under continuous and or-dered and fluctuation congregating and through uphill diffusion and congregation a single phase is disaggregated into Sm2(Co, Cu, Fe, M)17 phase with a metastable size about 50nm and Sm (Co, Cu, Fe, M)5 phase with a size about 10nm. The former is surrounded by the latter and these two phases are coherent. Any small


fluctuation would make concentration amplitude heightened continually and to form new phase. The fluctuant concentration continuously changes with time, at beginning stage the interface between above mentioned 2:17 phase and 1:5 phases is not obvious; after disaggregating the size distribution of the 1:5 phase and the 2:17 phase becomes comparatively regular with highly continuity (Xu, 2000; Xiao, 2004).

It was observed from the specimen of this experiment that the nucleus of em-bryo was diffusion nucleation: the course is to form new phase core at first, namely giving birth of nucleus. The fluctuation of nuclei embryos or new phase points rely on diffuse transition of individual atom by thermal activation and gra-dient of free energy as driving force of phase transformation. Solid state phase transformation belongs to inhomogeneous nucleation.

3.1.4.8 Metastable transitional phase

After the metastable transitional phase being aging for 20h at 830 its cellular texture become non-integrated any more that part of them have been destroyed and a few of them become trip structure, this trip structure might belong to a tran-sitional phase. When temperature was raised to 1050 or lower to 820-800 the transitional phase transforms into stable phase of 1:5 phase and 2:17 phase with cellular texture.

The experiment indicated: in order to get rapidly increase of mHc aging time should not be less than 2h at 840 .

Many researchers considered that the influence of long time aging on coercivity was attributed to appearance of zirconium enriched long lamellar phase (some-times called as Z phase) being perpendicular to c axis. The experiment in this sec-tion indicates that this long lamellar zirconium enriched phase began to form at temperature heated to 780 . And when holding for 0.5h at 840 this long lamel-lar phase has run through the whole visual field, seeing Fig. 3.17 and Fig. 3.18. This fact does not fit the result that coercivity enhanced greatly with holding at equal temperature of 850 for 2h. Therefore, author believes that appearance of Z phase does not have direct contribution to high coercivity and is most possible to have indirect function, such as accelerating component diffusion as a passage or resulting in formation other new phase (Zhou, 1990; Pan, Ma, Li, 1993; Pan, Zhao, 1988). The function of zirconium sees the section 3.2 of this book.

3.1.5 Conclusions

1. The coercivity of Sm(Co, Cu, Fe, Zr)7.4 permanent magnetic alloy increases with quantity increases and upgrowth of cellular texture, and decreases with de-struction and subtraction of cellular texture. Coercivity of the alloy is very sensi-tive to the pattern, size, upgrowth degree and quantity of cellular texture.


2. From room temperature to 800 the process of temperature rising is a heat activation course. The heat activation is motivity for appearance, upgrowth and formation of cell texture. It controls atomic diffusion in components of the alloy. Owing to this heat activation function cupper in alloy diffuse to cell wall gradu-ally. When temperature is raised to above 700 , the cellular texture would grow up adequately and cell wall would be interconnected, cupper content in cell wall and interior of cell becomes big enough thereby it would result in that the domain wall energy difference between inter-cell and cell wall domain becomes sufficient big. Because coercivity is proportional to the difference of them thus coerciv-ity increases the Sm(Co, Cu, Fe, Zr)7.4 permanent magnetic alloy could obtain various types of alloys with different coercivity through different heat treat condi-tion.

3. Cellular texture of Sm(Co, Cu, Fe, Zr)7.4 permanent magnetic alloy appeared at 460 and its shape was spherical. It began to form at 500 , grown and devel-oped at 500-700 , integrated structure formed at 700 , and grown up perfectly at 790 . The cellular texture would change when holding at 840 for 40min.

4. It was found in the in situ and dynamic observation of Sm(Co, Cu, Fe, Zr)7.4 permanent magnetic alloy by transmission electronic microscope that the zirco-nium enriched lamellar phase (or called Z phase) being perpendicular to c axis appeared at 780 and ran through the whole visual field and cellular texture at 830 , but this lamellar phase had no direct contribution to coercivity of the alloy.

5. The precipitation and growth of the cellular texture of Sm(Co, Cu, Fe, Zr)7.4 permanent magnetic alloy are a complete process for cellular texture and phase transformation belonging to diffusion continuous phase transformation.

6. From precipitation of the cellular structure of the Sm(Co, Cu, Fe, Zr)7.4 per-manent magnetic alloy to growth up perfectly at 840 the coercivity of the alloy is determined by pinning field. After the cellular texture was destroyed coercivity of the alloy is determined by nucleation field.

7. Coercivity of Sm(Co, Cu, Fe, Zr)7.4 permanent magnets increases with the increase in the number, development, growth of cellular structure, decreases with the destruction and reduction of the cellular structure. The coercivity of alloy is sensitive to the morphology, structure, size, development process and the number of the cellular structure. Alloys with different coercivity can be obtained through different heat treatment conditions. From room temperature to 800 , the process is a thermally activated process. This process is the driving force of the emer-gence and development of cellular structure, which dominates the diffusion of the atoms in the alloy composition.

8. For (Sm1-xGdx)(Co, Fe, Cu, Zr)z magnets, with the increase of z magnet re-manence temperature coefficient firstly increased then reduced under the same temperature. When z =7.87, remanence temperature coefficient reaches a mini-mum, the most stable magnet has the largest coercivity (Pan, 2011).


3.2 Function of Zirconium on Sm(Co, Cu, Fe, Zr)7.4 Permanent Magnetic Alloy & Observation and Analysis by Electron Microscope *

This section is about directly observing precipitation and development of zirco-nium enriched long lamellar phase (or called Z phase) by 1000kV HVEM and combining with magnetism measurement to research coercivity mechanism of this alloy; and making use of Mössbauer effect to study impact of zirconium on the occupation status of each atoms crystal site and its relationship with coercivity.


Specimen preparation for electronic microscope: Specimen components in ex-periment: 26.01%Sm, 51.38%Co, 4.36%Cu, 15.19%Fe, 2.99%Zr; solid solution treatment for 0.5h at 1180 and then quenching to room temperature. The specimen was sliced to lamellae of the thickness of 0.25mm along the direction parallel to c axis. Then was electrolysis thinned by 20% perchloric acid and 80% glacial acetic acid to almost perforation. Afterwards the specimen was thinned and cleaned at last on ion thinning device.

Experimental method: The observation of filmy specimen was carried out using JEM-1000 HVEM with operating voltage 1000kV, output voltage 333V, current 6.2A, vacuum 3.32�10�4Pa and 9.3�10�5Pa after adding liquid nitrogen, and ion beam 10A. In experiment at first put specimen into JEM-1000 side insertion heating experimental dais, and then carried out heating from room temperature to 850 and held under electronic microscope.

Mossbauer experiment was carried out on constant acceleration Mossbauer Spectrometer, the specimen was powder. The experiment was proceeded at room temperature, radioactive source is 57Co(Rh) and the intensity is about Zomci, took count of was above 2.5�106, velocity demarcate adopted c-Fe standard spectra.

3.2.2 Research on function of Zirconium

The in situ and dynamic observation using electronic microscope:At room tem-perature alloys appeared as single-phase status without precipitated phase in crys-tal granule except crystal boundaries and a small quantity of defects. When tem-perature rose to 460 the cellular texture appeared on base phase. When tem-perature rose to 780 the lamellar phase being perpendicular to c axis appeared on base of original cellular texture (Fig. 3.11). It can be seen from Fig. 3.11 that

Cooperators of this study are professors Ruzhang Ma and Zhengwen Li of University of Science and Technology of Beijing.


this lamellar phase being perpendicular to c axis appeared in only part of whole visual field. At 780 this lamellar phase being perpendicular to c axis appeared at most of all the visual field and at 790 the lamellar phase being perpendicular to c axis appeared on base of original cellular texture. Along with continual rising of temperature the lamellar phase was increased gradually. Fig. 3.14 shows the long lamellar phase grown in the in situ at 810 . Fig. 3.16 shows the long lamel-lar phase grown in the in situ at 830 and Fig. 3.17 shows the lamellar phase at 840 and holding for 0.5h. It can be seen from Fig. 3.17 that the lamellar phase ran through the whole visual field.

The Mössbauer Effect experiment: Ascription of sub-spectral crystal site of Mossbauer spectroscopy was fitted and determined basis on analysis of RE2Co17

and RECo5 crystal structures, RE2Fe17 Mössbauer effect experiment, NMR (Nu-clear Magnetic Resonance) experimental result of RE2Co17 and RECo5. Crystal site marker, crystal sites Co1, Co2, Co3 and Co4 corresponds to 18g, 18h, 6c and 9d in RE2Co17 diamond structure and crystal site 12j, 12k, 4f and 6g in RE2Co17 hexagonal structure.

Average hyperfine field Hhf was obtained by weighted sum of each crystal site Hhf (f) and area percentage A3, that is

hf hf1

( ) ( )N

iH H i A i

�� %

The 2:17 phase sub-spectral area increased from 75.1% in 0Zr specimen to 78.6% in specimen of 2.4Zr. It can be seen that iron atoms entered into the 2:17 phase from 1:5 phase by adding zirconium. The average hyperfine field Hhf of the 2:17 phase increased from 298.4kOe in specimen of 0Zr to 301.0kOe in specimen of 2.4Zr, and the average hyperfine field Hhf of the 1:5 phase decreases from 272.4kOe in specimen of 0Zr to 270.2kOe in specimen of 2.4Zr. These indicated that zirconium could make nonmagnetic atoms, zirconium and cupper, decom-posed from the 2:17 phase to the 1:5 phase because exchange interaction of Fe-Fe in the 1:5 phase (or the 2:17 phase) reduced (or increased) magnetic moments of iron and thus reduced (or increased) hyperfine field at 57Fe nuclei. This is consis-tent with the result of energy spectrum analysis that zirconium enriched phase appeared in specimen of 2.4Zr. According to opinion of Panl, coercivity of alloy was related to different in magnetic characters between two phases (magnetocrys-talline anisotropy, saturation magnetization intensity and exchange interaction) as well as size and distribution of defects. Therefore, addition zirconium increased nonmagnetic atoms in the 1:5 phase and reduced magnetism of the 1:5 phase, and made domain wall intensified and the coercivity enhanced.

Addition of zirconium can promote iron atoms enter into Co3 crystal site from Co1 crystal site. In the 2:17 phase Co1 crystal site has strong axial anisotropy, Co3 crystal site is of strong plane anisotropy, and Co2 and Co4 crystal sites only have


weak plane anisotropy. And that iron atom replacing cobalt will weaken anisot-ropy. Therefore increase in number iron atoms occupying Co3 crystal sites will weaken plane anisotropy of these crystal sites and decrease in number iron atoms occupying Co3 crystal sites will enhance the axial anisotropy of these crystal sites. As the reduction it is to enhance the single axial anisotropy in these crystal sites. The result is increasing alloy’s single axial anisotropy and so that to be avail of enhancement of coercivity (Strnat, 1983; Pan, Xiao, 1989). The process proved that zirconium played a key role in improving alloy coercivity. It is advantaged also to enhance quadrate degree in demagnetization curve. Adding zirconium can increase solubility of samarium in the alloy that the solubility of Samarium in Sm2(Co,Cu,Fe,M)17 is increased by 27%.

It can be seen from Fig. 3.30 that coercivity mHc of the 25.2Sm-Co-6Cu-15Fe-Zr alloy increased along with increasing of zirconium content, and the ratio of Hk/mHc also increased with increasing of zirconium. Hk/mHc indicated ridgy degree, or called quadrate degree of demagnetization curve. The quadrate degree has di-rectly related to big and small of Br, Hc and Hk and also related to components and heat treatment process of the alloy. People always want to get demagnetization curve with good quadrate degree. In Sm(Co,Cu,Fe,M)z alloy mHc was trebled to reach 2000kA/m with zirconium content of 2% (at.) in compared with content of zirconium (at.) 0% (in condition TbCu7 structure, and single-phase region 11.5%-12.5%(at.)Sm).

Fig.3.30 The relation of mHc, Hk/mHc vs Zr content of 25.5Sm-Co-6Cu-15Fe-Zr alloy

(Fukui, Nishio, Iwama, 1987)

Adding zirconium into zirconium enriched 2:17 phase made 2:17 phase extend-

ing to cobalt enriched and samarium enriched sides. Then solubility of samarium


in the 2:17 phase was enhanced to 13% (at.) from 9.5% (at.) at 1370 and thus changed structure of high temperature phase and fully exerted effect of zirconium. For the 2:17 alloys containing zirconium multistage aging were adopted (as shown in Fig. 3.20). When content of zirconium is less than 1.5% isothermal ag-ing can be controlled for 30min to 40min. When Zr > 2.0% (at.) isothermal aging is controlled in 9-35h. Respective aging temperatures and holding times are at 700-600 for 1h, 500 for 2h, and 400 for 4h. For alloys with high content of zirconium generally cool to 400 then holding for 3h with the cooling speed of 0 .3%-1.5%/ . For permanent magnet with Zr 3% (wt.) isothermal aging at 850 makes against enhancement coercivity. The alloy no matter with a low-zirconium or a high-zirconium have to process classifying aging after isothermal aging at 820-850 and the cooling speed is 0.3-1.2 /min in order to obtain high performance permanent magnetic alloy.

From above analysis, we can know that zirconium plays a key role for en-hancement of coercivity of the Sm (Co, Cu, Fe, M)7.4 alloy.

3.2.3 Conclusions

The following conclusions are derived based on above researches: 1. Dynamic observation on Sm (Co, Cu, Fe, Zr)7.4 using 1000kV TEM proves,

appearance temperature of zirconium enriched long lamellar phase is 790 . Holding at 840 for 0.5h this texture runs through whole view field. This long lamellar structure is not the direct reason to cause produce high coercivity.

2. According to the Mössbauer research addition of zirconium can promote iron enter into the 2:17 phase from the 1:5 phase. Meanwhile nonmagnetic ions, cop-per, zirconium, etc. enter into 1:5 phase from 2:17 phase. These increased the difference in chemical composition and magnetic character between two phases. Addition of zirconium promote iron atoms in the 2:17 phase enter into Co3 crystal sites from Co1 crystal sites, consequently enhance uniaxial anisotropy of the alloy. And these two points both are favorable to enhance coercivity of the alloy.

3. Zirconium in Sm (Co, Cu, Fe, Zr)7.4 alloy acts enhancing coercivity, increas-ing solubility of samarium in the alloy, and improving quadrate degree of demag-netization curve.

3.3 Magnetism of High Coercivity Sm(Co, Cu, Fe, M)7.4 Per-manent Magnetic Alloy at 1.5-523K

High coercivity Sm(Co, Cu, Fe, Zr)7.4 permanent magnetic alloy is wildly used because of its high Curie temperature (T c =810 ) and high saturation magneti-zation intensity (Pan, Ma, Li, 1993; Pan, Ma, Ping, et al, 1991; Pan, Jin, 1990; Pan, Zhao, 1988). With the special requirement in the areas of aviation and navi-


gation material properties must be measured at high temperature and low tem-perature. High properties material, especially rare earth permanent magnetic al-loys with high coercivity, used in power device, magnetic force damper, magnetic force driver, etc. operating in media of liquid hydrogen and liquid nitrogen of spaceflight and operating at low temperature being specified by user. The de-signer must give the magnetism and the magnetism under repeated variation through high temperature to low temperature in order to effectively and reasona-bly selecting magnet. The permanent magnetic alloy used for aerospace perma-nent magnetic motor must be made with good stability in temperature and time at 523K. Therefore measuring and studying the magnetism of material at 1.5-523K are necessary.

3.3.1 Preparation of specimen and magnetism measurement apparatus and measurement method

Composition range of alloy: 23.5%-27.5%(wt.) Sm, 13.5%-25.5%Fe (mass frac-tion), 4.0%-8.2%Cu (mass fraction), 1.0%-4.1%Zr (mass fraction), the surplus was cobalt. For example, Sm(Co, Cu, Fe, M)7.4 specimen composition is 26.01% Sm 51.38% Co 4.36% Cu 15.19% Fe 2.99% Zr. The purity of raw metal materials was 99.5%-99.8%. Put above material in vacuum induction furnace and melt in argon atmosphere, cast into water cooling copper crucible, the melted al-loy ingot was crushed and pulverized grossly, intermediately and finely to powder of 3-5m; the powder was formed in magnetic field greater than 1.5T and the forming pressure was 2.5-3T/cm2 with a direction being perpendicular to mag-netic field; the above pressed billet was sintered for 1h at 1150-1210 , sintered billet was through solid solution treatment at 1130-1200 for 1-2h, then through isothermal aging (720-900 ) or step aging. The billet was cut as per require-ments, then was magnetized in a pulse magnetic field greater than 5T and meas-ured magnetism with CL-6 direct current parameter measurement instrument.

Measuring sequence is as follows: 25 , 100 , 150 , 200 , 250 , 25 , 60 , 196 , 25 . The specimen was measured by vi-brating specimen magnetometer in condition of 1.5K and 40K (Zhang, Cheng, 1982; Pan, Ping, Liu, Ma, 2003; Ma, et al, 1999).

3.3.2 Measurement results and discussion

Specimens with different composition and heat treatment system were measured. The curve of one composition of specimens measured at 196 to 200 is shown in Fig. 3.31, the demagnetization curve under 1.5K is described in Fig. 3.32 and magnetic properties under 1.5K are given in Table 3.4 and Table 3.5.


Fig. 3.31 The demagnetization curves of high coercivity Sm(Co, Cu, Fe, Zr)7.4 permanent

magnet alloy measured from 200 to 196 1 200 ; 2 150 ; 3 100 ; 4 25 ; 5 60 ; 6 196

Fig. 3.32 Demagnetization curve of Sm(Co, Cu, Fe, Zr)7.4 at 1.5K

(It is shown from the demagnetization curve that spin realignment appears at low temperature) Magnetic properties at 25 are as follows: Residual magnetic induction intensity

r 1.05T(10.5kGs)B �

m c 1,592kA/m(20kOe)H �

b c 716kA/m(9kOe)H � 3

max( ) 206.96kJ/m (26MGs Oe)BH � �

Magnetic properties at 196 : r 1.10T(11.0kGs)B �


m c 1,990kA/m(25kOe)H � b c 765kA/m(9600Oe)H �

max( ) 214(27MGs Oe)BH � �

Magnetic properties at 200 :

r 0.92T(9.2kGs)B �

m c 119.4kA/m(15kOe)H �

b c 655.0kA/m(8200Oe)H �

max( ) 195.0(24.5MGs Oe)BH � �

Table 3.4 Magnetic performance of Sm(Co, Cu, Fe, Zr)7.4 at 1.48-1.53K Havant/Oe M/emu·g 1 Mmoy/emu·g 1 Tbas/K T/K H/Oe

97 93.829 93.829 1.61 1.53 92

402 94.124 94.124 1.61 1.52 399

700 93.861 93.861 1.60 1.52 698

1,983 91.200 91.200 1.60 1.51 1,980

3,996 85.403 85.403 1.59 1.51 3,994

6,010 80.128 80.133 1.59 1.51 6,005

7,984 75.113 75.122 1.58 1.50 7,984

9,994 70.419 70.409 1.58 1.50 9,996

12,001 63.652 63.640 1.58 1.49 12,005

13,999 55.574 55.574 1.58 1.49 14,004

15,998 47.567 47.548 1.58 1.49 16,004

17,998 39.664 39.664 1.57 1.49 18,005

19,999 32.055 32.055 1.57 1.48 20,007

20,476 30.515 30.515 1.57 1.49 20,482

21,002 22.758 22.758 1.57 1.48 21,009

21,497 22.343 22.340 1.57 1.48 21,505

21,991 21.526 21.519 1.57 1.49 21,998

22,497 19.962 19. 962 1.57 1.48 22,504

23,000 19.185 19.185 1.57 1.48 23,007

23,494 17.718 17.718 1.57 1.48 23,500

23,990 15.993 15.993 1.57 1.48 23,996

24,498 6.243 6.243 1.57 1.48 24,506

25,000 5.886 5.886 1.57 1.48 25,008

25,494 5.086 5.086 1.58 1.48 25,501

25,990 4.042 4.042 1.58 1.48 25,996

26,498 2.060 2.060 1.57 1.48 26,506

26,999 1.040 1.040 1.57 1.48 27,006

27,497 11.249 11.250 1.58 1.48 27,504


Table 3.5 Magnetic performance of Sm(Co, Cu, Fe, Zr)7.4 at 1.48-1.55K Havant/Oe M/emu·g 1 Mmoy/emu·g 1 Tbas/K T/K H/Oe

87 91.336 91.336 1.64 1.55 82

400 91.499 91.501 1.64 1.55 398

698 91.322 91.327 1.64 1.54 696

2,011 88.831 88.831 1.63 1.54 2,008

3,980 85.877 85.880 1.62 1.53 3,978

5,980 82.021 82.019 1.61 1.52 5,979

7,979 78.618 78.618 1.60 1.52 7,978

9,994 74.070 74.077 1.59 1.51 9,998

11,975 68.805 68.793 1.59 1.50 11,978

13,999 60.445 60.432 1.58 1.50 14,005

16,000 51.362 51.362 1.58 1.49 16,004

17,999 40.716 40.708 1.57 1.49 18,005

19,999 31.367 31.367 1.57 1.49 20,006

20,478 30.066 30.066 1.57 1.49 20,484

20,999 27.489 27.490 1.57 1.48 21,007

21,500 26.320 26.316 1.57 1.49 21,506

21,999 24.672 24.672 1.57 1.49 22,005

22,487 18.284 18.284 1.57 1.49 22,493

23,000 17.395 17.395 1.57 1.48 23,006

23,498 16.720 16.720 1.57 1.48 23,506

23,998 15.438 15.438 1.57 1.48 24,004

24,488 13.857 13.857 1.57 1.48 24,495

25,005 7.449 7.449 1.57 1.48 25,012

25,500 6.834 6.834 1.57 1.48 25,505

25,996 5.475 5.475 1.57 1.48 26,003

26,488 0.095 0.095 1.57 1.48 26,495

27,004 1.698 1.699 1.57 1.48 27,012

27,502 8.269 8.269 1.57 1.48 27,507

27,994 8.981 8.981 1.57 1.48 28,001

3.3.3 Conclusions

Conclusions are drawn base on the above researches as follows 1. Magnetic properties of Sm (Co, Cu, Fe, M)7.4 varied along with change of

heat treatment system.


2. The Sm(Co, Cu, Fe, M)7.4 permanent magnetic alloy with different ingredient has different magnetic properties. Zirconium content have large influence on co-ercivity, iron content impact on Br and (BH)max, copper can promote precipitation of the 1:5 phase and improve coercivity.

3. Properties of permanent magnetic alloy at low temperature are better than that at room temperature.

4. Properties of Sm(Co, Cu, Fe, M)7.4 permanent magnetic alloy can be reversed to original value at room temperature through reversing low temperature to room temperature.

5. It is testified from demagnetization curve that spinning and then orientation appeared in measurement of Sm(Co,Cu,Fe,M)7.4 permanent magnetic alloy under 1.5K.

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Chapter 4 The Third Generation Rare Earth Permanent Magnet

This chapter discusses the 3rd generation rare earth permanent magnetic alloy. The 3rd generation rare earth permanent magnetic alloy has a world-record mag-netic energy product, mainly due to its tetragonal crystal system of Nd2Fe14B, which broke through the hexagonal system and rhombohedral system of perma-nent magnetic alloys of the 1st and 2nd generations.

Nd2Fe14B compound is the base phase of RE-Fe-B family permanent magnetic alloy. This chapter stresses on the formation of Nd2Fe14B phase, analysis of micro area of Nd2(Fe,Co)14B crystal, and the in situ and dynamic observation by TEM (transmission electron microscope) on Nd2Fe14B and Nd2(Fe,Co)14B at high tem-perature. Research on B-rich phase in Nd-Fe-B alloy was carried out as follows: the in situ and dynamic observation on Nd1.11Fe4B4 under TEM, phase analysis on Nd1.11Fe4B4, and relationship between B-rich phase and coercivity. The alloy of Nd16Co16Fe59Ga2B7 could raise the Curie temperature from 312 of Nd16Fe76B8 alloy to 450-500 , raise the intrinsic coercivity to jHc = 716-955 kA/m, and raise the maximum magnetic energy product to (BH)max = 223-262 kJ/m3. The in situ and dynamic observation and analysis were conducted on magnetism and high temperature phase transformation of Nd15Fe78B7 and Nd16Fe69Co8B7 permanent magnetic alloys. Lastly, this chapter discusses the model of coercivity and nano-meter microstructure of NdFeB permanent magnetic alloy containing niobium and gallium.

4.1 Improvement of the Properties of NdFeB Permanent Mag-nets Due to Element Substitutions

NdFeB permanent magnets set a record of magnetic energy product. The product of NdFeB reaches 446.24 kJ/m3 (55.78 MGs�Oe), Br=1.514 T (15.14 kGs) and mHc=694.4 kA/m3. However, the Curie temperature is rather low and only 312 .


Therefore, the temperature stability is poor and the working temperature is rather low (<80 ). These drawbacks limit the application range of NdFeB permanent magnets. In order to overcome these drawbacks, the effect of different elements on the properties of NdFeB permanent magnetic alloys using substitution methods. The related studies can be classified into three types:

Substitution of other rare earth elements for Nd; Substitution of other metal or non-metal elements for Fe; Substitution of other non-metal elements for B.

The part substitution of heavy rare earth, such as Dy, Tb and Er, for Nd can in-crease the coercivity. However, as the cost is more expensive for Dy than for Nd, the cost of NdFeB will increase about 50 % even if the substitution of Dy is about 10%. To reduce the cost, the coercivity of NdFeB alloys can be increased by the substitution of rare earth oxides, Dy2O3 and Tb2O3, for Nd through an appropriate technology(Sagawa, Hirosawa, Tokuhara, et al, 1987; Pan, Ping, Liu, et al, 2003; Pan, Pan, Ma, 1994).

The substitution of most transition elements in the Period Table for Fe has been investigated. The substitution is mainly due to an increase in Curie temperatures and a decrease in temperature coefficients. From experiment results, Al, Si and Ga substitutions for Fe can achieve above goals in a certain sense. In addition, only Co can be substituted for Fe in the whole composition range and, at the same time, Nd2Fe14B phase is maintained. In Nd2Fe14-xTxB, possible substituted elements and their numbers x are 1.5Al, 1Si, 0.2Cu, 1Ni, 3Mn, 3Cr, 1V, 0.5Nb, 0.2Zr and 0.2Ti. If iron is substituted by other elements, not only the residual inductivity will be reduced greatly but also the magnetic energy product of NdFeB will be reduced (Pan, et al, 1987; Pan, Li, Ying, et al, 1990; Pan, Li, Li, et al, 1989; Li, Ping, Ma, 1988). Cobalt substitution for Fe can significantly increase the Curie temperature, thus greatly reducing the reversible temperature coefficient and improving the thermal stability. The Curie temperature is increased in a certain scale by Ni sub-stitution, and is not improved by Al and Si substitutions. However, an appreciate substitution of Al and Si can improve the coercivity. Mn and Cr substitutions are not of benefit to the permanent magnetic-properties due to the antiferromagnetic coupling between Fe and Mn or Cr.

Based on the analyses of X-ray diffraction and SEM with EDAX, as Co is partly substituted for Fe, there are at least four phases in Nd15(Fe1-xCox)78B7 at x0.2, i.e., the tetragonal phase Nd2(Fe, Co)14B, B-rich phase Nd1+� (Fe, Co)4B4 (�=0.1-0.3), Nd-rich phase, and Laves phase Nd(Fe, Co)2 appearing at the grain boundary of the tetragonal phase. Among them, the tetragonal phase is dominant, and the others are insignificant.

The part substitution of Co for Fe improves the temperature stability, because the exchange constants of JCo-Co and JCo-Fe are larger than JFe-Fe. Consequently, the


permanent magnetic-properties can maintain to higher temperatures. The substitu-tion also improves the corrosivity-resistant of NdFeB alloys. The magnets pre-pared from Nd12Dy0.6Co5.8Zr0.1Ga0.3Fe74.5B7 has the magnetic properties: Br=0.87 T, mHc=1530 kA/m and (BH)max=137 kJ/m3. The temperature coefficients are

rB� =�0.11 %/ and cH� =�0.43 %/ (Sagawa, Hirosawa, Tokuhara, et al,

1987; Zhou, 1995). An appreciate substitution of Mo can increase the coercivity and reduce the ir-

reversible loss of open-circuit magnetic flux of NdFeB. Based on the analysis, Mo exists the tetragonal phase and the RE-rich phase and aggregates in the B-rich phase (Liu, Zhou, 1990).

The fast-quenched permanent magnets Nd9Fe85-xMnxB6 (x=0, 0.5 and 1) are prepared using Mn substitution. A small amount of Mn can improve the perma-nent magnetic-properties of the fast-quenched samples. The coercivity is in-creased from 339.6 kA/m to 398.2 kA/m (Xie, Yin, Jiang, et al, 2002).

For Nd8-xDyxFe83.5Co2Nb1Ga1B4.5, with increasing Dy concentration, the coer-civity first increases, reaches the maximum at x=0.5-0.8 and then decreases at x>0.8. The increase in the intrinsic coercivity with increasing Dy concentration is because the Dy substitution increases the anisotropy field of samples, and the decrease at x>0.8 has its origin in the large grains and weak exchange interaction (Cheng, Gao, Zhu, et al, 2002).

Al substitution can increase the coercivity. The combination of Al and Co sub-stitutions can increase both the Curie temperature and the coercivity. Nd16(Fe0.96Al0.04)76B8 reaches the coercivity of 1,114.4 kA/m, the remanence of 1.10 T and the energy product of 238 kJ/m3. Nd16Co10Fe67-yAlyB7 with y=0, 2, 4 and 8 is prepared by melting, ball-mill, shaping in a applied magnetic field, sinter at high temperature, annealing, and magnetizing at the field of larger than 4 T. The results of magnetic measurement showed that Br=0.94-1.28 T, and at y=3 the coercivity achieves 1,130 kA/m, and the maximum energy product maintains 242.7 kJ/m3. The Curie temperature of the magnetic alloy is 710 K due to 10 % Co substitution for Fe. It is note that a soft magnetic phase appears as the Co sub-stitution is more than 20 %. Mössbauer spectra showed that the line-widths be-come broader and the hyperfine fields decrease. Al atoms preferentially occupy the j2 site, which has strong easy-planar anisotropy. The Al substitution for Fe and Co can decrease the easy-planar anisotropy of the j2 site, and thus increasing easy-axial anisotropy of the tetragonal phase (Pan, Pan, Ma, 1994; Zhou, 1995; Zhang, Ma, et al, 1985).

The two-phase technology has become a new method for manufacturing sin-tered Nd-Fe-B magnet in recent years (Zhang, Liu, Zhao, et al, 2009). The two-phase composite, Nd2Fe14B/ �-Fe, has potential good magnetic properties. The doping of Cu, Nb, Mo, Cr, Co and La can effectively modify the structure and


improve magnetic properties. The doping of Cu and Mo can effectively improve properties of permanent-magnet alloys by the grains. The part substitution of other rare earth alloys for Nd can increase the coercivity. The part substitution of Co for Fe can increase the saturation magnetization and the Curie temperature. The part substitutions of Al for Fe can greatly magnetic properties of the alloys after heat-treatment. In the same quenched technology, amorphous state signifi-cantly increases with increasing Al concentration. The addition of Al elevates the crystallization temperature of Nd2Fe14B, restrains the growth of grains and grain. As a result, the magnetic properties are improved (Cheng, Gao, Zhu, et al, 2002; Wang, Zhao, Cui, et al. 2001).

Zn is added into Nd2Fe14B by mechanical alloying. During ball-mill, Zn is added to ball-milled machine and is homogenously mixed with Nd2Fe14B. Then the mixture is annealed in high vacuum ((2.0-4.0)×10�5 Torr or (2.67-5.33)× 10�3 Pa). The addition of Zn of 5% (wt) can significantly increase the coercivity. As the addition of 20%, mHc reaches the maximum; it is about 2.3 times (from 75.3 kA/m to 179 kA/m), as compared to original sample. Zn is mainly introduced to the Nd-rich phase and increases the pinning to the reverse-phase domains.

The addition of Sn or FeSn2 can increase mHc. For the addition of 25%, mHc is 43.78 kA/m and increases 60%, as compared to the original powders.

Using Zn, Sn, etc. of low melting point as bond, processing heat treatment at near the melting point, and using ball miller to add Zn and Sn in pre-melting of Nd2Fe14B and Sm2Co17 can obtain above results (Xiong, Xiong, 1993).

After a small amount of Nb is substituted for Fe, the coercivity of rare earth iron permanent-magnet alloys is increased. Most of Nb exists in phase as impurity. Based on the analysis of TEM, besides the Nd-rich phase, B-rich phase (Nd1+�Fe4B4, �=0.1-0.3) and main phase Nd2Fe14B, one observed a large amount of fine phase with homogeneous distribution exists stacking fault exists in the interior of the grains. In addition, Laves phase Fe2Nb is also observed in Nd2Fe14B. The addition of small amount of Nb can fine the grains, and thus in-crease mHc and Hk. The sintering technology of sample containing Nb is different from that of Nd15Fe77B7. The aging temperature is 550-900 . It is better to be divided into three steps for aging treatment, 900 for 2 h and 600° for 1-2 h. Another effect of Nb addition is to restrain the of �-Fe phase. Moreover, the addi-tion of Nb can reduce the irreversible loss. Fe2Nb has MgMn2 structure with lat-tice parameters a=0.482 nm and c=0.787 nm. The Laval phase exists in the matrix of permanent-magnet alloy as particles, which contain Nb of 22%-44% and have a diameter of 2m. The Laval phase is non-magnetic, is structure with main phase Nd2Fe14B and is more Nb than main phase. As Nb>4 %, the main phase Nd2Fe14B is destroy. Based on the experience of the author, it is better for the addition of Nb to be 1%-2% (at.) (Pan, Ping, Liu, et al, 2003; Liu, Pan, Luo, et al, 1991).


The substitution of a small amount of Ga for Fe can greatly increase coercivity and reduce the irreversible loss. The irreversible loss is 1.8 % for the ageing at 200 and is 1.2 % for the ageing 400 . The sample shows good stability. The combining substitution of Ga and Nb can significantly improve the stability of permanent-magnet alloys. Nd16Co16Fe60Ga2B6 has relatively large coercivity when the annealing at 640 is used and the addition of Ga is 2% (at) (Pan, Li, Li, et al, 1989; Liu, Pan, Luo, et al, 1991).

Substitution of Gd for Nd in Nd17-xGdxFe76.5B6.5 decreases �-Fe precipitation. Nd-Fe-B with Gd substitution has less magnetic saturation 4�M and HA. Not more than 5% Gd substitution can get proper magnetic performance (Liu, Zhao, Zhao, et al, 2010).

Zhou Guojun, et al have applied CALPHAD (Calculation of Phase Diagram) to assess thermodynamical Nd-B binary system which gives guide in composition design of Nd-Fe-B alloy (Zhou, Zeng, 2010).

4.2 Magnetic Properties and the Occupancy of Co and Ga At-oms for NdFe(Co, Al, Ga)B Permanent-Magnetic Alloys *

NdFeB permanent-magnet alloys have excellent magnetic properties. However, as compared to Sm-Co alloys, the relatively low Curie temperature (only 312 ) and therefore poor thermal stability limit the applications of NdFeB. In order to im-prove the thermal stability, it is necessary to increase the Curie temperature of NdFeB. Soon after NdFeB was discovered, it has been known that Co substitution for Fe can greatly increase the Curie temperature (Sagawa, Fujimura, Togawa, et al, 1984; Pan, et al, 1987; Xu, Ping, Li, et al, 1986; Croat, Herbst, Lee, et al, 1984; Ma, et al, 1987; Tokanaga, Tobise, Meguro, et al, 1986; Szafra�ska-Miller, Ptusa, Wystocki, et al, 1987; Mizoguchi, Sakai, Niu, et al, 1986). However, the substitu-tion reduces the coercivity of NdFeB. The coercivity can be increased by the small addition of Al and Ga in NdFeCoB. Based on Mössbauer spectra, the distri-butions of Co, Al and Ga atoms at the sites and the relationship between the dis-tribution and magnetic properties, such as Curie temperature, have been studied.

4.2.1 Preparation and method

Samples of Nd16Fe77-xCoxB7 with x=0, 4, 8, 16 and 24, Nd16Co16Fe61-xGaxB7, Nd16Co10Fe67-xAlxB7, and Nd16Co16Fe61-xAlxB7 with x=0, 1, 2, 4 and 7, were pre-pared by arc melting under an Ar atmosphere. The starting materials were 99.8% Nd, 99.5% B, 99.9% Fe in purity. The ingots were remelted three times in order to

Cooperators of this study are professors Ruzhang Ma, Jueyun Ping and Zhengwen Li of Univer-sity of Science and Technology Beijing who conducted most of the Mössbauer analysis.


achieve homogeneity. Then, the ingots were crashed and ball-milled to powders with about 3.5 m. The powders were shaped in a magnetic field with 1.5 T. Finally, the samples were sintered at 1100 and annealed at 500-630 , followed by quenched at room temperature. The prepared samples were magnetized in a pulse strong magnetic field. The permanent magnetic parameters were measured using DC Parameter Measurement Instrument and the Curie temperatures were obtained from magnetic balance.

X-ray diffraction was performed using APD-10X diffract-meter with Cu K ra-diation. Mössbauer spectra were collected in a conventional constant-acceleration spectrometer. The X-ray source was 57Co in Rh matrix. The calibration was made using the spectrum of �-Fe spectrum at room temperature (Xu, Ping, Li, et al, 1986). All Mössbauer absorbers were the powdered samples and contained about 5-10 mg/cm2 of natural iron. The Mössbauer spectra were analyzed using a singlet, a doublet and six sextets, which are associated with Nd-rich phase, B-rich phase and six inequivalent Fe sites in the tetragonal phase.

The occupancy of Fe atoms at six sites in the tetragonal phase Nd2Fe14B is N1=2/4, N2=4/14, N3=2/14, N4=4/14, N5=1/14 and N6=1/14, where Ni (i=1-6) represents the j2, k2, j1, k1, e and c site, respectively. If it is assumed that the recoil-free fractions on the six sites are the same, the occupation fractions Ni(Fe), Ni(Co) and Ni(Al) of Fe, Co and Al atoms can be estimated by the following formulae:

1

61 2

1

1(Fe) 100%

(Co) 100% (Fe)

ii

ii

i

i i

SCNC C N

S

N N�

� � � � ! � �#

% (4.1)

for NdFeCoB permanent-magnet alloys, and

16

1 2 3

1

,0 ,0

1(Fe) 100%

(Co) (Co) 100% (Fe)(Al) 100% (Fe) (Co)

ii

ii

i

i i i

i i i

SCNC C C N

S

N N NN N N

�

� � � � � ! � � � � � �#

% (4.2)

for NdFeCoAlB permanent-magnet alloys, where C1, C2 and C3 denote the con-centrations of Fe, Co and Al in the alloys (Croat, Herbst, Lee, et al, 1984), Si de-notes the area of Mössbauer subspectral at the i site, Ni,0(Fe) and Ni,0(Co) denote the occupancy in the alloys without Al element.

4.2.2 Nd16Fe77-xCoxB7 alloy

The magnetic properties of Nd16Fe77-xCoxB7 are listed in Table 4.1.


Table 4.1 Magnetic performance and Tc of alloy Nd16Fe77-xCoxB7

Cobalt content, x Br/T jHc/kA�m�1 (BH)max/kT�m�3 Tc/ 0 1.26 811.72 295.24 318

0.05 1.23 987.84 280.12 374 0.1 1.22 373.23 275.35 433 0.2 1.11 344.58 227.60 490 0.3 1.08 289.67 215.66 553

From Table 4.1, with Co substitutions, the Curie temperature significantly in-

crease, but the coercivity, remanence and maximum energy product will greatly decrease.

The Mössbauer spectra of NdFeCoB are shown in Fig. 4.1(a). Based on the curve-fitted results, the occupation fractions of Co atoms in the tetragonal phase Nd2(Fe, Co)14B are shown in Fig. 4.1(b). Co atoms preferentially occupy the k2 and j2 sites. In the six sites, the k2, j1 and j2 sites play important roles in elevation of the Curie temperature. First, the j2 site has most neighboring atoms (the num-ber is 12), and thus contributing strong positive exchange interaction. On the other hand, the distances of j1-k2 and j1-j1 sites are 0.2396 nm and 0.2433 nm, which are less than the critical distance 0.245 nm between positive and negative exchange interactions. These two sites contribute the negative interaction. The Co atoms preferentially occupy the j2 site, which enhances the positive interaction, due to the exchange interaction of JCo-Co>JCo-Fe>JFe-Fe. The Co atoms prefer the k2 site, which decreases the negative interaction of k2-j1 sites. Consequently, the total interaction is enhanced and thus increasing the Curie temperature with Co substi-tutions. In addition, with Co substitutions, the occupancy of the j2 and k2 sites in-creases rapidly at x 8 than at x>8. Therefore, the increase in the Curie tempera-ture is elevated much fastest (Pan, 1986; Li, Ping, Ma, et al, 1988; Yang, et al, 1985; Yang, et al, 1989; Ma, et al, 1992).

Fig.4.1 The Mössbauer spectra of NdFeCoB (a) and the occupation fractions of

Co atoms in the tetragonal phase Nd2(Fe, Co)14B (b)


In the tetragonal phase Nd2Fe14B, the j2 site has strongest planar anisotropy. The contribution on anisotropy is larger for Co atoms than for Fe atoms. There-fore, the preferential occupancy of Co atoms at the j2 site enhances the easy-plane anisotropy, and thus reducing the total uniaxial anisotropy. This leads to a de-crease in coercivity. In addition, the appearance of soft magnetic phase Nd(Fe, Co)2 at high Co substitution also reduces the coercivity of NdFeCoB.

With Co substitution, the hyperfine fields on the six sites slight decrease.

4.2.3 Nd16Co10Fe67-yAlyB7 and Nd16Co16Fe61-yAlyB7 alloys

The addition of an appropriate amount of Al into NdFeB or, on the other word, the partial substitution of Al for Fe can increase the coercivity. The combined substitution of Co and Al can not only elevate the Curie temperature, but also increase the coercivity. To avoid the soft magnetic phase Nd(Co, Fe)2, the concen-tration of Co is taken 10%.

For Table 4.2, with Al substitutions, the remanence Br decreases, and the coercivity has the maximum; the coercivity is 1,130 kA/m at y=4 for Nd16Co10Fe67-yAlyB7. The corresponding maximum of the Curie temperatures is found to be 437 . However, the value is not high enough. An increase in Co concentration of 16, the permanence magnetic properties are as shown in Table 4.3 for Nd16Co16Fe61-xAlxB7.

Table 4.2 Magnetic performance and Tc of alloy Nd16Co10Fe67-yAlyB7

y 0 2 3 4 6 8 10

Br/T 1.28 1.19 1.13 1.07 0.99 0.94 0.62

jHc/kA�m�1 572 796 1050 1130 726 432 427

(BH)max/kJ�m�3 302 271 242.7 215 196 143 58

Tc/ 417 431 437 400.5

Table 4.3 Magnetic performance and Tc of alloy Nd16Co16Fe61-xAlxB7

x 0 1 2 4 7

Br/T 1.22 1.17 1.11 1.04 0.98

jHc/kA�m�1 565.0 764.0 988.0 923.1 350.2

(BH)max/kJ�m�3 270.1 246.7 238.7 198.8 127.3

Tc/ 501 478 481 451

From Table 4.3 and Fig.4.2, the coercivity of Nd16Co16Fe61-yAlyB7 has the

maximum at y=2, and as y>2, the coercivity decreases. The remanence and energy product reduce monotonously with Al substitutions. The demagnetizing curves of NdFeCoAlB do not obviously change, as compared to NdFeB (Fig. 4.3). In addi-tion, for series of NdFeCoAlB, as the Co concentration increases from 10 to 16, the Curie temperature is also elevated.


Fig.4.2 Curves for relation of remnant magnetization Br and intrinsic

coercivity jHc to Al content x 1 Curve of Br; 2 Curve of jHc

Fig.4.3 Demagnetization curves of Nd16Fe61-xAlxCo16B7 alloys 1 x=0; 2 x=1; 3 x=2; 4 x=4; 5 x=7

The Mössbauer spectra of NdFeCoAlB are shown in Fig.4.4(a). Based on the

curve-fitted results, the occupation fractions of Al atoms in the tetragonal phase Nd2(Fe, Co, Al)14B are shown in Fig. 4.4(b). The results showed that Al atoms mainly occupy the j2 sites as y 4, and also preferentially occupy the k1 site at y=8.

Al is a non-magnetic atom. Al substitution for Fe and occupation of the j2 site will decrease the planar anisotropy, thus increasing the total uniaxial anisotropy. This is an important reason that the coercivity increases from 572kA/m to 1, 130kA/m, as Al concentration y varies from 0 to 4 for Nd16Co10Fe67-yAlyB7

(Ping, Li, Ma, et al, 1986). As different from Co substitution, Al substitutions lead to rapid decrease in the

hyperfine fields at all sites; however, the slope is different. The maximum slope is found at the j1 site, and minimum slope at j2 and c sites. Al is non-magnetic atoms and mainly occupies the j2 site. The j1 site has four adjacent j2 atoms, and the j2 and c sites have no j2 atoms. Therefore, the hyperfine field decreases rapid for the


j1 site and slowly for the j2 and c sites with Al substitutions (Pan, Zhao, Ma, 1988).

Fig.4.4 The M ssbauer spectra of NdFeCoAlB (a) and the occupation fractions of Al

atoms in the tetragonal phase Nd2(Fe, Co, Al)14B (b)

4.2.4 Nd16Co16Fe61-xGaxB7 alloy

As Ga and Al are the same group in the Periodic Table, the substitution of Ga for Fe is shown to be able to improve the permanent magnetic properties of NdCoFeB, as similar to Al. The properties as well as the Curie temperature are listed in Table 4.4.

From Table 4.4, with Ga substitutions, the intrinsic coercivity increases, reaches to maximum, mHc=938.8kA/m, at x=2 and then decreases. The remanence slightly decreases with Ga substitutions.

The Curie temperature reduces slowly for the Ga substitution x 2; it decrease from 501 to 475 as x varies from 0 to 1, and only decreases to 455 as x increases to 7. As compared to the combined substitution of Co and Al, the Curie temperature reduces much slowly for the substitution of Co and Ga. In conclusion, the partial Ga substitution for Fe opens a way that the coercivity increases, but the Curie temperature only slightly decreases.


Table 4.4 Magnetic performance and values of Tc of alloy Nd16Co16Fe61-xGaxB7 corre-sponding to Ga content

x 0 1 2 4 7

Br/T 1.22 1.19 1.10 1.10 0.99

jHc/kA�m�1 680.2 708.9 938.8 766.4 718.5

(BH)max/kJ�m�3 270.6 262.6 230.8 214.8 175.1

Tc/ 501 475 478 465 455

X-ray diffraction showed that there exists the Lavas phase in the whole substi-

tutions of Ga. The phase is soft magnetic and forms nucleation points in the de-magnetization process. Consequently, the increase in coercivity is limited. The lattice parameters are listed in Table 4.5 for the tetragonal phase Nd2(Fe, Co, Ga)14B, the Laves phase and the Nd-rich phase.

Table 4.5 Lattice constants of the matrix phase, Nd-rich phase and the Laves phase in

alloy Nd16Co16Fe61-xGaxB7

Nd2(Fe, Co, Ga)14B Nd-rich phase Laves phase Content of gallium, x a c c/a a a

0 0.8762 1.213 1.384 0.510 0.740

1 0.8766 1.216 1.387 0.510 0.740

2 0.8772 1.218 1.389 0.510 0.740

4 0.8766 1.220 1.392 0.510 0.740

7 0.8757 1.222 1.396 0.510 0.740

From Table 4.5, partial substitutions of Ga for Fe in the tetragonal phase

Nd2(Fe, Co)14B have considerable influences on the lattice parameters. With Ga substitutions, the parameter a increases first, reaches to the maximum at the sub-stitution x=2, and then decreases, while the parameter c and the value of c/a in-crease linearly. The characteristics are very similar to the results of X-ray diffrac-tion of Al substituted NdFeCoB alloys (Xu, Ping, Li, et al, 1986). The Ga concen-trations have a little influence on lattice parameters of Nd-rich and Lavas phases. The lattice parameters of the B-rich phase are a=0.714 nm, cFe=0.391 nm and cNd=0.352 nm.

Based on the fitted areas of Mössbauer subspectra Si, the occupancies of Ga at-oms at the six sites in the tetragonal phase were calculated and the results were given in Fig. 4.5. The dotted line in the figure represents a random occupancy of the Ga atoms at the sites. The calculated occupancies of Ga atoms at e and c sites are negative numbers with almost zero, which indicate that the Ga atoms do not occupy the two sites. The occupancies are above the dotted line for the j2, k2 and j1


Fig.4.5 Occupation probability of Ga atoms via Ga content x

sites, and on the dotted line for the k1 site. Therefore, the Ga atoms preferentially occupy the j2, k2 and j1 sites. At the substitution x 2, the occupancy at the j2 site rapidly increases, as compared to the occupancies at the k2 and j1 sites. As the Ga atoms are non-magnetic, the exchange interaction between Fe/Co and Ga is equal to zero. As a result, the occupancy of Ga atoms at the j2 site leads to a significant decrease in the Curie temperature. As x>2, the increased ratio of occupancy at the j2 site reduces about 60 %, as compared to the case of x 2. In addition, with in-creasing the occupancies of Ga at k2 and j1, the magnitude of negative interactions between Fe atoms for j1-j1 and j1-k2 sites is decreased. The two factors lead to smaller decrease in Curie temperature for x>2 than for x 2. From experiments, the Curie temperature only reduces by 23 , as x varies from 2 to 7. On the other hand, Ga mainly occupies the j2 site, and thus decreasing easy planar anisotropy. Consequently, the coercivity increases with Ga substitutions.

The rare earth permanence-magnet alloys with six elements are prepared by the substitution of a small amount of Ga for Fe and Co. The effect of Co is to increase the Curie temperature and to decrease the irreversible temperature coefficient, and the effect of Ga or Al is to increase the coercivity. For Nd13.0Dy0.3Fe80.27Al0.2Ga0.08Cu0.05B6.1 permanence-magnet alloy sintering at the temperature of more than 1100 followed by ageing at 500-600 , the magnetic properties achieve: Br=1.44T, mHc=1,048kA/m (13.10kOe) and (BH)max=408 kJ/m3 (51.0 MGs�Oe). Such large energy product has its origin in the amount of main phase Nd2Fe14B as high as possible; in the alloy, the Nd2Fe14B accounts for 97.1%, while non-magnetic phases only for 2.32%. Therefore, Br=1.44 T and (BH)max>50 MGs�Oe are obtained. At the same time, the additions of Ga, Al, Dy and Nb are helpful for the increase in coercivity. The above formula and techno-


logical process can be also modified to obtain the permanence-magnet alloys with high coercivity. For example, the magnet possesses the parameters: mHc=2,400 kA/m (30 kOe) Br=1.04 T (10.4 kGs), and (BH)max=208.2 kJ/m3 (26.2 MGs�Oe)

(Pan, Zhao, Ma, 1988; Pan, Ma, Ping, et al, 1991; Pan, Ping, Liu, 2003). In order to increase the coercivity, the magnet energy product has to be sacrificed in terms of the current technology; it is right, as reversed. The company has reported on a rare earth permanence-magnet alloy with (BH)max=446.24 kJ/m3 (55.78 MGs�Oe), Br=1.514 T (15.14 kGs), and mHc=694.4 kA/m (8.6 kOe).

The substitution of a small amount of Ga for Fe not only increases the coerciv-ity, but also decreases the irreversible loss of magnetic flux. The combined substi-tutions of Ga and Nb decrease the irreversible loss to hirr< 5%. The loss will be larger than 40 %, increased by eight times, for the corresponding alloy without Ga and Nb. After the ageing of 260 , the operating temperature of the alloy can be elevated to 200 ; it increases 120 , as compared to Nd15Fe77B8 (Pan, Ma, Ping, et al, 1991).

Some of Ga enters the main phase Nd2Fe14B, and others are located in the grain boundary. Fig. 4.6 shows the TEM images of atom lattice for Nd15Co16Fe61-xGaxB7 with x=2. In the alloy, the homogeneous layer dislocations were observed, as shown in the TEM image of Fig. 4.7. The TEM image of the alloy with x=1 is shown in Fig. 4.8, from which, the boundary between the grains becomes clear after addition of Ga. From Fig. 4.9, it is observed that the grain boundary is greatly modified, as compared to that of Nd15Fe78B7. When a reversible magnetic-field exists, a strong pinning on domain wall happens and the wall is firmly pinned at the grain boundary. As a result, the coercivity is increased. A granular crystalline is densely distributed at an intersection between the grain boundaries, as shown in Fig. 4.9. Fig. 4.10 is an enlarged image of TEM. A strip-shaped boundary with gathering many small spheres is observed. After analysis, the

Fig.4.6 Electron micrograph of atom crystal lattice of Nd15Co16Fe60Ga2B7 alloy

(The electron micrograph of the complete atom crystal lattice is observed)


boundary consists of B-rich and Nd-rich phases. From Fig.4.9 and 4.10, we can see that there is diagonal contrast in interior of the matrix phase and closes to crystal boundary.

Fig.4.7 Electron micrograph of Nd15Co16Fe60Ga2B7 alloy

(Homogeneous stacking fault can be seen)

Fig.4.8 Electron micrograph of Nd15Co16Fe61GaB7 alloy

(Clear and complete and crystal grain boundary favorable to coercivity can be seen)



(Clear and orderly grain boundary of Nd2Fe14B without Ga-rich phase (Ga2Nd) is observed which is favorable for coercivity)


(Microstructure of Nd15Fe77B8 changes with substitution of Ga and Co which changes the triangular grain boundary formed by Nd2Fe14B and Nd-rich phase. The substitution of Ga and Co raises Curie

temperature from 312 to 450-550 and does not decrease coercivity obviously)

Ga and Co substituted NdFeB alloys, there are, in general, three phases, namely the Nd-rich, the B-rich, and the tetragonal phases. The eutectic temperature of these phases is significantly different from the melting temperature of Fe and Ga. When the alloys are quenched from high temperature to room temperature, the Nd-rich and B-rich phases are inhomogenously distributed; some is the straight grain boundary, as shown in Fig. 4.8. Due to deep eutectic, the eutectic composi-tions are aggregated in amorphously deeply-eutectic states at the intersection of boundaries. The ageing in 580-630 for 2-5 h leads to a homogenous distribu-tion of the boundary compositions, and accelerates the process from deep eutectic compositions to crystallization. Consequently, particles with the B-rich and Nd-


rich phase are densely distributed along the grain boundary. The sample after ag-ing has large coercivity if the optimum aging temperature is selected.

There are several Nd-rich phase with the compositions as follows, seeing Table 4.6 (Zhou, Dong, 1999).

Table 4.6 Ga-rich phase in NdFeCoGaB

Lattice constant Ga-rich phase Crystal system

a b c

Ga2Nd Hexagonal 0. 43 0.43

Ga3Fe11Nd6 Tetragonal 0. 81 2.3

GaNd Orthorhombic 0. 44 1.13 0.42

4.2.5 Conclusions

1. NdFe(M)B permanence-magnet alloys with high Curie temperature, high coer-civity and low temperature coefficient have been developed using powder metal-lurgy method. The intrinsic coercivity of 1130 kA/m, the maximum magnetic-energy product of 262 kJ/m3, the reversible temperature coefficient Br of �0.04%/ , and the Curie temperature of 450-550 have been achieved.

2. The partial substitution of Al and Co for Fe can increase the Curie tempera-ture from 312 to 450-500 . The substitution of 4%-10% Co for Fe increases the Curie temperature to 400 . For the substitutions of 16% Co and of 1%-2% Al, the Curie temperature is larger than 480 , and the coercivity is 988 kA/m and the maximum energy product is 239 kJ/m3.

3. Mössbauer spectra have shown that the Ga atoms mainly occupy the j2 site, then the j1, k2 and k1 sites, and excludes the e and c sites. The Al atoms mainly occupy the j2 and then the k1 site. The occupation of Ga atoms at j1 and k2 sites leads to a decrease in the negative interaction of the j1-j1 and j1-k2 sites. On the other hand, the Al atoms do not preferentially occupy the two sites, and the nega-tive interaction of the j1-j1 and j1-k2 sites is retained. Consequently, the Curie tem-perature decreases greatly for the Al substituted alloys than for the Ga substituted alloys.

4. All of the Ga substituted alloys, Nd16Co16Fe61-xGaxB7 with x=0, 1, 2, 4 and 7, contain the soft-magnetic Laves phase, which limits further increase in the coer-civity. For the tetragonal phase, the lattice parameter a increases first, to reaches the maximum at x=2 and then decreases; the parameters c and c/a monotonously decrease with Ga substitution. In addition, the coercivity has a peak at the Ga sub-stitution x=2.


4.3 The Studies of Main Phase Nd2Fe14B and Nd2(Fe,Co)14B in NdFeB Permanent-magnet Alloys *

NdFeB permanent-magnet alloys mainly consist of three phases, namely the tetragonal phase Nd2Fe14B, Nd-rich and B-rich phase. The Nd2Fe14B phase be-longs to the space group P42/mnm, which has been studied by Herbst, et al. The structure and properties of the B-rich phase have been studied by Givord, et al. The coercivity of NdFeB alloys has its origin in the anisotropy of Nd2Fe14B. Therefore, the studies of the Nd2Fe14B phase have been attracted much attention. The authors have studied the effect of Co substitution on Curie temperature for Nd2(Fe, Co)14B using TEM with super-high voltage of 1,000kV and Mössbauer spectroscopy (Pan, 1987; Filder, Huo, 1985; Pan, 1986; Ping, Li, Ma, Pan, et al, 1986; Ma, Jiang, Xu, 1999; Sagawa, Fujimura, fogawa, et al, 1984).

4.3.1 The preparation of samples and experimental methods

Samples of Nd15Fe85-xBx with x=0, 3, 5, 7 and 11 were prepared by arc melting under an Ar atmosphere. The starting materials were 99.5% Nd, 99.5% B, 99.9% Fe in purity and FeB alloy with B of 14.52 %. The ingots were remelted three times in order to achieve homogeneity. Then, the ingots were crashed and ball-milled to powders with about 3-5 m. The powders were shaped in a magnetic field with 1.5 T. Finally, the samples were sintered at 1090 and annealed at 600 for 1 h. The prepared sample was cut into thin plate along the c axis, which was thinned to 0.025 mm by the mechanical method and followed by the ion thinned method to make the thin film for TEM experiment. TEM with super-high voltage was performed using JEM-1000 kV. The powder samples were used in Mössbauer spectrum experiment.

4.3.2 SEM analysis

Fig. 4.11 is the SEM image of (Nd0.9Dy0.1)16Fe75Nb2B7. The dark range is the tetragonal phase Nd2Fe14B and the white range is the Nd-rich phase. SEM with EDAX showed that the relative concentrations of Fe and Nd(Dy) are 66.62%(wt.) and 33.38 %(wt.), respectively.

Mössbauer analysis and part of words were completed by Professors Ruzhang Ma, Zhengwen

Li and Jueyun Ping.


Fig.4.11 Electron micrograph of (Nd0.9Dy0.1)16Fe75Nb2B7 alloy at room temperature (Sintering 1090 ×1.5h, annealing 600 ×1h, the precipitation of Fe2Nb (black, different shape)

observed in matrix phase (Szafra�ska-Miller B, P usa D, Wys ocki J J, et al, 1987))

4.3.3 The formation of Nd2Fe14B

In order to study the formation of Nd2Fe14B, the samples Nd15Fe85-xBx with x=0, 3, 5, 7 and 11 have been prepared. X-ray diffraction showed that there are only �-Fe and Nd2Fe17 phases without B, and the new diffraction lines of Nd2Fe14B appear at B concentration of 3% (at.). As B>5% (at.), �-Fe and Nd2Fe17 phases com-pletely disappear and there are main phase Nd2Fe14B, the Nd-rich and B-rich phases. SEM with EDAX showed that the mass fractions of Fe and Nd are 65.62 % and 34.38 %, respectively, which are consistent with the fractions of Fe and Nd in Nd2Fe14B. In conclusion, the B concentration is the decided factor in the forma-tion of the tetragonal phase Nd2Fe14B. The ferromagnetic coupling of Fe and Nd is along the c axis in the tetragonal structure. The strong uniaxial anisotropy is the main cause of high coercivity for NdFeB alloys. The anisotropy has its origin in the energy splitting of rare earth ions due to the crystal field.

4.3.4 Mössbauer spectra at room temperature

The Mössbauer spectra of Nd15Fe85-xBx with x=0, 3, 5, 7 and 11 indicated that at x=0, there are only the Mössbauer components of �-Fe and Nd2Fe17 phases. At x=3, the Nd2Fe14B appears, and the same time, �-Fe and Nd2Fe17 still exist, as pointed by their 1st and 6th characteristic lines a and a� and by b and b�, respec-tively. At x=5, the spectral lines of Nd2Fe17 almost vanish and at x=7, the lines of �-Fe completely disappear. There exist the sextet of the main Nd2Fe14B, and the paramagnetic doublet of Nd-rich phase.

Table 4.7 lists alignment sequence of hyperfine field, which can determine sites adscription of ferromagnetic sub-spectra.


Table 4.7 Fe crystal sites and neighbor circumstance of Nd2Fe14B tetragonal phase Adjacent atom number Site of

crystal Wykolff notation

Fe atom number in unit cell Fe Nd B

Order of hyperfine field

Fe1 e 4 9 2 2 5 Fe2 c 4 8.8 4 0 5 Fe3 �1 8 9 3 0 3 Fe4 �2 8 12 2 0 1 Fe5 k1 16 9 2 1 4 Fe6 k2 16 10 2 0 2

4.3.5 Composition analysis and the studies of Mössbauer spectra for Nd2(Fe, Co)14B

The composition analysis of electron probe has been performed for Nd2(Fe1-x, Cox)14B. At x=0, the compositions of Nd and Fe are 19.4% (at.) and 84.6% (at.), respectively. The Curie temperature is 312 ; at x=0.3, the compositions of Nd, Fe and Co are 15.4% (at.), 62.5% (at.) and 22.1% (at.), respectively. Then the Curie temperature is heightened to 550

4.3.6 In situ and dynamic observation of TEM on Nd2Fe14B and Nd2(Fe, Co)14B

The in situ and dynamic observation was conducted on filmy sample under elec-tric microscope from room temperature to 130 . At this temperature range there is no precipitate in Nd2Fe14B, as shown in Fig. 4.12. At 260 precipitates appeared

Fig.4.12 Electron micrograph of Nd15Fe78B7 permanent magnet alloy at room

temperature and electron diffraction of B-rich phase (right above) (Observed Nd-rich phase (A), matrix phase Nd2Fe14B (B) and B-rich phase (C) and

high density of defect in B-rich phase)


in the Nd2Fe14B and at 312 highly dispersed precipitate phase was observed clearly in the all vision and diffraction ring of polycrystal appeared in diffraction pattern of Nd2Fe14B. At 500 more diffraction ring of polycrystal appeared, which indicated that polycrystal turned up in Nd2Fe14B but the crystal boundary is very clear, as shown in Fig. 4.13.

At 400 there are tiny precipitates stated to come forth in Nd2Fe14B and at 500 the precipitates grew up like needles and be perpendicular to each other, such as shown in Fig.4.14.

Fig.4.13 Electron micrograph of Nd15Fe78B7 permanent magnet alloy at 500

(Observed clear boundary between Nd2Fe14B grains at 500 , B-rich phase, heating from room tem-perature to 500 , variation of its defect density (C), and obvious variation in microstructure

of Nd-rich phase (A), a few points of Nd2O3 (D) appears in matrix phase Nd2Fe14B (B))

Fig.4.14 Electron micrograph of Nd15Fe70Co8B7 permanent magnet alloy at 500

(Observed precipitation point growing up into needle-like phase in Nd2(Fe,Co)14B matrix phase, lot of needle-like phase perpendicular to c axis in bottom crystal plane appears; high density of crystal de-

fect, twin crystal, dislocation, stacking defect etc. as sub-structure appear with phase transition)

At 600 a lot of precipitates appeared in Nd2Fe14B, which was Nd2O3 by elec-tronic diffraction, as shown in Fig. 4.15 to Fig. 4.17.


Fig.4.15 Electron micrograph of Nd15Fe78B7 permanent magnet alloy at 600

(Observed high temperature phase transition at 600 and 500 with much higher heat activation temperature: lot of Nd2O3 (A) appears in Nd2Fe14B, B-rich phase in grain boundary increases (B), (E),

(F) give the results of electron diffraction of matrix phase)

Fig.4.16 Electron diffraction pattern of precipitation phase in Nd2Fe14B of Nd15Fe78B7

permanent magnet alloy at 600

Fig.4.17 Electron diffraction pattern of precipitates in the matrix phase of Nd2Fe14B in

Nd15Fe78B7 permanent magnetic alloy at 600


4.3.7 Conclusions

The conclusion was derived through aforementioned investigation: 1. Tetragonal crystal of Nd2Fe14B compound is the only hard magnetic phase.

Measurement of magnetism showed Br = 1.61 T; by observation of TEM crystal-line structure is intact with a = 0.88 nm, c = 1.22 nm. To enhance magnetic per-formance it is necessary to increase volume fraction of the matrix phase Nd2Fe14B as much as possible.

2. Increase of Boron content results in increase of phase Nd2Fe14B. When Bo-ron content x is increased to 5 the Mössbauer spectra for Nd2Fe17 phase is almost disappeared, and paramagnetic spectrum corresponding to B-rich phase appears synchronously.

3. Increase of cobalt content leads to enhancement of Curie temperature. Area of sub-spectra in the Mössbauer spectra reflects comparative amount of distribu-tion of iron atoms at crystal sites, that indirectly expresses that cobalt atom has different preferential in distribution at crystal sub-sites. Analysis indicates that cobalt atoms have priority to occupy sites j2 and k2.

4. Dynamic observation on filmy specimen under 1000 kV HVEM discovered that precipitate starts to appear in Nd2Fe14B phase at 260 , the highly dispersed precipitate phase can be found clearly in all vision at 312 and polycrystalline diffraction ring appears in electronic diffractive ground pattern. At 500 more polycrystalline diffraction ring appears in diffractive ground pattern of Nd2Fe14B.

5. For Nd2(Fe, Co)14B phase the precipitate phase appears at 400 , the precipitation temperature is about 100 higher than that of Nd2Fe14B compound.

4.4 Studies on B-rich Phase in NdFeB Alloy *

People have made many studies since the birth of high performance NdFeB per-manent magnet. This permanent magnetic alloy has the Nd2Fe14B compound as the base phase and contains Nd-rich phase, B-rich phase and sometimes a small amount of �–Fe. The studies on Nd2Fe14B and RE2Fe14B compounds have been relatively in-depth. The relationship between B-rich phase and magnetic hardening is attracting people’s attention. D. Givord, et al elaborated the B-rich crystal struc-ture under a long periods and ultra-structure model of Nd5Fe18B18 (Givord, Moreau, Tenaud, 1985). Regarding the sensitivity of the coercivity with micro-structure author observed the microstructure of NdFeB alloy, especially variety of the microstructure of B-rich phase with change of temperature, and discussed coercivity mechanism of NdFeB permanent magnetic alloy by combining other

Cooperators of this test include: Professor Ruzhang Ma and Doctor Zhibo Zhao, University of Science and Technology Beijing.


results.

4.4.1 Preparation process and experimental method

The alloy was melted in a non-consumable vacuum arc furnace with the constitu-ents of high purity material of metallic neodymium, boron powder, and metallic iron. And the ingot was pulverized into powder of 3m under protection of argon atmosphere and was formed under 1.5 T magnetic fields. Then it was sintered at 1100 and proceeded aging at 600 . The above sample was sliced into lamellas of 0.25 mm in a direction vertical to c axis and than was thinned mechanically to 0.025 mm. Afterwards the observable filmy specimen was obtained by ion thin-ning. The observation was carried out in a JEM-1000 kV HVEM equipped with a heating apparatus. The specimen for X-ray diffraction analysis and Mössbauer experiment was all in fine powder with a size passing through 325 mesh sieve (0.043 mm).

4.4.2 The in situ and dynamic observation of Nd1.11Fe4B4 by TEM

4.4.2.1 Pattern at room temperature

Appearance of B-rich phase at room temperature is shown in micrograph of Fig. 4.18.

Fig.4.18 Electron micrograph of Nd15Fe78B7 permanent magnet alloy at room temperature

(Observed B-rich phase with high density dislocation, Nd-rich phase in triangle grain and Nd2Fe14B phase)

A B-rich phase; B Matrix phase; C Nd-rich phase It can be seen from the micrograph that B-rich phase locates in the triangle

crystal boundary of the base phase, where Nd-rich phase also exists. Some part of B-rich phase directly contact with the base phase, and some where is an immingle transition area of B-rich phase and Nd-rich phase. High dense defects exist in the interior of B-rich phase, which represents as striae, alternated with bright and dark,


paralleled in direction vertical to c axis. The electronic diffraction pattern of B-rich phase at room temperature is shown

in the top right part in the micrograph of Fig. 4.18. The tetragonal system was demarcated, and as the result the Fe-B substructure was obtained with crystal lat-tice parameters as: a = 0.712 nm, cFe-B = 0.389nm. This result basically accords with that by X-ray diffraction (Givord, Moreau, Tenaud, 1985). Meanwhile a phenomenon can be observed that reciprocal pole was elongated.

4.4.2.2 Microstructure at high temperature

The Nd15Fe78B7 specimen was heated in the HVEM and the whole process was videotaped. And the observation was carried out for variation of microstructure of the permanent magnetic alloy. It was found that there almost had no change at 25-100 (Fig. 4.19). But at 140 the defect was observed in interior of B-rich phase, and especially at 280 that the precipitates appeared in B-rich phase. The further analysis was not carried out for the precipitates because the precipitate particles ware too small. Electronic diffraction indicated that the microstructure also vitiated remarkably at 322 , but its appearance recovered again to be similar to that at room temperature at 500 , as shown in the micrograph of Fig. 4.20.

Fig.4.19 Electron micrograph of Nd15Fe78B7 at 322

(Where new phase points of Nd2O3 appeared in microstructure of B-rich phase) B Matrix phase; C B-rich phase; K Electronic diffraction pattern of B-rich phase at 322 ;

J Electronic diffraction pattern of B-rich phase at 25

Fig.4.20 Electron micrograph of Nd15Fe78B7 alloy at 500

(Observed variation of the microstructure of B-rich phase and the growth of Nd2O3 phase and appearance of twin crystal)

A B-rich phase; B Matrix phase


4.4.3 Study on Nd1+�Fe4B4 by X-Ray diffraction and Mössbauer effect

In order to study structure and specialty of B-rich phase of Nd1+�Fe4B4 the alloy was confected as per composition of Nd1.1Fe4B4 and annealed at 1,100 for 25 days. It discovered by X-ray diffraction that the powder diffraction pattern was composed by 2 set of lines. It was revealed by strength computing that the rela-tively sharp line came from the contribution of Fe-B substructure; and that the relatively dispersed one was correspondingly the neodymium substructure. The result can be obtained by self edited computer indexing program that: a = 0.7116 nm, cFe-B = 0.3894 nm, cNd = 0.3523nm.

Mössbauer spectrum of Nd1.1Fe4B4 is fitted by Lorentz line type using the least square method, the detailed result analysis can refers to reference (Pan, Ma, Ping, et al, 1991). The Mössbauer spectrum of Nd1+�Fe4B4 compound is two fourth grade split peaks, as shown in Fig. 4.21. This indicated that the Nd1+�Fe4B4, i.e., the B-rich phase, appeared as paramagnetic, ultra-subtle parameter is isomerous shift I. S. = 0.03 mm/s and the fourth grade split is Q. S. = 0.56 mm/s. Split of ferromagnet still was not found from the Mössbauer spectrum measured at temperature of liquid ni-trogen. That indicated that the atomic order temperature for ferromagnet is al least below 77 K. And the result of relevant magnetic measurement shows that the Curie temperature is about 14 K for B-rich phase (Givord, Moreau, Tenaud, 1985; Pan, Ma, Ping, et al, 1991; Bezinge, Braun, Muller, et al, 1985).

Fig.4.21 Mössbauer spectrum of B-rich phase (Nd1.1Fe4B4) in

casted Nd15Fe78B7 alloy at room temperature

4.4.4 Analysis on Nd1.1Fe4B4 phase

Based on elaborated crystalline structure, D. Givord, et al indicated that chemical


formula of the B-rich phase could be formulated as Nd1+�Fe4B4, where �=cFe-B/cNd -1 which can be derived from data of X-ray diffraction. It is noticed that there is not a simply multiple relation between cFe- B and cNd, thus this compound may belong to an one dimension composition varying type incommensurate structure. The Nd1+�Fe4B4 compound may also be described as “infinitely adaptive struc-ture”(Anderson, 1973). The phenomenon that double periods subsisting in the same crystal structure had been discovered in MnSi2-x and the interrelated Nowotny phase (Zwilling, Nowotny, 1971). In order to project the speciality of double periods that is normally called as vernier structure or chimney - ladder structure.

It can be observed from X-ray diffraction pattern that the lines in sublattice of neodymium appeared widening obviously. This apparently may not be attributed to interior stress or over fined crystal particle because there were only a few lines being widened but not all diffraction lines of Nd1+�Fe4B4 compound being wid-ened. We consider that this may be a result of interaction of inter-sublattice. Modulation of Fe-B sublattice to neodymium sublattice made some direction of neodymium sublattice lose periodicity, which resulted in unequal space among crystalline planes and resulted in widened lines. In fact the Fe-B sublattice is also modulated by neodymium sublattice. A. Bezinge, etc. observed the kink modula-tion of Fe-B sublattice (Bezinge, Braun, Muller, et al, 1985). That may be related to interaction of sublattice structures.

At the same time we observed that the interior crystals of Nd2Fe14B base phase appears relatively high integrity, where have no defect, nor precipitate. Therefore, no matter nucleation center of reverse domain or pinning point all unable to exist in crystals of Nd2Fe14B. The magnetic hardening related to microstructure in crys-talline boundary. The B-rich phase is paramagnetic at room temperature so that it can act as function of inclusion pinning on domain wall and heighten the coerciv-ity. Table 4.8 indicates relationship between magnetic property and boron content in the alloy. It can be seen from the table that the property reaches optimum when the boron content is 7% (at.). The superfluous boron would cause decline in value of Br and (BH)max, but the coercivity is heightened. Apparently, when boron con-tent exceeding 7% (at.) and reaching 11% (at.) the saturated magnetization strength will decline and so that depress the remanence because iron and neodym-ium are reduced comparatively in proportion of the alloy. After formation of suf-ficient Nd2Fe14B and B-rich phase the remnant will mainly form B-rich phase. The quantitative analysis on the phase proceeded in our early work indicated that boron content in the B-rich phase of the alloy will increase from 4.6% to 10.4%, more than doubled, when boron content is increased from 7% (at.) to 11% (at.). Thus the coercivity is heightened 130.35kA/m accordingly. We deemed that in-crease of B-rich phase enhanced the pinning ability of crystalline boundary for anti-magnetization process and so that resulted in enhancement of coercivity. This


is because energy of domain wall (�) is proportional to volume V of additive phase. Therefore, it can be explained by formula 4.3 for enhancement of Hc by increasing B-rich phase:

cs max

12

HM x

-�� . � �� (4.3)

Table 4.8 Relationship of magnetic performance of Nd15Fe85-xBx alloy and Boron content

Boron content, x Br/T jHc/kA�m�1 (BH)max/kJ�m�3

0 0.05 3.95

3 0.21 51.35

5 0.81 434.50 115.42

7 1.35 758.40 330.34

11 1.07 888.75 223.68

4.4.5 Relationship between B-rich phase and coercivity

The high dense defects in B-rich phase can be observed obviously, and that the phenomenon of elongated reciprocal pole in electronic diffraction spectrum may be related to these defects in B-rich phase. Existence of high dense defects can enhance pinning ability on domain wall so that further heightens the coercivity. It can also be seen the Nd-rich phase and filmy belt besides the B-rich phase in NdFeB. Thus mechanism of coercive force is a complicated problem and the rela-tive studies are under way (Pan, Zhao, Ma, 1988; Pan, Ping, Liu, et al, 2003; Pan, Liu, Luo, 1990; Pan, Ma, Li, 1993).

K. J. Strnat, et al (Strnat, Li, 1985) discovered that the coercivity of NdFeB permanent magnetic alloy almost disappeared at 265 . This change is very pos-sibly related to variation of microstructure with change of temperature, besides the contribution by heat activation. For different pinning instance the dependency relationship of coercivity on temperature is different. Our observation indicated that the B-rich phase normally locates at trigonal crystalline boundary and dis-tributes very dispersed. The pinning points are of low density thus they may be-come weak pinning, then the relation between Hc and temperature may be repre-sented as:

2 /3

1/ 2c o

75( / ) 1 KTH HAbr

� ��

(4.4)

This formula is consistent with experimental date by Handjipanayis, et al (Had-jipanayis, hawless, Dickerson, 1985).


4.4.6 Conclusions

It is concluded as below based on above studies: 1. The chemical formula for B-rich phase is Nd1+�Fe4B4, it possesses a very

special crystal structure because there is not a simply multiple relation between cFe-B and cNd. Thus the B-rich phase may be an one dimension component varia-tion type commensurate structure and this compound is similar to Nowotny phase of MnSi2-x, etc., or a chimney-ladder structure.

2. The B-rich phase becomes paramagnetic phase at room temperature. They locate in a triangular crystal boundary and can become pinning point on the do-main wall of anti-magnetization nucleus that avails enhancement of coercivity.

3. There are high dense defects existent in B-rich phase. These defects are per-pendicular to c axis, and that may enhance pinning ability of B-rich phase on do-main wall.

4. The coercivity is related to microstructure but the microstructure varies with change of temperature. Therefore, the further observation of relationship between microstructure and temperature will be helpful for understanding of coercivity mechanism.

4.5 Influence of Boron Content in NdFeB on Nd2Fe14B Phase and Magnetic Property *

�� investigated Ce-Fe-B phase diagram in 1972. Afterwards �� worked out Y-Co-B, Ce-Co-B and Sm-Co-B ternary phase diagrams one after the other, in 1976 he worked out Co-La-B phase diagram, and in 1979 worked out Fe-Nd-B, Fe-Sm-B and Fe-Gd-B phase diagrams.

In 1980 Croat, Koon, Beckor, and C. C. Hadjipanayis investigated Fe-Pr and Fe-Nd series microcrystal permanent magnets. It found the coercivity of 7.5 kOe (596.8 kA/m), (BH)max= 3-4 MGs�Oe (23.9-31.8 kJ/m3) in Nd0.4Fe0.6 and Pr0.4Fe0.6 Alloys prepared by J. J. Croat. He considered that the produce of coercivity is related to formation of metastable microcrystal structure. In 1981, N.C. Koon, C. M. William, B.N. Das reported that the coercivity reached 9 kOe (716.2 kA/m) after crystallizing amorphous Tb0.05La0.05(Fe0.82B0.18)0.9 at 930K. In 1983, C.C. Hadji-panayis and R.C. Hazelton of Kansas University, USA worked out Pr16Fe78B5Si3 with the intrinsic coercivity of mHc = 15 kOe (1193.7 kA/m), (BH)max = 13 MGs�Oe (103.4 kJ/m3), and the specimen only containing boron (Pr16Fe78B6) also

Cooperators of this study are Professors Ruzhang Ma, Zhengwen Li, Jueyun Ping and doctor-course students Bo Zhao and Baoguo Li.


reached the similar level. They pointed out that the crystallized specimen became preferable permanent magnetic phase because Fe20RE3B6 with the square phase has a great of magnetic anisotropy. In June of 1983, Sumitomo Special Metal Corporation announced that they successfully developed the Nd15Fe17B5 perma-nent magnet with a high magnetic energy product.

This section is to report that study the function on formation of phase texture by addition of boron and study the relation between variation of alloy microstruc-ture and the coercivity by using electronic microscope, 1000 kV HVEM, X-ray microscope analysis instrument cooperating with magnetic measurement instru-ment.


NdFeB alloy for experiment was made by using normal powder metallurgy method. At first the high purity neodymium, high purity iron and high purity bo-ron was melted in non-consumable arc furnace. In the melting process the furnace was pumped to for vacuum in advance and filled argon latter. In order to ensure component homogenization the alloy was melted thrice continually. The melted alloy was cracked roughly at protection of argon atmosphere and pulverized in a ball grinder into about 3m. The powder was formed under 1.5 kGs magnetic field and sintered at 1100 . After aging at 600 it was cooled to room tempera-ture as sample for electronic microscope.

The above sample was sliced into lamellae of 0.25 mm in direction perpendicu-lar to c axis, and thinned to 0.025 mm mechanically, and then the lamella were electrolyzed to open a hole basically in an electrolyte of 20% perchloric acid and 80% glacial acetic acid and finally be thinned and cleaned on ionic thinning appa-ratus. Then the specimen was ready for observation.

Observation of the filmy specimen was carried out in JEM-1000 HVEM. The operating voltage was 1000 kV, output voltage was 185V, electric current was 6.6 A, vacuum was 0.7×10-7 Torr (9.3×10-6 Pa) and the ion beam was 10 A. The specimen was inserted at first into the side inserting type heating dais of JEM-1000 in the observation and observed by electronic microscope in the condition of room temperature and heat condition.

4.5.2 Influence of boron content on alloy magnetic property and phase structure

This section investigated the relationship between the B content and the magnetic property. The addition of B (6%-8%(at.)) facilitates the formation of hard mag-netic phase resulting in high performance magnetic materials.


4.5.2.1 Variation of magnetism with change of boron content

It can be seen from Table 4.9 and Fig.4.22 that boron content has important influ-ence on property of magnet of NdFeB permanent magnetic alloy. It is seen from the condition of Table 4.9 that the magnetism will be optimal when boron content is 7% (at.), in fact preferable magnetic performance in all can obtained the in the boron content range of 6%-8% (at.). Addition of boron enable NdFeB alloy to form a square phase.

Table 4.9 Relationship between magnetic performance and B content

Boron content/%(at.) Br/Gs mHc/Oe (BH)max/MGs�Oe

0 500 50

3 4100 620

5 8000 5300 13.2

7 13500 9650 41.5

11 10700 11250 28.5

Fig.4.22 The influence of boron content of Nd-Fe-B on its magnetic characteristic

(Demagnetization curve) Using Mössbauer effect to Nd15Fe85-xBx (x=0, 4, 8, 11) we conclude: follows

boron (B) content increase, Nd2Fe14B magnetism look gradually increase in alloy. As x<4, there is no rich boron look apearance, and Nd2Fe14B look descend 24%, magnetic performance descends too. If x=11, Nd2Fe14B descend 5%. That does not affect intrinsic coercivity. Enhance Nd2Fe14B quantity as to enhance saturation magnetization intension, properly control boron (B) content and alloy ingredient, properly increasing Fe is to enhance Nd2Fe14B look effective means (Pan, Zhao, Li and Ma, 2011).


4.5.2.2 Influence of boron content on phase structure

In order to study the formation of strong single axis anisotropy the following work has been conducted:

(1) Confected specimen of Nd16BxFe84-x, while x = 0, 3, 5, 7, 11; (2) Carried our phase analysis by X-ray diffraction for above specimen (Fig. 4.23).

Fig.4.23 The result of X-ray diffraction of casted Nd15Fe85-xBx alloy

(a) x=0; (b) x=2; (c) x=8 (Where signs *, ** and / present the index of the crystal plane of Nd2Fe17, Nd-rich phase and metal Nd,

unlabelled is the crystal plane index of square phase Nd2Fe14B, the black upside down arrow indicates the diffraction line of B-rich phase)

The result of component analysis of micro area being carried out by electronic

probe is shown in Table 4.10 (the boron element here is unable to be analyzed by this instrument).

Table 4.10 Quantity analysis of phase composition in NdFeB alloy

Boron content/%(at.) Phase Fe/%(wt.) Nd/%(wt.)

Nd-rich phase 38.29 61.71

The matrix phase 41.73 58.27 0

�-Fe, Nd2Fe17 93.46 6.54


Continued Table 4.10 Boron content/%(at.) Phase Fe/%(wt.) Nd/%(wt.)



�-Fe, Nd2Fe17 88.90 11.10



�-Fe, Nd2Fe17 82.29 17.71



�-Fe, Nd2Fe17



�-Fe, Nd2Fe17



�-Fe, Nd2Fe17

60.88 39.12

62.47 37.53 Nd2Fe14B

69.41 30.59

Note: Data of Fe% (wt.) and Nd% (wt.) are the average of several measurements and are a compa-rable percent with all Fe and Nd (excluding boron content).

It can be seen from above experiment result that without boron the alloy was

composed of �-Fe, Nd2Fe17 base phase and Nd-rich phase. The composition of Nd-rich phase, �-Fe and Nd2Fe17 base phase was changed when boron content of 3% (at.) was added into the alloy. And that a new peak, being analyzed as Nd2Fe14B, appeared besides the original �-Fe and Nd2Fe17. While the alloy with a boron content of 5% (at.) the base phase composition change to be Fe 65.62%, Nd 34.38%, being consistent with Nd2Fe14B phase, besides the change in percentage of iron and neodymium. Thus increase of boron content promotes formation of quadrangle crystal system and strong magnetic phase of Nd2Fe14B. Also because the neodymium and ferromagnetic moment of ferromagnetic-coupling is along c axis all of crystalline structures are square and so that as an anisotropy structure conduced for the high coercivity. The main part of intrinsic anisotropy inside of Nd2Fe14B phase originated from the split of crystal field with 4f energy level of rare earth materials. Thus it can be seen from Fig. 4.22 that the coercivity height-ened rapidly because of increase of boron content, and neodymium element en-riched degree in crystal boundary increased largely from planar distribution phase of neodymium. When boron content was 7% the �-Fe and Nd2Fe17 would disap-


peared referring to X-ray diffraction spectrum, which was also verified by micro-area analysis experiment using electronic probe. There are no �-Fe and Nd2Fe17 in

Nd16B7Fe77 (i.e., 7% (at.) boron content) in Table 4.10, and that neodymium con-tent raised to 82.45% in Nd-rich phase and the neodymium enrichment becomes more obvious in crystal boundary. Then the main phase is the base phase, i.e., Nd2Fe14B. It can be seen from Fig.4.22 and Table 4.9 that this boron content en-ables the magnetic performance to reach the optimum status. J. F. Herbst, J. J. Croat, et al conducted studies on cells on crystal of Nd2Fe14B and found: there are four Nd2Fe14B and total 68 atoms in each crystal cell, among them iron atoms are 56, neodymium atoms are 2 and boron atoms are 4; and the boron atoms occupied triangle prism constituted of 3 nearest iron atoms in upside and downside of base plane. All moments of neodymium and iron paralleled with c axis of square crys-tal cells. The more perfectible configuration is the reason to improve Tc (for Nd2Fe14B is 627 K, for Nd2Fe17 is 330 K). That only praseodymium and neodym-ium can provide a higher product of magnetic energy. M. Sagawa believed that boron in NdFeB square phase acted enlarging atomic space of Fe-Fe and reduced effect of the nearest neighbor atomic number of iron, and that addition of boron was the reason to induce rising of Curie temperature. Therefore, controlling ap-propriate boron became a key subject to obtain excellent magnetic performance. It can be seen from Fig.4.22 and Table 4.9 that when boron content is higher than 7% (at.) i.e., increases to 11% (at.) Br will be declined and mHc will be enhanced somewhat, that indicated decline in the product of magnetic energy which indi-cates integrated magnetic performance. And that enriched degree of neodymium in boundary of prismatic crystals (casting status) was decreased referring to pla-nar distribution of neodymium by study using scanning microscope. Nd2Fe14B peak was weakened by referring to X-ray diffraction spectrum of Nd16Fe73B11 alloy and the phase composition in micro-area was changed, as shown in Table 4.10.

The saturated magnetization intensity of square phase was 15.7 kGs at room temperature (Bilonizhko, Kuzma, IZV, et al, 1974), and that the remanence Br can be estimated in expression as below as per the square phase in alloy and orienta-tion angle of magnetic moment.

r s ( / )cosB M V V 0/� (4.5)

where V � and V represent the volume of square phase and volume of the whole specimen, respectively; Ms represent saturation magnetic intensity. The rema-nence of alloy estimated for x = 3, 5, 7, and 11 by the Eq. 4.5 is well consistent with the result in actual measurement and the difference is within 7%.

4.5.3 Conclusions

Conclusion was derived through above studies:


1. Boron content in NdFeB permanent magnetic alloy has important influence on the magnetic performance of the alloy. Appropriate boron content may result in a good magnetic performance.

2. When B = 0% (at.) the NdFeB alloy is mainly constituted of Nd-rich phase �-Fe and Nd2Fe17. When B content increased from 0%-7% (at.) the �-Fe and Nd2Fe17 disappeared and Nd2Fe14B phase formed gradually. While B=11% (at.) the strength of base phase weakened but there still has �-Fe and Nd2Fe17. That resulted in degradation in magnetic induction intensity and magnetic performance (but mHc can be raised). Thus controlling boron content may obtain NdFeB per-manent magnetic alloys with different brand, different performance in practicality.

3. Addition of boron into NdFeB alloy may promote formation of strong mag-netic performance phase.

4. Observation of NdFeB alloy specimen using 1000 kV HVEM found a filmy belt between crystal boundaries of two hard magnetic square phases. This filmy belt was widened regularly in proportion with rising of temperature ranged from 25-600 .

4.6 High Curie Temperature NdFeCoGaB Permanent Magnetic Alloy *

To improve low Curie temperature, coercivity and bad thermal stability of ternary alloy Nd15Fe78Co8B7 using Co and Ga to replace part of iron to constitute varia-tion order and effect of element dynamic cross of quinary alloy.

4.6.1 Preparation process and experiment method

Sample was prepared as per Nd16(Fe1-xCox)77B7, where x= 0, 0.05, 0.1, 0.2, 0.3. Samples with Ga was prepared as per Nd16Co16Fe61-xGaxB7, where x = 0, 1, 2, 4, 7. Prepared material was melted in vacuum non-consumable are furnace with argon protection. Melted material was crushed roughly in crushing appliance, and pul-verized to 3.5 m, and then the powder was molded under magnetic field of 1.5T and pressure of 2T/cm2. Molded roughcast was sintered in furnace at 1100 for 1 hour and edged at 600 , and then quenched to room temperature. Samples of good sintered and edged were magnetized in strong pulse magnetic field. Mag-netic parameters were measured using CL-6 magnetic measurer and Curie tem-perature was measured by magnetic balance.

Condition for X-ray diffraction experiment: cupper target with wave length 0.15418nm. The Mössbauer spectrum was measured under iso-accelating the

* Cooperators of this work include: Professor Ruzhang Ma, Jueyun Ping and Professor Zheng-wen Li, University of Science and Technology Beijing.


Mössbauer spectrum. Experiment was carried out at room temperature, radiation source was 57Co (Rh) and used �-Fe for to measure velocity and took the central route as zero point of velocity. All Mössbauer experiment sample were in state of powder. All Mössbauer spectrum used a group of single-line spectrum, a group of dual-line spectrum and 6 groups of hexed-line spectrum to fit by the least square method, and six inequivalent iron crystal sites corresponding to Nd-rich phase and B-rich phase and tetragonal phase (Pan, 1986; Ping, Li, Ma, Pan, et al, 1986; Ma, et al, 1984).

4.6.2 Using cobalt to replace part of iron

Table 4.1 is the result of magnetism and Curie temperature measurement for sam-ple of Nd16(Fe1-xCox)77B7, with different cobalt contents. It can be seen from Table 4.1 that Curie temperature was heightened with increasing cobalt content, but the coercivity (jHc) and remanence (Br) declined with increasing cobalt, thus magnetic energy product (BH)max was also declined.

There are 6 inequivalent iron sites in Nd2Fe14B, Curie temperature is raised with increase of atomic distance of Fe-Fe. Among 6 sites j1 and j2 have most im-portant function on Curie temperature. At first j2 has the most neighboring atoms (12) and its distance with the neighboring atoms is the biggest (268×10-12 m) and so that it has the strongest exchange function. The next iron atomic distance be-tween j1 site and k2 and j1 sites are 2,936×10-13 m and 2,433×10-13 m, respectively, that are smaller than critical distance 245×10-12 m of plus minus exchange integral. When cobalt content x > 0.2 cobalt atom occupies k2 site preferentially (Ping, Li, Ma, Pan, et al, 1986), the minus exchange function is improved because exchange integral jCo-Co > jCo-Fe > jFe-Fe. Simultaneously, when x < 0.2 Co atom occupies j2 site preferentially, which will further increase function of plus exchange. The re-sult leads to increase of Curie temperature with increase of cobalt content. Site j2 has a strong plane anisotropy, thus the preferential replacement of cobalt atoms will enhance plane anisotropy, but hypolattice anisotropy of neodymium is bigger than that of 3d transit metals so that increase of cobalt content decreases single axis anisotropy of tetragonal phase Nd2(Fe1-x, Cox) 14B, but in whole range of re-placement by cobalt the tetragonal phase is still axial anisotropy. In addition, along with increase of cobalt magnetic soft phase Nd(Co, Fe)2 appeared and it may form nucleation point in process of reverse magnetization. Therefore, coer-civity of the alloy degrades repidly with increase cobalt content. In preparation of samples it is necessary to control condition of process to avoid appearance of Nd(Co, Fe)2 phase to enhance the coercivity.

4.6.3 Use Ga to replace part of iron in NdFeCoB alloy

Table 4.4 is measuring result of magnetism and Curie temperature of Nd16Co16


Fe61-xGaxB7. It can be seen that Curie temperature declines fast with increase of x when x < 2, and Curie temperature declines slowly when x > 2. Author used Mössbauer method to study this phenomenon.

As described in section 4.6.2 Nd2Fe14B tetragonal phase had 6 inequivalent iron sites, when Co and Ga are added they will occupy some of the sites and their oc-cupation probability is:

6

Fe Fe Fe Co1

Co Fe

Ga Fe Co

14[ /( )] /

1 (Compiling Group, 1979)

ii i i

ii i i

i i i

P C C C S Sn

P P

P P P

�

� � ��

� � !

� � � �#

% (4.6)

where PiFe, Pi

Co and PiGa present probability to occupy the site No. i respectively;

CFe and CCo present iron content and cobalt content in the alloy, respectively; Si is the area parameter in the Mössbauer fitting; value ni (i = 1-6, represent sites j2 , k2, j1, k1, e and c) is decided by formula below:

2 2 1 1: : : : : : : : :2 4 2 : 4 1 1j k j k e cn n n n n n � (4.7)

where �i is the atomic rate of iron to cobalt when gallium content is 0 and is a constant relative to a certain site.

By Eq.4.6 and Eq.4.7 and the Mössbauer fitting result that the curve of gallium atom occupation probability changes with change in gallium content x is obtained as shown in Fig. 4.24. Where the broken line represents random distribution status

Fig.4.24 Relation of Ga atom’s probability of occupying of crystal lattice place in

tetragonal phase vs Ga content x in Nd16Co16Fe60GaB7


of gallium in each site. It was obtained by calculation that the occupation prob-ability of sites e and c is minus, almost be zero that indicated gallium does not occupy that two sites.

It is obviously that when x 2 gallium mainly occupies site j2, when x > 2 gal-lium has a certain probability to occupy each site of k2, j1 and k1 yet. The very interesting thing is that at position x = 2 occupation probability curve of gallium in site j2 presents slope variation. Exchange function of j2 site is positive, and ex-change function of j1 and k2 sites is negative. When gallium occupies j2 site pref-erentially Curie temperature declines quickly with increase of gallium content when x 2 because positive exchange function is replaced by zero exchange func-tion. Also when x > 2 for sites j1 and k2 the part of negative exchange function is replaced by zero exchange function, that is in favor of raising of Curie tempera-ture. But generally, the probability to occupy j2 site by gallium is bigger than that to occupy other sites thus Curie temperature still declines when x > 2, and only it declines slowly compared with that when x 2. Therefore, using the Mössbauer method can well explain the variation tendency of the Curie temperature with change of gallium content.

Experiment result of X-ray diffraction indicates that X-ray diffraction patterns of 5 alloys with different composition are similar, the different is just some perks have some displacement. Fig. 4.25 shows the X-ray diffraction pattern when x = 1. It can be seen from the figure that Laves phase exists after adding gallium, which limits further enhancement of coercivity. Experiment result indicates that when x is increased from 4 to 7 X-ray diffraction peak offsets to direction of small angle, that indicates gallium enters into crystal lattice and lattice parameters a and c vary with increase of gallium content, the relative data is shown in Table 4.11.

Fig.4.25 The X-ray diffraction pattern of Nd16Co16Fe60GaB7 powder

Nd-rich phase;� B-rich phase; � Laves phase

Table 4.11 Relationship between gallium content and lattice parameters

Gallium content, x 0 1 2 4 7

a/nm 0.876,16 0.876,59 0.877,18 0.876,59 0.875,9

c/nm 1.212,72 1.215,70 1.217,98 1.219,8 1.222


It can be seen from these data that lattice constant a of the quadrilateral phase increases with increase of gallium content at beginning and reaches the maximum when x = 2; if x > 2 the constant decreases gradually; but lattice constant c in-creases with increase of gallium content.

It can be known from above analysis that for Nd16Co16Fe61-xGaxB7 gallium can enter into the quadrilateral phase and coercivity reaches the highest when x = 2.

Nd16Co16Fe61-xGaxB7 was edged at 580 for 3 h and then quenched to room temperature and prepared to film for electronic microscope as per procedures mentioned in section 2.3.2 of this book. Variation in crystal interface and interior of crystal was observed under electronic microscope. Three points were found as follows:

1. Crystal interface was flat and straight after edging, seeing Fig. 4.26 and Fig. 4.27.

Fig.4.26 Electron micrograph of Nd16Co16Fe59Ga2B7 alloy at room temperature

Fig.4.27 Electron micrograph of Nd16Co16Fe59Ga2B7 alloy at room temperature, observed

Ga-rich phase (Ga2Nd) between Matrix phase and crystal boundary phase


2. Ga-rich phase (Ga2Nd) enters into the matrix phase Nd2Fe14B and figure of Ga-rich phase appears in the matrix phase is shown in Fig. 4.28 to Fig. 4.30, of which microstructure looks like trip in the matrix phase indicating that Ga-rich phase and Nd2(Fe, Ga)14B, Nd2(Fe, Ga)14B and Nd2Fe14B are deep coherence.


Ga-rich phase precipitated in matrix phase Nd2Fe14B


Ga-rich phase (Ga2Nd) precipitated in Nd2Fe14B phase



block of Ga2Nd phase precipitated in Nd2Fe14B phase 3. Gallium is non-magnetic atom and gallium enters into magnetic phase to

form Nd2(Fe,Ga)14B, which seems to have magnetic dilution effect, thus molecu-lar magnetic torque is somehow declined, lattice constant c is increased and coer-civity of alloy is heightened.

4.6.4 Conclusions

NdCoFeGaB permanent magnetic alloy of high Curie temperature was prepared by means of powder metallurgy. Curie temperature of NdFeB permanent mag-netic alloy was heightened from 312 to 450-500 ; intrinsic coercivity of the alloy (iHc) reached to 716-955 kA/m, and magnetic energy product (BH)max = 223-262 kJ/m3. The study result using Mössbauer effect, etc. on sample of Nd16Co16Fe61-xGaxB7 indicated that when x = 2 gallium atoms mainly occupy j2 site; when x > 2 gallium atoms also have a certain probability to occupy sites k2, j1, k1 besides j2. X-ray diffraction experiment attested to that of the tetragonal phase of this alloy the lattice constant a became the maximum at x = 2 and the constant c increases with increase of x (from 1 to 7). The result of magnetic measurement found that the intrinsic coercivity presented a peak at x = 2.

4.7 Influence of Adding Element Dysprosium on Performance of NdFeB Alloy

On sintering of performance of NdFeB permanent magnetic alloy the biggest in-


fluence is oxygen content and the milling process is a key procedure causing high oxygen content in the alloy. After hydrogenated and crushed the powder is pulver-ized to 3-4 m is an effective method to reduce oxygen content in alloy (Pan, Ma, Ping, et al, 1991; Zhou, 1990).

It was found in study that the powder was hydrogenated and crushed to confect NdFeB permanent magnetic alloy where the crystal granule was abnormally big. The abnormal big size of the granule may reach 400 nm in magnet sintered at 1000 for 1.5 h. The formation of abnormal big granule will lower magnetic performance of the alloy. This abnormal big granule is very hurtful for magnetism, especially for coercivity of the alloy, which made the intrinsic coercivity largely degrade.

However, hydrogenated milling is a good method for milling of NdFeB perma-nent magnetic alloy because it can make oxidized content be lowered to the minimum limit. If a good method can be found that the intrinsic coercivity will not be degraded, i.e., to prevent the granule from growing big, then the purpose can be reached that the oxygen content can be lowered and the intrinsic coercivity does not be degraded. This section is to study this problem.


The alloy was melted in a middle frequency induction furnace with 99.5% pure metallic neodymium, iron and B-Fe alloy containing 18.4% of boron; the alloy ingot was hydrogenated and crushed, and then was pulverized into powder in ball milling. The powder with the granule size of 3-4 m was orientation molding pressed under magnetic field over 1.5 T and the molding pressure was bigger than 2 T/cm2. The molded billet was sintered in vacuum at 1100 for 100 min. For contrast Dy2O3 was added into hydrogenated coarse powder which was then pul-verized in ball miller. After sintered at high temperature the billet was put the magnet for aging at 500-600 and its magnetism was measured using loop in-strument of closed magnetic route.

4.7.2 Experiment result using SEM

Two types of samples with and without adding Dy2O3 after hydrogenated milling were observed by using SEM. The micro texture of permanent mag-netic alloy without Dy2O3 is shown in Fig. 4.31(a) and the micro texture with Dy2O3 is shown in Fig. 4.31(b). It can be seen from Fig. 4.31 that the per-manent magnetic alloy with adding Dy2O3 has comparative smaller crystal granule.


Fig.4.31 SEM structure of NdFeB permanent magnet alloy and SEM structure of

NdFeB with addition of Dy2O3 (a) No addition of Dy2O3; (b) With addition of Dy2O3

4.7.3 Measurement of magnetism

The intrinsic coercivity iHc, of the permanent magnetic alloy after adding Dy2O3

was enhanced to 1474 kA/m from the original 660.5 kA/m. The data of Br, bHc and (BH) max are shown in Table 4.12. It can be seen from comparison of 4 groups data in the table that after addition of Dy2O3 the iHc was enhanced from 660.5 kA/m, 851.2 kA/m, 774.3 kA/m and 802.3 kA/m to 1,474.0 kA/m, 1,451.0 kA/m, 1,254.0 kA/m and 1,273.0 kA/m respectively; bHc was enhanced from 635.0 kA/m, 778.0 kA/m, 721.0 kA/m and 248.9 kA/m to 776.1 kA/m, 791.5 kA/m, 794.3 kA/m and 808.7 kA/m, respectively. And that Br and (BH) max were degraded slightly. Br was degraded from 1.206T, 1.228T, 1.193T and 1.209T to 1.035T, 1.057T, 1.053T and 1.076T respectively; and (BH) max was degraded respectively from 274.0kJ/m3, 285.0kJ/m3, 268.0kJ/m3 and 277.0 kJ/m3 to 202.7kJ/m3, 204.3kJ/m3, 204.4kJ/m3 and 212.7 kJ/m3. The demagnetization curve is shown in Fig. 4.32.

Table 4.12 Magnetic parameters of NdFeB alloy before and after adding Dy2O3

Features Br/T bHc/kA�m�1 jHc/kA�m�1 (BH)max/kJ�m�3 1.206 635.0 660.5 274.0

1.228 778.0 851.2 285.0

1.193 721.0 774.3 268.0 Without adding

Dy2O3

1.209 748.9 802.3 277.3

1.035 776.1 1474.0 202.7

1.057 791.5 1451.0 204.3

1.053 794.3 1254.0 204.4 With adding

Dy2O3

1.076 808.7 1273.0 212.7


Fig.4.32 Demagnetization curves of the specimen with addition of Dy2O3 and

specimen without addition 1 No addition of Dy2O3; 2 With addition of Dy2O3

4.7.4 Experiment result using transmission microscope

Microstructure of NdFeB permanent magnetic alloy by hydrogenated milling with and without addition of Dy2O3 is shown in Fig. 4.33. It can be seen from Fig. 4.33 that the microstructure of NdFeB permanent magnetic alloy, no matter base phase or crystal boundary, was changed because of addition of Dy2O3. By looking at the microstructure of NdFeB alloy without addition of Dy2O3 its crystal boundary and base phase are all very clean and simple. The small triangle in Fig. 4.33(a) is Nd-rich phase. And most of this type of Nd-rich phase appears in triangle. After addi-tion of Dy2O3 the figure of Nd-rich phase was also changed so that it is no more a simple triangle crystal boundary figure. Because dysprosium appeared both in crystal boundary and base phase the microstructure become complicated, that appeared many twin-crystals and even Nd2O3.

4.7.5 Distribution of Dy2O3

There were many disputes on distribution of Dy2O3 since a long time ago. Somebody considered that the Dy2O3 is distributed in crystal boundary and some-one thought that mainly distributed in the base phase. Therefore, we conducted analysis on the observed base phase and crystal boundary using electronic probe at the same time when we carried out the observation by electronic transmission. It indicated in energy spectrum of probe analysis that Dy2O3 and Dysprosium all distributed in both base phase and Nd-rich phase.


Fig.4.33 Microstructure of NdFeB prepared with hydrogen pulverization

(a) No addition of Dy2O3; (b) With addition of Dy2O3

Dy2O3 was added in milling process. After high temperature sintered at 1100 the Dy2O3 powder was reacted with NdFeB ternary alloy powder as follows:

16 77 7 2 3 16-2 2 77 7 2 3Nd Fe B + Dy O Nd Dy Fe B + Nd Ox xx x++, (4.8)

Thus discussing that Dy2O3 entered to what type of phase or distribution is to discuss tendency and distribution of dysprosium. The Table 4.13 expressed ana-lyzed data of Nd-rich phase by electronic probe.

Table 4.13 Probe analysis on B-rich phase of NdFeB alloy after adding Dy2O3

Element Mass fraction/% Mole fraction/%

Nd 86 79

Dy 7 5

Fe 7 16


It can be seen from Table 4.13 that the dysprosium was enriched in original 4% of the alloy to 5% in mol fraction (mass fraction 7%) in Nd-rich phase. And from Table 4.14 the dysprosium in base phase was reduced 1% from 4% (mole fraction) of the alloy to 3% (mole fraction). Above mentioned indicated that dys-prosium distributed more in Nd-rich phase than in base phase. Already existent studies indicated that when added dysprosium is less than 4% (mole fraction) the situation would exhibit different. While dysprosium content is 1.5% the distribu-tion of dysprosium will be equal in both base phase and Nd-rich phase; and that if the content is less than 1% then the dysprosium distributes more in base phase than in Nd-rich phase (Pan, Ma, Ping, et al, 1991; Zhang, Lu, 1992).

Table 4.14 Probe analysis on the matrix phase of NdFeB alloy after adding Dy2O3

Element Mass fraction/% Mole fraction/%

Nd 42 23

Dy 6 3

Fe 52 74

From abovementioned tests it can be analyzed that Dy forms (Nd, Dy)2O14B

phase in NdFeB alloy, and it was known from the reference (Pan, 1996) that the anisotropy of Dy2Fe14B (11.94 MA/m) was much bigger than that of Nd2Fe14B (7.16 MA/m); and taking Dy2Fe14B as the matrix term the DyFeB permanent magnetic material had jHc = 3.98 MA/m, which is the highest coercivity obtained up to date (Zhou, 1990); thus using part of Dy to substitute Nd can highly en-hance the intrinsic coercivity. In addition, it can be seen from Fig.4.31 that adding Dy2O3 can make crystal grain of alloy fine. The grain size shown in Fig.4.31(b) is about 3 m in average. Studies shown that adding content of Dy became 0.4% (mole fraction) the average grain size was 7 m; when Moore fraction increased to 2% then the average grain size was 5.1 m; and the Moore fraction was 3% the average grain size became 3 m. As the result of fined grain the area percentage of grain boundary will increase comparatively; and percentage of Dy distributing in Nd-rich phase and grain boundary is bigger than in the matrix phase with in-crease of Dy content, which makes difference of Dy between grain boundaries and the matrix phase bigger and thus changes the microstructure of the alloy (Fig.4.33). And because of the increase of such difference the move of domain becomes difficult; to move the domain it is necessary to increase the reverse field so that the coercivity of the alloy is enhanced.

X-ray diffraction analysis indicated that compared adding 1% Dy (mole frac-tion) with not adding Dy into Nd15Fe77B8 alloy, the lattice constant a was reduced from 0.883 nm to 0.870 nm, c was reduced from 1.216 nm to 1.211 nm.

Obviously, adding Dy2O3 in hydrogenated pulverization of NdFeB permanent


magnetic alloy preparation process can effectively control grain growth up in high temperature sintering of this alloy; and adding appropriate amount of Dy2O3 in hydrogenated pulverization can increase anisotropy of (Nd, Dy)2Fe14B not only, but also make grain fine, thus to largely enhance the intrinsic coercivity of the alloy.

4.7.6 Conclusions

By aforementioned studies we can conclude as follows: 1. Hydrogenated pulverization in NdFeB alloy preparation process resulted in

abnormal growth up of grain size in high temperature sintering, which declines the intrinsic coercivity of the alloy.

2. Adding appropriate amount of Dy2O3 can act function to prevent growth up of grain size, and due to forming (Nd, Dy)2Fe14B tetragonal phase with very high anisotropy thus to enhance the intrinsic coercivity of the alloy.

3. When adding Dy to enhance the coercivity of NdFeB alloy if Dy is higher than 4% (mole fraction) Dy distributes in Nd-rich phase and grain boundary higher than in the matrix phase.

4.8 Nanocrystalline Microstructure and Coercivity Mechanism Model of NdFeB Alloys with Nb and Ga

Addition of Nb and Ga can increase the coercivity of NdFeB alloys. But the coer-civity is far less than that of the theoretical value. So up to now, the research on the coercivity mechanism of NdFeB alloys is still the research focus. Previous investigations pointed out that there were mainly three coercivity mechanism of rare earth iron-based permanent magnets: (1) nucleation hardening model, (2) boundary local pinning model and (3) uniform pinning model (Hadjipanayis, et al, 1988; Kronmüller, Durst, Sagawa, 1988; Durst, Kronmüller, 1987; Pan, Pan, Ma, 1994).

Hadjipanayis, et al (Liu, Pan, Luo, et al, 2004) believed that the boundary local pinning model played the main role on the coercivity mechanism. However, Kronmüller, et al (Kronmüller, Durst, Hock, et al, 1988) thought the coercivity was mainly determined by nucleation hardening model. They both had some ex-periment proofs, but neither of them was flawless. Kronmüller supported the nu-cleation hardening model that based on the alloys could reach the saturation mag-netization in a lower magnetization field. But these phenomena could also be ex-plained by boundary local pinning model. Therefore, these models could not ex-plain comprehensively the behavior of the domain wall in the boundary.

In this section, we will highlight the reason for the addition of Nb and Ga to


improve the coercivity of NdFeB alloys. From the angle of nanocrystalline micro-structure and magnetic properties, we study the coercivity mechanism of the al-loys and put forward some new experiment proofs.

4.8.1 Experimental procedure

The NdFeB alloys are prepared by vacuum melting and mold casting (protected by argon) using the high pure rare earth Nd and Dy (99.99%), high pure metal iron, cobalt and niobium (99.9%), ferroboron alloy with 18% (wt.) boron accord-ing to the alloy composition proportion. The alloys are cool down to room temperature. The alloy ingot is milled into powder with a particle diameter about 3-4m under the protection of organic solution. The powder is molded in mag-netic field (1.5T) under the press (2 t/cm2). The compacts are sintered at 1,090-1,100 for 1.5-2 h, then cooled down to 900 and solutionized for 1 h, aged at 600 for 50 min and cooled to room temperature. Thus thermal-demagnetized samples are obtained.

In order to study on the effects of Nb and Ga addition amount on the magnetic properties, Nd13Dy2(Fe1-xNbx)79B6 was prepared with x = 0.01, 0.02, 0.03, 0.04, 0.05, 0.08 and 0.16, respectively.

Samples were magnetized in the magnetizing field (> 4T). The magnetic prop-erties, namely, remanence and coercivity are measured using a magnetic parame-ter instrument. The energy product is calculated by demagnetization curve. The Curie temperature of the alloys is measured by vibrating sample magnetometer.

For the scrutinization of the alloy microstructure by ultra-high voltage (1000kV) microscope, the samples are cut into 0.25mm slice vertical to c axis, milled and cleaned using ion beam thinner. The ultra-high voltage microscope model is JEM-1000kV: operation voltage 1000kV, vacuum degree 2.5×10-4Pa and 0.7×10-5Pa with addition of liquid nitrogen.

The Mössbauer spectroscopy of the alloys is measured by model Oxford-ms500 with 57Co/Rh radioactive source and �-Fe velocity calibration.

4.8.2 Magnetic properties measuring

The samples are magnetized in pulsed magnetic field (>4.5T), then measured us-ing magnetic parameter instrument, seeing Fig.4.34. When x=0.02, the intrinsic coercivity increased by 48%, which is the maximum value. When x is above 0.02, the intrinsic coercivity, remanence and energy product monotonically decrease. The magnetic properties of Nd13Dy2(Fe0.98Nb0.02)79B6 are: Br=1.38T, iHc=1326 kA/m, (BH)max=380 kJ/m3.


4.8.3 Study of Mössbauer effect

This section studied the magnetic properties of Nd15Fe85-xBx alloys, Nd15(Fe1-xNbx)78B7 alloys and Nd2Fe12-xCo2NbxB (x=0.05, 0.10, 0.15 and 0.20) alloys at the atom level by means of Mössbauer spectrum.

4.8.3.1 Mössbauer analysis of Nd15Fe85-xBx alloy at room temperature

It can be seen from the Mössbauer spectrum of Nd15Fe85 alloy (x=0) that there are six rays of �-Fe and new absorption rays of b� with the velocity of -1.8mm/s and 1.2mm/s. There were obvious broadening and dissymmetry at the second and fifth rays of �-Fe. So the velocity of the new phase absorption rays will be �3-3mm/s. The hyperfine field is 180kOe. The absorption rays of d and d� of Nd2Fe14B phase (magnetic phase) appear in the Mössbauer spectrum after adding a little boron in the alloy (Nd15Fe83B2). Comparing with Nd15Fe85 alloy, the Mössbauer absorption rays intensity of the Nd15Fe83B2 alloy decrease evidently. When x=4 (the nominal composition is Nd15Fe81B4), the rays of Nd2Fe17 nearly disappeared. The content of �-Fe is far less than that of Nd15Fe81B4. The intensity of �-Fe rays is very weak. When x=8 (the nominal composition is Nd15Fe77B8), there are only rays of Nd2Fe14B (magnetic phase, tetragonal phase). The rays of �-Fe disappear. Besides, there are new rays, which are paramagnetic rays of Nd-rich phase. Therefore, the addition of appropriate boron in the alloy can improve the magnetic properties of the alloy because Nd2Fe14B phase has high anisotropy. With increasing the boron content, the contents of �-Fe and Nd2Fe17 decrease. It can be seen from the Möss-bauer spectrum of Nd2Fe77B8 alloy that Fe6 grain positions occupy 27%, para-magnetic phases grain positions occupy 6%. Every Nd2Fe14B crystal cell has 16 Fe5 & Fe6, 82 Fe3 & Fe4, 4 Fe1 & Fe2 (Ping, Pan, 1985; Zhao, Xia, Ma, Pan, 1989). The neutron diffraction of the alloy by Herbest et al shows that every Nd2Fe14B crystal cell had 68 atoms: 56 Fe, 8 Nd and 4 B.

According to the size of the hyperfine field of Nd2Fe14B, the order should be

2 1 2 18 16 16 8 4 4B >B , B > B , B >Bj k k j e c . According to the size of the intensity, the order

should be 1 2 1 216 16 8 8 4 4, > , > , k k j j e cI I I I I I . The neighbour atom number of RE2Fe14B

was shown in Table 4.15 (RE represents Y, Sc, La, Ce, Pr, Nd, Sm, Eu, Ga, Tb, Dy and Ho) (Qi, 1998).

There are many neighbour RE atoms in grain positions 8j1 and 4c while there are few neighbour RE atoms in grain positions 16k1, 16k2, 8j2 and 4e. So the RE elements have little effect on the hyperfine field of RE2Fe14B. The hyperfine field of RE elements increases gradually from Y to Gd. But the hyperfine field of RE elements decreases accord to the order: Gd, Tb, Dy and Ho. The hyperfine field B

——

hf of Nd2Fe14B is 29.55T.


Table 4.15 Neighbor atoms number of RE2Fe14B

Grain position Neighbour atom number Grain

positions Fe16k1 Fe16k2 Fe8j1 Fe8j2 Fe4e Fe4c RE4f RE4g B4g Total Fe RE

Fe16k1 2 3 1 2 1 1 1 1 1 13 10 3

Fe16k2 3 3 1 2 0 1 1 1 0 12 10 2

Fe8j1 2 2 1 3 1 0 1 2 0 12 9 3

Fe8j2 4 4 3 0 1 0 1 1 0 14 12 2

Fe4e 4 0 2 2 1 0 2 0 2 13 9 2

Fe4c 4 4 0 0 0 0 2 2 0 12 8 4

RE4f 4 4 2 2 2 2 0 2 2 20 16 2

RE4g 4 4 4 2 0 2 2 1 1 20 16 3

B4g 4 4 0 0 0 2 0 1 0 7 6 1

4.8.3.2 Mössbauer analysis of Nd15(Fe1-xNbx)78B7 alloys at room temperature

The occupying probability of Fe and Nb in tetragonal phases of Nd15(Fe1-xNbx)78B7 alloys is shown in Table 4.16. It indicated that Nb primarily occupy e and c grain positions. There are few neighbour atoms in grain positions e and c, which can not result in the hyperfine field change. When x 0.04, many Nb atoms enter into grain boundary, Nb-rich and Nd-rich phases rather than enter into tetragonal phases. Therefore, Nb has little effect on the saturation magnetization Ms of Nd2Fe14B phase. When x 0.08, Nd2Fe14B phase is destroyed and is replaced by new phase.

Table 4.16 Occupying probability of Fe and Nb in tetragonal phases of Nd15(Fe1-xNbx)78B7 alloys

x Occupation probability j2 k2 k1 j1 e c

PFe 1.006 1.007 0.999 0.985 0.948 0.934 0.01

PNb �0.006 0.007 0.001 0.015 0.052 0.066

PFe 0.982 0.993 0.986 0.958 0.972 0.958 0.02

PNb 0.018 0.007 0.014 0.042 0.028 0.042

PFe 0.951 0.968 0.965 0.958 0.951 0.923 0.04

PNb 0.149 0.032 0.035 0.042 0.049 0.077

4.8.3.3 Mössbauer analysis of Nd2Fe12-xCo2NbxB(x=0.05, 0.10, 0.15 and 0.20)

The Mössbauer spectrum of 57Fe is formed by six rays. It indicates that hyperfine field is digression according to 8j2, 16k2, 8j1, 16k1, 4e and 4c. Nb primarily occu-pies j2 grain position. The distance of neighbor Fe atoms is above 0.244nm (Pan, Ma, Li, 1993). The exchange effect of j2 grain position is positive. Nb enters into 8j2 grain position and the exchange effect is positive, which decreases Tc of the alloys.


4.8.4 Study of nano-microstructure of NdFeB alloys with Nb

The Hc of NdFeB alloys can increase to 1325 kA/m by adding appropriate content of Nb. Nb has little effect on the anisotropic field HA because there are few Nb in the tetragonal phases of the alloys. Fe2Nb phases emerge in the microstructure of Nd15Fe77B8 alloys with Nb. Fe2Nb is MgMn2 type structure, whose lattice con-stants a and c are 0.482nm and 0.787nm, respectively. The size of Fe2Nb is only 2-4 nm, which mainly disperses in Nd2Fe14B phase and Nd-rich phase coherent with Nd2Fe14B phase. The Nb mainly enriches in grain boundary, which changes the triangle microstructure of the grain boundary of Nd15Fe77B8 alloys. The re-search shows that there are small Nb-rich precipitates in the grain boundary be-sides Fe2Nb phase. The percentage composition of Nb in the precipitates is 89%-97%. It indicates that Nb diffuse into Nb-rich phases and form new particle dur-ing aging treatment. The diffusion is propitious to form the intact grain boundary. The crystal defects of the alloys increase after aging treatment. According to the nucleation hardening model, every defect is a position that is easy to form inverse nuclear. In general, the defects’ nucleation field is low, which will decrease the intrinsic coercivity of the alloys. However, the addition of Nb can increase the intrinsic coercivity of the alloys. So the result contradicts with the nucleation hardening model. Therefore, the nucleation hardening model can not explain the phenomena (Pan, Liu, Luo, 1990; Zhao, Xia, Ma, Pan, 1989).

By measuring the grain sizes of the NdFeB alloys with Nb, we find that the Nb can inhibit the growth of the grain and decrease the sizes of the grains. The total surface area of the grains and the domain walls pinning increase, which improves the intrinsic coercivity of the alloys. Therefore, Nb has cross effect on the alloys, which can be explained by authors using the theory of “dynamic cross – complementary combination”. The appropriate addition of Nb can decrease the sizes of the nano-grains and increase the grain boundary and domain walls pinning, which increases the intrinsic coercivity. Above the discussion can be summarized “grain finery local pinning”.

4.8.5 Dynamic cross and microstructure of the NdFeB alloys with Nb and Dy

The magnetic properties of the NdFeB alloys with addition of Nb and Dy are show in Table 4.17.

Table 4.17 Magnetic properties of NdFeB with Nb and Dy No. Nominal composition Br/T mHc/kA·m-1 (BH)max/kJ·m-3 HA/T

1 Nd13.5Dy1.5Fe80B6 1.16 1428.6 259.8 8.47

2 Nd13.5Dy1.5Fe78Nb2B6 1.11 1639.8 252.6 8.48

3 Nd12.5Dy1.6Fe78Nb2B6 1.08 1765.9 248.7 8.49


The contents of Nb of the samples No.1, No.2 and No.3 are 0%, 2% (at.) and 2% (at.), respectively. It can be seen that the mHc of samples No.2 and No.3 in-creases 13% and 19% than that of sample No.1, respectively. The anisotropic field HA of the alloys has few changes. We have studied the microstructure using TEM in order to explain the effect of Nb on the alloys. There are strip phases in the microstructure of the alloys with Nb (shown in Fig. 4.34). The strip phase is Fe2Nb, which is called “Laves phase”. The average size of the grain of the alloys with Nb is 5.6 nm while that of the alloys without Nb is 9.3 nm. So Nb can inhibit the grain growth and decrease the grain size, which increases the grain surface area and domain walls pinning.

Fig.4.34 Electron micrograph of (Nd0.9Dy0.1)15Fe76Nb2B7 permanent magnet alloy at

room temperature, observed Laves phase Fe2Nb precipitated in Nd2Fe14B phase (Liu, Luo, Pan, et al, 1991)

The effect of Nb increasing jHc is due to its finer grain structure. Hu Jiafa etc.

add 1.75% Nb powder to Nd16Fe77B7 powder with inter-grain alloying. jHc reaches 14.6kOe(1162.16kA/m) compared to 11.7kOe(931.32kA/m) with no addi-tion of Nb powder.

4.8.6 Dynamic cross and microstructure of the NdFeB alloys with Nb, Ga, Co and Dy

Addition of Ga in NdFeB alloys replaces Fe, which can obviously improve the coer-civity. There is appropriate ratio of combined addition of Nb and Ga in the alloys. The NdFeB alloys with nominal composition of Nd13Dy2Co4Nb0.8Ga1.3B7Fe71.9 are pre-


pared by reasonable process. The intrinsic coercivity is above 2000kA/m. The sintering temperature was T 1105 . Aging treatment is key factor that affects the coercivity of the alloys. The appropriate aging temperature is according to the composition of the alloys. When the aging temperature is 525-630 , the NdFeB alloys have high coercivity. The existing forms of Ga and Nb in the alloys are difference. Nb elements mainly exist in matrix phases with the forms of precipi-tates and inclusion morphology. The grain boundaries have a few Nb elements. Besides, there are some strip Fe2Nb Laves phases (Fig. 4.34). Nb elements can decrease the grain size when the content is less than 1.5%-2% (at.). Ga elements exist in the matrix phases and grain boundaries with the form of Ga-rich phases (Fig. 4.35). The content of Ga for replacement of Fe should be 1%-2% (at.). The existing form of Ga in the grain boundaries is same to that in the matrix phases, which inhibits the formation of soft magnetic phases (Fig. 4.36). The addition of Ga elements can decreases the grain size and increases the grain boundary area, domain walls pinning effect and the coercivity of the alloys.

Fig.4.35 Electron micrograph of Nd15Co16Fe60Ga2B7 permanent magnet alloy at room

temperature, observed Ga-rich phase (Ga2Nd), Nd2Fe14B phase and grain boundary between Nd2Fe14B crystal boundaries A Ga�rich phase; B Matrix phase; C Grain boundary

4.8.7 Curie temperature of the NdFeB alloys with Nb

The Cure temperature of the NdFeB alloys is decreased by addition of Nb. The re-search of Mössbauer on the alloys shows that the matrix phases Nd2Fe14B of the alloys have exchange action: JFe-Fe>JFe-RE>JRE-RE. So the Curie temperature is deter


Fig.4.36 Electron micrograph of Nd15Fe78B7 permanent magnet alloy at room temperature, observed Nd-rich phase (C), Nd2Fe14B phase (B) in the

triangle grain boundary of Nd2Fe14B mined by the distance of neighbour Fe. JFe-Fe, JFe-RE and JRE-RE are the exchange con-stants of Fe-Fe, and Fe-RE and RE-RE, respectively. The interatomic exchange integral of Fe is negative when the distance is less than 0.244nm. When the distance is above 0.244nm, the exchange integral is positive (Herbst, Croat, Yelon, 1985). The distances of iron atom in the site j2 between in the sites k1, k2 and j1 are 0.2748nm, 0.2640nm and 0.2784nm, respectively. Therefore, the exchange integral of iron atoms in the site j2 is positive. When the niobium atoms occupy the site 8j2, the exchange integral is change into zero, which results in the decrease of Curie temperature of the alloys. The hyperfine field parameters of the alloys are shown in Table 4.18.

Table 4.18 Hyperfine field parameters of Nd2Fe12-xCo2NbxB(x=0, 0.2) alloys

Samples Grain positions Hyperfine field Hhf /kOe 4 split �/mm·s�1 Displacement

�/mm·s�1

16k2 275 0.325 �0.005

16k3 296 0.329 0.024

8j1 287 0.170 �0.280

8j2 333 0.650 �0.001

4e 269 �0.43 �0.190

Nd2Fe12Co2B

4c 260 0.110 �0.280



Samples Grain positions Hyperfine field Hhf /kOe 4 split �/mm·s�1 Displacement

�/mm·s�1

16k2 271 0.315 �0.03

16k3 291 0.320 �0.025

8j1 282 0.12 �0.480

8j2 328 0.668 �0.020

4e 265 �0.480 �0.290

Nd2Fe11.8Co2Nb0.2B

4c 255 �0.260 �0.002

4.8.8 New coercivity mechanism model of multi-component NdFeB alloys

The coercivity mechanism of the NdFeB alloys is still research focus in recent decades. Kronmüller, et al had studied systematically the coercivity mechanism from theoretical derivation to experimental verification. They published many papers and pointed out a nucleation hardening model. Hajipanayis, et al. put for-ward boundary local pinning model, which could explain many questions and brought great influence on the research of the mechanism. The third model is uni-form pinning mechanism model, which is similar to the second model in some aspects (Ding, Pan, Luo, 1990; Zhou, 1995; Pan, Li, Li, et al, 1989; Liu, Luo, Pan, et al, 1991; Liu, Pan, Luo, et al, 1991; Ping, Li, Ma, Pan, et al, 1986; Pan, Pan, Ma, 1994; Liu, Pan, Luo, et al, 1990).

Is the coercivity of NdFeB alloys determined by nucleation hardening or uni-form pinning? The effect of temperature, magnetization field and anisotropic field on the coercivity should be studied in order to make clear the question.

The intrinsic coercivity jHc of the alloys increases greatly after aging at 600�900 . jHc linearly increases with increasing magnetization field. The main cause is that the coercivity is determined by nucleation hardening and pinning hardening at different temperature, respectively. When the magnetization field is not less than the saturation field of the coercivity, the coercivity is determined by the nucleation field of the magnetization reversal. The coercivity is determined by pinning field when the magnetization field reaches pinning field. It seems that the statement is comprehensive. But the coercivity mechanism of the isotropic sin-tered NdFeB and rapid-quenched NdFeB magnets is not explained by nucleation hardening model. Mishra et al. pointed out that the coercivity of rapid-quenched NdFeB magnets derived from the pinning field of the grain boundary to the Bloch domain wall.

The nucleation field of the matrix phase Nd2Fe14B of the alloys should be equal to the anisotropic field of the alloys. But jHc is far less than the anisotropic field of


the alloys. The conceivable reason is that the matrix phase Nd2Fe14B has many surface defects and forms anti-magnetization domain. The surface defects and roughness of the matrix phase is improved after aging treatment, which increases the coercivity.

Both “nucleation hardening model” and “uniform pinning model” can not ex-plain the coercivity mechanism of the NdFeB with above ternary alloy elements (additive elements such as Nb, Al, Ga etc.). Author gives out “grain finery local pinning” model, which can explain the coercivity mechanism of the NdFeB with above ternary alloy elements especially the alloys with Nb. The coercivity mechanism has been tested by experiments (Pan, Li, Li, et al, 1989; Liu, Luo, Pan, et al, 1991; Liu, Pan, Luo, et al, 1991; Liu, Luo, 1989).

4.8.9 Conclusions

Following are the conclusions of above investigations: 1. The Mössbauer analysis of the Nd2Fe12-xCo2NbxB(x=0.05, 0.10, 0.15, 0.20)

alloys indicates that HA (anisotropic field) increases while Ms (saturation magneti-zation field) and Tc (Curie temperature) decrease with Nb content increasing. The nano-microstructure of (Nb0.9Dy0.1)15Fe78B7 alloys shows that there are Fe2Nb phases in the structure. The lattice constants are: a=0.482nm, c=0.787nm. The addition of Nb can decrease the grain size and mainly rich in grain boundary, which increases greatly the intrinsic coercivity (jHc). Nb has little effect on the anisotropic field HA because the elements mainly exist in the grain boundary and form Laves phases rather than enter into the crystal lattice. The microstructure of the alloys is dynamic tested by TEM with 1000kV from room temperature to 400 . It shows that there is not precipitates and change in the microstructure. But some precipitates emerge from Nd2Fe14B phases of the NdFeB alloys without Nb at 280 . Therefore, Nb and Ga can improve the thermal stability of the alloys. The nucleation hardening model and pinning model can not comprehensively explain the coercivity mechanism of NdFeB alloys with Nb and Ga. Author gives out a new “grain finery local pinning” model. This model can explain the phe-nomena of any addition element increases the coercivity and improves thermal stability of NdFeB ternary alloy, which was the result of grain refinement and increment of the boundary local pinning.

2. The Mössbauer analysis of the Nd15(Fe1-xNbx)78B7 (x 0.04) indicates that Nb primarily occupy e and c grain positions. But there are few neighbor atoms in e and c grain positions. Therefore, there are few Nb atoms in e and c grain posi-tions, which can not result in the hyperfine field change. Nb atoms mainly occupy 8j2 grain position.

3. The microstructure of NdFeB ternary alloys and NdFeB alloys with Nb are dynamic analyzed by TEM with 1000kV. It indicates that the ternary alloys with


nominal composition Nd15Fe77B8 produce precipitates (Nd2O3) when the tempera-ture reaches 280 . But the alloys with Nb begin to produce precipitates when the temperature is above 400 . Nb can decrease the grain size and improves the magnetic properties of the alloys.

4.9 In Situ and Dynamic Observation on Magnetic and Phase Transformation of Nd15Fe78B7 Permanent Magnet at High Temperature *

It is appropriate to research magnetic hardening mechanism and determine use temperature of NdFeB permanent magnet by dynamic observation on microstruc-ture of Nd15Fe78B7 permanent magnet at a high voltage microscope and trying to connect phase transformation and phase precipitation with macro magnetic prop-erty of the material. There are many experiments and theoretic studies on mag-netic hardening mechanism of NdFeB magnet and presented a crystal structure of Nd2Fe14B. There have two opinions on origin and mechanism of high coercivity of NdFeB permanent magnet alloy because NdFeB permanent magnet alloy con-tains a little more Nd and B than Nd2Fe14B compound formula weight.

1. High coercivity of Nd2Fe14B alloy is originated single- axis anisotropy. 2. Coercivity of single phase Nd2Fe14B is very low. Coercivity of NdFeB per-

manent magnet alloy is mainly controlled by nucleation process of anti-magnetic nucleus. The B-rich phase, Nd-rich phase and especially Nd-rich phase among crystals of Nd2Fe14B with low magnetic anisotropy act as nucleation center in the center of anti-magnetic nucleation. Once the anti-magnetic nucleus formed the low magnetic anisotropy area would act as function of nails in the domain wall. Expansion field of these nucleus further determine the strength of coercivity.

There still is another opinion on this problem that the intensity of coercivity is determined by B-rich phase.

Nevertheless, these opinions may have somewhat limitation because all of these opinions were derived in a static state (Sagawa, Fujimura, Yamamoto, et al, 1984; Sagawa, Fujimura, Togawa, et al, 1984).

We used an 1000 kV HVEM to make dynamic observations on change of mi-crostructure of sintered Nd15B7Fe78 permanent magnet alloy at temperatures from 25 to 600 . It is observed that the base phase of Nd2Fe14B have highly dis-persed precipitate and the precipitate phase grows continually along with rising of the temperature. At 600 the precipitation becomes polycrystals. We also ob-served that aberrance occurs in diffraction pattern of B-rich phase at 332 and

Cooperators of this study are Fengzuo Tian, Jixian Sun, Ansheng Liu, Shikuan Ren, Chengqin Huang, Guoqing Yao, and Qiming Ying, General Research Institute for Nonferrous Metal.


Nd-rich filmy belt in the boundary of crystals become wider proportionally along with rising of the temperature (Ping, Li, Ma, Pan, et al, 1985; Pan, 1986).

4.9.1 Preparation process of specimen and experiment method

The alloy was prepared by using a casting piece composed of high purity neodymium, high purity boron and high purity iron; and then pulverized the casting piece into powder of 3.2 m in medium of toluene. The powder was formed under 1.5 T magnet field and sintered for one hour at 1100 , then placed for aging at a lower temperature and quenched to room temperature as specimen for electron microscope.

The sample described above was sliced to into lamellae of 0.25 mm in the di-rection being vertical to c axis and thinned mechanically to 0.025 mm. Then the lamella was electrolyzed to open a hole basically in an electrolyte of 20% per-chloric acid and 80% glacial acetic acid and finally be thinned and cleaned on ionic thinning device. Then the specimen was ready for observation.

Observation of specimen film was conducted in JEM-1000 HVEM. The opera-tion voltage was 1000 kV, the output voltage was 185 V and the electrical current was 6.6 A. Vacuum was kept as 2.13 ×10�3 Pa (equivalent to 1.6×10�5 Torr), after adding liquid nitrogen the vacuum was adjusted to 9.33×10�6 Pa (equivalent to 0.7×10�7 Torr). Ionic beam was 10 A. The specimen was then inserted into the side inserting type heating dais of JEM-1000 and observed by electronic micro-scope in the condition of room temperature and heat condition.

4.9.2 Microstructure and phase in crystal boundary of NdFeB permanent magnet

Many types of crystal boundary exist in NdFeB permanent magnet: the 1st type of crystal boundary is flat and clean; ternary and multi-component alloy have this type of crystal boundary uniformly as shown in Fig. 4.37. Enriched materials of one or more components exist in the crystal boundary. It was confirmed to be en-riched neodymium by electronic diffraction. The enriched neodymium has a cubic (fcc) structure which was testified by analysis using energy spectrum. The grain boundary becomes clean through a proper annealing temperature system. The 2nd type is a multiphase crystal boundary; some even form big neodymium enriched clusters or multiphase neodymium enriched clusters, as shown in Fig. 4.38. There still exists pectinate or stria lattice shapes contrast inside of NdFeB base grain boundary by adding gallium, that is analyzed as iron enriched component by en-ergy spectrum.


Fig.4.37 Electron micrograph of Nd15Fe78B7 permanent magnet alloy at

room temperature, observed Nd-rich phase and favorable grain boundary after annealing in the triangle grain boundary of Nd2Fe14B

Fig.4.38 Electron micrograph of Nd15Fe78B7 permanent magnet alloy at

room temperature, observed Nd-rich phase in different pattern at the join point of crystal boundary of Nd2Fe14B, and observed crystal boundary of the

specimen annealed not at optimized temperature


Alloy Nd15Fe78B7 generally has magnetic phase of Nd2Fe14B and Nd-rich phase, and a little �-Fe (Fig. 4.39). Alloy containing gallium and niobium has Ga-rich phase and Fe2Nb phase (Laves phase) without exception. The ternary eutectic temperature of the above mentioned 3 types of phases has big difference in com-parison with the melting points of neodymium and iron; and there is a deep eutec-tic tendency by holding an out of order liquid structure to room temperature after aging and cooling. Nd-rich phase distributes unevenly. This inhomogeneous dis-tribution is related to uneven cooling. Eutectic components (Nd2Fe14B, Nd1.1Fe4B4 and Nd-rich phase) are enriched as an out-of-order deep eutectic state, such as the magnetic phase of NdFeB multi-component alloy containing niobium and gallium, and its magnetic phase becomes Nd2(Fe, Ga)14B, Nd2(Fe, Nb)14B, Fe2Nb, , en-riched in a form of the out-of-order deep eutectic status.

Fig.4.39 Electron micrograph of Nd15Fe78B7 permanent magnet alloy at room tempera-

ture, observed clean and clear Nd2Fe14B phase and B-rich phase with concentrate stacking fault, Nd-rich phase at join point of B-rich and Nd2Fe14B

Coercivity of NdFeB alloy is not high after sintering. And the coercivity may

be doubled by annealing at a temperature range of 550-650 . Effect of annealing behaves on the crystal boundary that result in homogenization of the neodymium enriched boundary and Nd-rich phase. Crystallization of deep eutectic structure results in that the neodymium enriched particles to form a dense distribution of granular material in cross coign of crystal boundary, and sometimes represents crystal nucleus of Nd2Fe14B that make the around lean iron also form Nd-rich phases. Iron content of Nd15Fe78B is the maximum in NdFeB magnetic alloys. The base phase contains a certain solid solute iron and a long time annealing lead to dissolution of iron atoms of Nd2Fe14B, that causes modulated decomposition on Nd2Fe14B (Fidler, Luo, 1985).

Composition of Nd-rich phase is complicated but it acts an important function for magnetic hardening of sintered NdFeB alloy. The typical Nd-rich phase is of a


twin- hexagon structure, its lattice constants are a = 0.365 nm, c = 1.180 nm, the neodymium atomic fraction occupies 97% and the balance is iron atoms. The bright belts in Fig.4.40 and Fig. 4.41 surround particles of Nd2Fe14B. Its atomic proportion is Nd 75%, Fe 25% and the ratio of Nd to Fe is around 3.5 to 1.2 in the Nd-rich phase.

Fig.4.40 Electron micrograph of Nd15Co16Fe60Ga2B7 permanent magnet alloy at

room temperature, observed Ga2Nd phase, plain and straight and clear grain boundary, clear Nd2Fe14B phase

GaNd, Ga2Nd and Ga3Fe11Nd6 phases exist in the crystal boundary and the base

phase of Nd15Fe60Co16Ga2B7 have orthorhombic system (a=0.44 nm, b=1.13 nm, c=0.42 nm), hexagon system and square system, respectively. It can be seen from Fig. 4.41 that the crystal boundary is not clean and has mixtures of Nd-rich phase and Ga-rich phase. Its annealing temperature is 590 on the low side. The Fig.4.42 presents a related clean phase of crystal boundary of gallium containing NdFeB alloy. Ga-rich phase (in a triangle shape and embedded to base phase) has an integrated flat boundary with a black meddle but a bright sides. This alloy pos-sesses high coercive field intensity.

Fig.4.41 Electron micrograph of Nd15Co16Fe60Ga2B7 permanent magnet alloy at room temperature, observed grain boundary with boundary angle of 120º between

matrix phase of Nd2Fe14B and clean Nd2Fe14B


Fig.4.42 Electron micrograph of Nd15Co15Fe61Ga2B7 permanent magnet alloy at

room temperature, observed plain and straight grain boundary in Nd-Fe-Ga-B alloy with Ga addition, and observed complete and clear microstructure of Nd2Fe14B

4.9.3 Phase transformation of microstructure of B-rich phase at high temperature

Fig. 4.43 shows electronic microscope photograph of sintered Nd15B7Fe78 specimen at room temperature. In the figure section C represents B-rich phase (Nd1.11Fe4B4), section D represents electronic diffraction pattern of selected area of section C. We can see that: (1) the B-rich phase is clamped among base phases of Nd2Fe14B, and, is close to Nd-rich phase; (2) Crystal defect, layer dislocation and plane defect exist inside of the B-rich phase, which can be seen in the section C; (3) character of layer dislocation is of certain orientation that exhibits bright and dark striae alternately. The B-rich phase accounts for a volume fraction of 5%~8% in permanent magnetic alloy of Nd15B7Fe78 so that it may not be found easily in electronic microscope like base phase of Nd2Fe14B and Nd-rich phase. It would be found more difficultly if the specimen film be thinned excessively. Therefore, using an observable thick film under 1000 kV HVEM made the observation experiment of B-rich phase succeeded, and that becomes able to observe the layer dislocation and plane defect of striae with a certain orientation; (4) it can be seen that electronic diffraction pattern of selected area appears reciprocal point elongated phenomenon. The reciprocal point elongated phenomenon at room temperature is most possibly related to amplitude modulation of high temperature phase. We derive the reciprocal plane from (310) and (031) and calculate out its crystal lattice parameter a0 = 0.7128 nm, c0 = 0.3894 nm. This result is close to the data of reference that gave a0 = 0.7128nm, c0 = 0.3894 nm for the Nd1.11Fe4B4 (Givord Moreau, Tenaud, 1985). Iron and boron in ternary compound of Nd1.11Fe4B4 together compose a substructure in square symme-try, the same as the substructure composed of neodymium atoms. All of these struc-


tures are described using Pccn orthorhombic space group (a0 = b0 = 0.7117 nm, c0 = 0.3507 nm). All parameters of sublattice of B-rich phase (Nd1.11Fe4B4) in NdFeB permanent alloy can be computed out as follows by using electronic diffrac-tion pattern of selected area in C location of Fig. 4.43.

Name of sublattice a0 c0 Fe sublattice 0.7182 nm 0.3894 nm Nd sublattice 0.7100 nm 0.3528 nm

Fig.4.43 Electron micrograph of Nd15Fe78B7 permanent magnet alloy at room tempera-ture, observed three main phases of the alloy: matrix phase Nd2Fe14B (B), Nd-rich phase (A), B-rich phase (C), and electron diffraction pattern of select area in matrix phase (E),

electron diffraction pattern of select area in B-rich phase (D) This specimen was heated in 1000 kV electronic microscope and dynamic ob

serve the precipitation process (made video record for dynamic observation in the same time). When the temperature raised to 322 micrograph of B-rich phase changed largely compared with that at the room temperature, as shown in Fig. 4.44. Precipitated black block appears inside of B-rich phase and oxidation occurs in the dislocation direction of precipitated black block. Direction of dislocation striae in other place does not be so obvious and so clear as that at room tempera-ture. Fig.4.45 changed obviously compared with Fig.4.43. The atomic ordered temperature of Nd1.11Fe4B4 is 14K and its magnetic intensity become strongly abnormal. The reference reported that at 322 there is another abnormity in magnetic intensity (Givord, Moreau, Tenaud, 1985). This opinion is consistent with phenomenon by this dynamic observation that the micrograph of B-rich phase has big change at 332 .

4.9.4 Phase transformation of microstructure of Nd-rich filmy belt in Nd15Fe78B7 crystal boundary at high temperature

It was discovered at first that there is Nd-rich filmy belt among crystals of


Nd2Fe14B at room temperature (Fig. 4.45). It was found further that the Nd-rich filmy belt only exists among crystals of Nd2Fe14B and between B-rich phase and Nd2Fe14B phase in a serious and careful observation.

Fig.4.44 Electron micrograph of Nd15Fe78B7 permanent magnetic alloy at 322

B Matrix phase; C B-rich phase; E Diffraction pattern for B-rich phase

Fig.4.45 Electron micrograph of Nd15Fe78B7 permanent magnetic alloy at room temperature B Matrix phase of Nd2Fe14B; G bcc filmy belt (Nd-rich belt)

It was observed that the Nd-rich filmy belt is widened when the temperature


rose to 140 , as shown Fig. 4.46. By continually raising temperature to 312 (the Curie temperature of this alloy) we found the Nd-rich filmy belt is further widened, as shown in Fig. 4.47. It can be seen from Fig. 4.47 and Fig. 4.48 that the particles inside of the filmy belt are inhomogeneous but the whole filmy belt is very regular. When the temperature was raised to 600 and held for 30 min, highly dispersed precipitates appears inside of base phase but the filmy belt keeps intact without any dispersed precipitates, as shown in Fig. 4.49.

Fig.4.46 Electron micrograph of Nd15Fe78B7 permanent magnet alloy at 140 , observed Nd2Fe14B matrix phase (B) and bcc thin strip of Nd-rich phase (G)

Fig.4.47 Electron micrograph of Nd15Fe78B7 permanent magnet alloy at 312 , observed Nd2Fe14B matrix phase (B), B-rich phase (C), and electron diffraction pattern of select area

of Nd2Fe14B matrix phase (E)


Fig.4.48 Electron micrograph of Nd15Fe78B7 permanent magnet alloy heating at 312

for 30min, observed Nd2Fe14B phase (B), B-rich phase (C), and electron diffraction pattern of B-rich phase (E), heating to 500 tiny Nd2O3 phase appears; comparing to Fig.4.15, the

addition of Co increases Curie temperature and thermal stability of the alloy obviously

Fig.4.49 Electron micrograph of Nd15Fe78B7 permanent magnet alloy at 600 , observed

electron diffraction pattern of matrix phase at 500 , as shown in the figure E, multi-crystal circle appears, it indicates transition of single crystal of Nd2Fe14B into multi-crystal.

F is electron diffraction pattern of the select area of the precipitation phase

4.9.5 Phase transformation of Nd2Fe14B base phase of Nd15Fe78B7 alloy at high temperature

Section B in Fig.4.43 is the micrograph of Nd2Fe14B base phase of Nd15Fe78B7,


section E is electronic diffraction pattern of the selected area. That (uvw) * = (�101) can be calculated from (010) and (101). The crystal parameters are worked out as a0 = 0.881 nm, c0 = 1.203 nm, that is similar to a0 = 0.879 nm, c0 = 1.219 nm reported by reference (Givord, Moreau, Tenaud, 1985).

It can be seen from Fig.4.43 that there has not any elaborate texture, and neither has crystal defect inside of Nd2Fe14B at room temperature.

The dynamic observation of precipitation process did not find any precipitate below 280 by heating this specimen; at 280 the specimen began to appear precipitates; at 312 it was observed clearly that a highly dispersed precipitation phase bestrews to all visual field, and a diffraction ring of polycrystals appears on diffraction pattern. The precipitates are growing continually along with rising temperature and cause a change in crystal lattice at temperature of 500 . Nd2Fe14B precipitated at 500 , as shown in Fig. 4.48 E and Fig. 4.43 D. Precipi-tates of Nd2Fe14B crystal granule are shown in Fig. 4.49 when temperature was raised to 600 . It can be seen from the figure that the precipitates grow up rap-idly. The diffraction pattern of base phase appeared many more polycrystalline rings compared with that at 300 or 500 , as shown in section E of Fig. 4.49. The electronic diffraction pattern of selected area of precipitate phase is shown in section F of Fig. 4.49, which can be seen as polycrystalline rings entirely. This indicates that the single crystals of Nd2Fe14B in the Nd15Fe78B7 precipitated to be polycrystals and the precipitate is in a scale of tenth nm.

The massive specimen was heated in the sintering furnace as per the above mentioned heating sequence, and held at 600 for 40 min, quenched and re-magnetized, and then measured its magnetism. As the result the coercivity is not only lowered but increased slightly. It is found in dynamic observation that the dislocation existing at room temperature reduced somewhat when raising tem-perature to 140 . Most of the dislocation is reduced at 312 , and almost all of the dislocations became unobservable at 600 . This observation reveals the func-tion of clearing up stress by low temperature aging at around 600 .

It is discovered in the dynamic observation that crystal boundary among Nd2Fe14B is still clear even if heating to 500 .

For Nd15Fe85-xBx, if x=11, Nd2Fe14B descends 5%. It does not affect intrinsic coercivity, but enhance, Nd2Fe14B quantity and the saturation magnetization in-tension, properly controlling boron (B) content and alloy ingredient, properly in-creasing Fe enhance Nd2Fe14B comprehensive effective (Pan, Zhao, Li, et al, 2011).

4.9.6 Conclusions

The following conclusions are derived thorough above studies: 1. Sintered NdFeB permanent magnetic alloy with unexampled high magnetic


energy product is due to good square symmetrical phase of Nd2Fe14B (base phase) formed by adding boron and a proper amount of Nd-rich phase and B-rich phase. Where the base phase of Nd2Fe14B made the biggest contribution to magnetic energy product. NdFeB alloy has a big magnetic anisotropy thorough function crystal field because Nd3+ possesses track magnetic moment. Furthermore, split of crystal field of rare earth 4 f energy grade is the source of intrinsic magnetic ani-sotropy in Nd2Fe14B. All of magnetic moments of neodymium and iron are paral-lel to c axis and the magnetic moment array is ferromagnetic. However, the square symmetry is broken because of raising temperature. The precipitate has been observed at 280 , and that the original non-defective base phase of Nd2Fe14B appears defects due to precipitation of tiny polycrystals. These defects become antimagnetic nucleation center that results in degradation of coercivity.

2. One of the functions by aging for 1 h at 600 is to eliminate internal stress. It is observed that dislocations disappeared along with rising of temperature in dynamic observation.

3. bcc filmy belt of Nd-rich phase surround homogeneously Nd2Fe14B phase, Nd-rich phase and B-rich phase in a specimen being aging for 1 h at 600 . It is observed in the dynamic observation that the bcc filmy belt is widened propor-tionally along with rising of temperature, meanwhile the intensity of coercive force degrade with rising of temperature and the coercivity degrade to 0 at 312 . The bcc filmy belt widens regularly and proportionally at rising tempera-ture up to 600 . And that there are polycrystalline precipitates fulfilling the field of vision in a size of teens nm on Nd2Fe14B but the bcc filmy belt appears no pre-cipitate. This fact indicates that bcc filmy belt make no contribution to the coer-civity and may not belong to hard magnetic phase.

4. B-rich phase in sintered NdFeB permanent magnetic alloy has crystal defects at 25 . There are two type of lattices: ferrous sublattice (a0 = 0.7128nm, c = 0.3894nm) and neodymium sublattice (a0 = 0.7100nm, c = 0.3528nm).

4.10 In Situ and Dynamic Observation on High Temperature Phase Transformation and Magnetism of Nd16Fe77B7 Permanent Magnetic Alloy

NdFeB permanent magnet has been studied widely (Sagawa, Fujimura, Togawa, et al, 1984) since it was found in 1983. Using cobalt to substitute part of iron can raise Curie temperature of the tetragonal phase of Nd2Fe14B considerably, and can improve thermal stability of the alloy (Sagawa, Fujimura, Yamamoto, et al, 1984; Matsuura, Hirosawa, Yamamoto, et al, 1985; Arai, Shibata, 1985; Hirosawa, To-kuhara, Yamamoto, et al, 1987; Fuerst, Herbst, 1988). Researches indicated that NdFeCoB magnet mainly consisted of quadrilateral phase Nd2(Fe, Co)14B and phase of crystal interface. The phase of crystal interface contains Nd-rich phase


and Laves phase of Nd(Fe, Co)20 (Arai, Shibata, 1985), and there was report that phase of Nd20(Fe, Co)80 was found as well (Koestler, Ramesh, Echer, et al, 1989). But B-rich phase seems hard to be found. Nevertheless, the remanence is reduced by adding cobalt into the tetragonal phase, and that appearance of magnetically soft phase results in dramatically degradation of the coercivity.

Microstructure has important influence on magnetic features, especially on the coercivity, of sintered NdFeB magnet. This section introduces a dynamic observa-tion on microstructure change of Nd16Fe69Co8B7, with optimal magnetic features, at temperature from 25 to 700 . Though the microstructure of NdFeCoB mag-net has been studied detailedly (Koestler, Ramesh, Echer, et al, 1989), there is no report about dynamic observation. A comparison is made between the dynamic observation result of NdFeB and that of NdFeCoB.

4.10.1 Samples preparation process and experimental method

Neodymium, iron, cobalt and ferroboron with purity above 99.5% was arc melted in vacuum crucible (1.5×10�2 Pa) to get ingot of the alloy, the ingot was milled to be powder of 3-5 m which was oriented in magnetic field of 1.5 T and was pressed out in axial direction at 2T/cm2, and then the pressed powder was sintered at 1100 for 1 h and edged at 550 . Magnetic measurement indicated that this magnetic possesses outstanding features: its remanence, coercivity and magnetic energy integral were 1.28 T, 589.0 kA/m and 302.4 kJ/m3.

The sintered alloy was sliced to lamella of 0.25 mm in direction perpendicular to c axis of the magnet for observation under HVEM, the lamella was polished mechanically thinned by ionic beam to prepare film sample which is perpendicu-lar to the axis of magnetic field. The observation was conducted under JEM-1000 HVEM which has temperature controlling unit. Operation voltage was 1000 V.

4.10.2 The in situ and dynamic observation on nanometer mi-crostructure and high temperature phase transformation

The study investigated the behaviors of additional cobalt and the Nd-rich phase at different temperatures, which benefits the understanding of the temperature de-pendence of the magnetic properties.

4.10.2.1 High temperature phase transformation of Nd2(Fe, Co)14B

Nd2(Fe, Co)14B crystal granule presents as integral crystals at room temperature, which almost exists no defect. Crystal lattice constants of its quadrilateral phase can be calculated by X-ray diffraction pattern as a0=0.882 nm, c=1.218 nm. But in some crystals a few big and many tiny globose impurities distributed in the matrix could be observed clearly, seeing Fig. 4.50(a). These impurities phases were de-


marcated as bcc structure with lattice constant a0=1.086 nm by electronic diffrac-tion pattern (Fig. 4.50(b)). They were analyzed as Nd2O3. Fidler, et al also found these impurities in the NdFeB magnet with addition of aluminum and determined them as non-magnetic hexagonal phase of Nd2O3 (Fidler, Knoch, Kronmüllera, et al, 1989). By analyzing their microstructure authors found that there was a non-crystalline phase (Fig. 4.50(c)) between impurities and the matrix. In NdFeCoB magnet the most important ferromagnetic phase is the tetragonal phase Nd2(Fe, Co)14B which takes on magnetically hard characteristic of magnet, and appear-ance of the non-ferromagnetic phase and the non-crystalline phase will reduce the remanence of the magnet. In addition, nucleation in non-crystalline phase though reverse domain will make coercivity of the magnet degraded sharply.

Fig.4.50 Electron micrograph of Nd2(Fe, Co)14B at room temperature

(a) Big and small black blocks are spheral impurity embedded in matrix phase; (b) Amorphous phase be-

tween matrix phase and spheral impurity; (c) Diffraction pattern of the amorphous phase in (b)

In heating process when temperature was below 400 the microstructure of

the quadrilateral phase had no remarkable change; when temperature was higher than 400 some tiny precipitates started to appear in the quadrilateral phase; and when reached 500 the precipitates grew like needle which presented to be per-


pendicular with each other and appeared evenly and largely on the matrix, as shown as Fig. 4.51; when temperature was further raised to 700 the quadrilat-eral symmetrical the matrix was demolished completely and its microstructure is shown as Fig. 4.52.

Fig.4.51 Electron micrograph of Nd2(Fe, Co)14B at 500 , observed phase transition at

500 , with heat activation twin crystal, dislocation sub-structure appear

Fig.4.52 Electron micrograph of Nd2(Fe, Co)14B at 700 , observed destruction of matrix phase

Author had conducted the in situ and dynamic observation on Nd16Fe77B7 mag-

net by HVEM and found that the precipitates started to be appear in Nd2Fe14B crystal granule at 280 (Pan, Liu, Luo, 1990). Grössinger, et al had measured that anisotropy field HA and coercivity field Hc changed with variation of tem-perature (Grössinger, Krewenka, Kirchmayr, et al, 1987). The HA and Hc were degraded rapidly along with rising of temperature, degradation of HA was fastest at around 280 which is the temperature corresponding to the appearance of the precipitates. It can be seen that the precipitates of the tetragonal phase is the main reason for degradation of HA and Hc and completed demolition of quadrilateral symmetry will lead to degradation of HA and Hc to zero.


The precipitates started to appear at 400 in this work and this temperature was about 100 higher than the precipitation temperature of samples without cobalt. It can be seen that addition of cobalt can improve thermal stability of the tetragonal phase. Deeply study on composition analysis and structure ascertain-ment of precipitated phase would be favorable for developing application of the magnet and understanding of the relation between magnetic features and change of temperature.

4.10.2.2 The in situ and dynamic observation of crystal interface

The crystal interface phases include Nd-rich phase and a phase of undefined structure. These phases can be observed in crystal interface lamella and interfaces among three or four granules of crystals, seeing in Fig. 4.53(a) and Fig. 4.53 (b). The Nd-rich phase at the interface of three granules has fcc structure with lattice constant a = 0.55. The other phase has not been defined because of the compli-cated electronic diffraction pattern, as shown as Fig. 4.53(b). Many studies indi-cated that magnetically soft phase of Nd(Fe, Co)2 existed in NdFeCoB magnet, but that whether the phase of undefined structure is belong to the magnetically soft phase still needs further study to determine it (Mizoguchi, Sakai, Niu, et al, 1986).

Fig.4.53 The structure of two kinds of crystal grain boundary in Nd16Fe69Co8B7

permanent magnet alloy at room temperature (a) Thin layer of the boundary of Nd-rich between two grains of matrix phase Nd2(Fe, Co)14B;

(b) A triangle grain boundary of Nd-rich enclosed by three Nd2(Fe, Co)14B grains The phase in crystal interface has significant influence on coercivity. The Nd-

rich lamella of non-magnetism acts as pinning function to move of the domain wall and so that will enhance the coercivity. But the magnetically soft phase at triangle crystal boundary would act as nucleation center of the reverse domain to produce demagnetization field so that to degrade coercivity sharply. According to the importance of the phase of crystal interface to the coercivity it is necessary to


observe them at a high temperature. The Nd-rich phase and crystal interface lamella phase had no remarkable varia-

tion during the whole process of rising temperature. Microstructure of the phase of crystal interface with undefined structure had not remarkable change when temperature raised from 25 to 600 . And when temperature was raised to 700 the diffraction pattern of its mircotexture was found to have a change. A in Fig. 4.54 is the microstructure of phase of undefined structure at 700 , B is the diffraction pattern which appears as multi-crystalline ring. When sample was held for 20 min at 700 the diffraction pattern consisted of a non-crystalline diffusing ring and some diffraction spots, as shown in C. This change process still waits to be studied referring to the result obtained at present. It can be seen from the other group of relationship curve between aging temperature and coercivity of Nd16Fe69Co8B7 alloy (in section 4.10.7) that coercivity of the alloy after sintering was not high, was measured as 379 kA/m; if try to draw a relationship curve be-tween aging temperature and coercivity it can be seen that the intrinsic coercivity of the alloy was the highest as 589 kA/m at 550 ; the intrinsic coercivity de-graded at 700 and enhanced slightly at 900 . Ternary NdFeB alloy appeared coercivity peak at 600 , which is considered to attribute to the dissolving of the unstable magnetically soft phase of NdFeB alloy by heat treatment at 600 , and this process can improve coercivity of the alloy (Schneider, Hening, Missell, et al, 1990; Pan, Ma, Ping, et al, 1991; Pan, Jin, 1990; Pan, Li, Li, et al, 1989; Pan, Zhao, Ma, 1988). In comparison, experiment by author found the aforementioned process appeared at 550 , that is to say, the magnetically soft phase of the alloy with cobalt dissolved at 550 . Moreover, the coercivity degraded at 700 con-siderably, that had not been studied and discussed in reports before. Observation with electronic microscope found that the diffraction pattern of the matrix was

Fig.4.54 Electron micrograph of triangle grain boundary in Nd16Fe69Co8B7

permanent magnet alloy at 700 A Uncertain phase; B Diffraction pattern of uncertain phase where multi-crystal circle is observed;

C A diffusion circle of amorphous and a few of diffraction spots of the uncertain phase changed from its electron diffraction pattern at 700


multi-crystal ring at 700 , but after holding at 700 for 20 min the diffraction pat-tern changed from the multi-crystal ring to a non-crystalline diffusion ring and same diffraction spots. Accordingly, the coercivity degraded correspondingly at 700 .

4.10.3 Function of cobalt in NdFeCoB alloy

It was known from measurement result of Curie temperature that using some co-balt to replace part of iron can raise Curie temperature of the alloy. The reason to raise Curie temperature is that the exchange function of Co-Co and Co-Fe is stronger than that of Fe-Fe, and Co occupies j2 and k2 crystal sites preferentially, which can improve negative exchange between j2 - k2 sites so that to enhance positive exchange function and weaken negative exchange function.

The Nd16Fe77B7 permanent magnetic alloy mainly consists of the matrix phase of Nd2Fe14B, Nd-rich phase and B-rich phase (Nd1+�Fe4B4), among of them the phase making the biggest contribution is Nd2Fe14B, accounting for 80%-85% of total volume. And Nd16Fe77B7 has the biggest magnetic anisotropy by action of crystalline field because Nd3+ has track magnetic torque; and the crystalline split of 4f energy grade of the rear earth element is the source of intrinsic magnetic anisotropy of Nd2Fe14B (Arai, Shibata, 1985). When using some cobalt to replace part of iron the matrix of the alloy becomes Nd2(Fe, Co) 14B. To study cobalt the effect of atom of cobalt entering into crystal lattice on Curie temperature Tc a mi-cro-area composition electronic probe analysis was made on the alloy of Nd16Fe69Co8B7. Analysis result indicated that interior of crystal of alloy Nd2(Fe, Co)14B, Fe=77.3%(at.), Nd=15.2%(at.), Co=7.5%(at.); in comparison with micro-area electronic probe analysis a conclusion can be achieved that enhancing cobalt content in the matrix phase results in rising of Curie temperature Tc. Nd16(Fe1-x

Cox)77B7 (x=0.3), Tc reaches 553 (Pan, Li, Xu, Ma, 2011). Generally, by aforementioned experiment we obtained precedentless valuable

information about the alloy in the in situ and dynamic observation that provided direct gist for further study of microstructure. In addition obtaining of knowledge about microstructure variation of phases around crystal interface at high tempera-ture would vail to clarify the mechanism between coercivity and microstructure.

4.10.4 Magnetic characteristic measurement result and analysis

CL6-1 magnetic parameter measurer was use to measure magnetism of Nd16Fe69Co8B7 alloy after magnetization saturation. The result measured was: remnant magnetic induction intensity is 1.28T, the intrinsic coercivity is 589.0 kA/m, the maximum magnetic energy product (BH)max = 302.4 kJ/m3. It can be seen from the result that this alloy has excellent magnetic performance. Its de-


magnetization curve is shown in Fig. 4.55.

Fig.4.55 Demagnetization curve of Nd15Fe69Co8B7 alloy

4.10.5 Curie temperature measurement result

Curie temperature of ternary alloy of Nd16Fe77B7 is 312 , after adding cobalt the Curie temperature is raised remarkably to be Tc = 510 , as shown in Fig. 4.56.

Fig.4.56 Measurement curve of Curie temperature of Nd16Fe69Co8B7

4.10.6 Phase analysis by X-ray diffraction, lattice constant and cell volume

X-ray diffraction pattern is shown in Fig. 4.57, lattice constant and cell volume are shown in Table 4.19.

It can be seen from Table 4.19 that a, c, V increased after adding cobalt to re-place part of iron.


Fig.4.57 The X-ray diffraction pattern of Nd16Fe69Co8B7 (the number indicates the crys-

tal plane index of Nd2(Fe, Co)14B and tetragonal phase) Nd�rich phase; B�rich phase; Laves phase

Table 4.19 Lattice constants of the tetragonal phase and cell volume of alloys

Nd16Fe77B7 and Nd16Fe69Co8B7

Alloy a/nm c/nm c/a V/nm3

Nd16Fe77B7 0.8785 1.2188 1.387 0.9406

Nd16Fe69Co8B7 0.880 1.270 1.38 0.9436

4.10.7 Relationship between aging temperature and coercivity of Nd16Fe69Co8B7

The intrinsic coercivity of Nd16Fe69Co8B7 alloy was measured as 379 kA/m only after sintered at 1100 , and after edging at 550 its intrinsic coercivity became 589 kA/m. The curve for relationship between aging temperature and coercivity was made at the same time, as shown as in Fig. 4.58.

Fig.4.58 Relationship of coercivity of Nd16Fe69Co8B7 permanent magnet alloy vs Aging temperature


4.10.8 Conclusions

Conclusion was achieved based on the aforementioned studies: 1. The in situ and dynamic observation on Nd16Fe69Co8B7 alloy through HVEM

found that using cobalt to replace part of iron altered modality of the matrix phase of ternary NdFeB alloy, i.e., some globose Nd2O3 impurities appeared in granule of Nd2(Fe, Co) 14B. When temperature was raised to 400 some tiny precipitates appeared in the matrix phase, when temperature reached 500 the precipitates grew up to be acerose and appeared largely. Nd-rich phase existed in crystal inter-face, which had no remarkable change in the process of raising temperature. Structure of crystalline state transferred at 700 . It can be seen from the diffrac-tion pattern that the patter of original multi-crystal ring became that of a non-crystalline diffusion ring and some diffraction spots at 700 and holding for 20 min.

2. The dynamic observation found that temperature the precipitates appeared in the matrix phase precipitates was raised 120 in comparison with ternary NdFeB alloy, i.e., the thermal stability was heightened 120 by addition of cobalt. This result is just responding to the measuring result that Curie temperature was heightened in 150 by addition of cobalt to replace part of iron.

3. X-ray diffraction and metallographic study result indicate that lattice con-stants a, c were increased and crystal cell volume was augmented by adding cobalt to replace part of iron.

4.11 Analysis on Lamella Phase of Grain Boundary in Micro-structure of NdFeB Permanent Magnetic Alloy

In resent years study on microstructure and coercivity of NdFeB alloy has achieved great progress, but for mechanisms of nucleation and pinning there still needs discussion. K. Hiraga, et al considered that there existed lamella phase of bcc in crystal interface surrounding Nd2Fe14B granule (Pan, Zhao, Ma, 1988). And some others considered that this bcc phase was formed in preparation process (Pan, Liu, Luo, 1990). In order to have a clearer understanding of microstructure around the grain interface, authors made systematically observation and analysis using HVEM. This section is not to explain these phenomena but is to issue information obtained in experiments for further discussion. 4.11.1 Experimental method

Two permanent magnets of Nd16Fe77B7 and Nd16Fe57Co16Ga4B7 were prepared by


normal powder metallurgy method. The magnet was thinned by machine and ion to prepare samples for electronic microscope. Rising of temperature and oxidation should be avoided by all efforts. The in situ and dynamic observation on prepared samples was conducted under JEM-1000 HVEM.

4.11.2 Magnetism measurement

Magnetic characteristic of sample of Nd16Fe77B7 was: Br = 1.22 T, Hc = 824 kA/m, (BH)max = 280 kJ/m3; that of Nd16Fe57Co16Ga4B7 was: Br = 1.1 T, Hc = 766.4 kA/m, (BH)max = 214.8 kJ/m3.

4.11.3 Analysis on result of the in situ and dynamic observation of samples

The matrix phase of Nd2Fe14B, B-rich phase of Nd1+�Fe4B4 (�= 0.1-0.3) and Nd-rich phase were observed under JEM-1000 HVEM. In addition a bcc lamella phase was found in crystal interface. The following is to analyze state of bcc la-mella in the crystal interface and its order of change with temperature. Fig.4.59 is the microstructure around boundary of Nd2Fe14B. Where A is the matrix phase, C is lamella in crystal interface; Fig. 4.60 is the other figure of the lamella texture in crystal interface, i.e., a clear lamella texture existed in crystal interface between two Nd-rich phases, and this type of lamella texture extended to interior of crystal granule in part of area. Author considered that this lamella belt might be formed in crystal interface and other interfaces (there was lamellar texture between in-cludes and the matrix phase, as shown in Fig.4.59), this type of filmy belt may extend to interior of the matrix phase of Nd2Fe14B if there are some other types of plane defects near crystal interface and connection locals of the interfaces. If the

Fig.4.59 bcc thin layer and impurity in the grain boundary of Nd15Fe77B8 and

bcc thin layer on the boundary of matrix body A Matrix phase; C Thin layer


filmy belt were formed by oxidation in preparation process it was difficult to extend to interior of the matrix phase evenly. Moreover, this filmy belt existed widely in samples and it was evenly in width. Therefore, author considered that this belt was impossible to form in preparation process, but it was inherent in magnet after annealing process.

Fig.4.60 bcc thin layer between Nd15Fe77B8 alloy and Nd-rich phase

A Nd-rich phase; B Matrix phase; C Thin layer By analysis this filmy belt is Nd-rich phase with a thickness of 20nm, which

appeared in crystal interface and cross corner after disappearance of Nd-rich phase by annealing at 580-635 ; it belongs to fcc crystal structure with lattice constant 0.56 nm; the content of neodymium in the lamellar belt is about 85%.

This lamellar Nd-rich phase acts important function for magnetic hardening of sintered NdFeB alloy. If Nd-rich phase is insufficient in the alloy the coercivity would be very low (though Br is high) and the general magnetic performance would not be good.

Fig.4.61 is the magnified photo of the filmy belt in Fig.4.60. The figure of the filmy belt can be seen more clearly that it is not an integral crystal but are many small interfaces existed which can be taken as microcrystal. Fig.4.62 is Nd-rich phase in intersectant locus of three filmy belts, equivalent to a trifurcate interface, and each interface is occupied by bcc lamella. As shown in Fig.4.63 that the in-cluded angle among three filmy belts is about 1201. Among them two are linked with each other and the other one is independent, they are Nd-rich filmy belts between the matrix phases.


Fig.4.61 Amplified graph of Fig. 4.60: thin layer of Nd15Fe77B8 alloy

Fig.4.62 The pattern of the join point of three bcc thin layers in Nd15Fe77B8 alloy, at the

join point there is Nd-rich phase A, B Matrix phase; C Nd-rich phase, grain boundary and between grain boundaries

Fig. 4.64 is the lamellar phase in crystal interface observed 280 . Fig. 4.65 the

filmy belt between includes and the matrix phase. There was uneven microstruc-ture in the filmy belt. The filmy belt was also observed in sample of Nd16Fe57Co16Ga4B7, as shown in Fig. 4.66. The formation of filmy belt can pre-


vent growth up of crystal granule, and crystal interface have big resistant for movement of the domain wall. In addition, the belts have low anisotropy thus they may act as effective pinning stand to improve coercivity of material.

Fig.4.63 Electron micrograph of Nd15Fe77B8 permanent magnet alloy at room temperature

B Matrix phase Nd2Fe14B; C bcc thin layer

Fig.4.64 Filmy belt between the matrix phase and inclusions of Nd15Fe78B7 alloy

observed at 280 B Matrix phase; C Filmy belt

Fig.4.65 The thin layer between matrix and impurity in Nd15Fe77B8 alloy (C)

B Matrix phase


Fig.4.66 The thin layer structure in Nd16Co16Fe57Ga4B7 specimen B Matrix phase; C Thin layer

What is showed in Fig. 4.62 belongs to grain boundary of multiple phases, the

Nd-rich globules in crystal interface of multiple phases congregated to be larger Nd-rich conglomerations, and the larger ones existed in cross corner of three crys-tal granules of Nd2Fe14B. When annealing at 630 neodymium in Nd-rich con-glomerations extended along crystal interface to the interface so that led to ho-mogenization. The thickness of lamella of Nd-rich phase in the interface is differ-ent because of different technical system.

Fig. 4.60 shows magnified of ashine Nd-rich filmy belt. It can be seen that there are small granules in the belt. For formation of the filmy belt many articles have presented different opinions as per their experiments. Some thought that is the broken off of the Nd-rich phase in process preparing the filmy sample for electronic microscope, possibility of this opinion is not excluded because the sample for observation under transmission electronic microscope must be very thin. In the experiment JEM-1000 HVEM was used with super-high voltage for accelerating voltage, and thus with high transmission rate can observe compara-tively thicker film. One of characteristics of HVEM is to observe filmy sample with thickness below 100 nm. As shown in Fig. 4.67 to Fig. 4.70, Nd-rich filmy belt in the interface mention before has ascertain thickness and this ashine Nd-rich lamella was had extended in 5 times at 312 , its magnified photo is shown as in Fig. 4.70. It can be seen that there were Nd-rich tiny particles from several nm to more than ten nm. Samples with the same composition but with different aging program were designed in order to describe the contribution of Nd-rich


filmy belt. The result revealed that the sample with such Nd-rich filmy belt in microstructure had higher coercivity in measurement. This type of Nd-rich filmy belt has fcc structure, its atom fraction is Nd 75%, and iron in balance. Its elec-tronic diffraction photo is shown as in Fig. 4.71, and its photoelectron energy spectrum is as Fig. 4.72.

Fig.4.67 Electron micrograph of broadened thin layer in Nd15Fe77B8 alloy observed at 450

B Matrix phase; C Thin layer (Nd-rich phase)

Fig.4.68 Electron micrograph of broadened thin layer in Nd15Fe77B8 alloy observed at 600

B Matrix phase; C Thin layer (Nd-rich phase)


Fig.4.69 Electron micrograph of Nd15Fe77B8 alloy observed at room temperature B Matrix phase; C Thin layer (Nd-rich phase)

Fig.4.70 Electron micrograph of Nd15Fe77B8 alloy at 312 B Matrix phase; C Thin layer (Nd-rich phase)


Fig.4.71 Electron diffraction of Nd-rich phase

Fig.4.72 Photoelectron energy analysis of NdFeCoGaB This type of lamella in crystal interface has almost disappeared in NdFeB alloy

more than ternary system, as shown in Fig.4.66 for alloy of Nd16Co16Fe57Ga4B7. The Ga-rich phase (Ga2Nd, GaNd, Ga3Fe11Nd6) of this type in crystal interface of multi-phases congregated herein. In crystal interface of alloys exceeding ternary elements it is hard to find three ashine lamellas distributed in 1201. Nd-rich la-mella with thickness of 20-30 nm and fcc structure of lattice constant 0.56 nm, which has contributes to coercivity of the alloy.

There are several types of Nd-rich phases. But compared with analysis of phase diagram Nd-rich phases can be taken as Nd-Fe binary eutectic alloy (Zhou, Dong,


1999). Phase diagram indicates that the eutectic temperature of Nd-Fe binary al-loy is 640 , that of ternary alloy of NdFeB is 655 and Nd-rich phase in cross corner of crystal interface became liquid phase at 640 . In experiment the edg-ing program technique, with important effect on coercivity of the alloy, were dif-ferent with different alloy composition, not all at 600 . Nd-rich crystal interface for NdFeB alloy with element more than three was not pure Nd-rich phase yet. To bring the multi-phases to be homogenized it needs to look at the combination ability of the adding element with iron. Determination of appropriate aging pro-gram needs to refer to phase diagram and a lot of experiments of different aging times and aging temperature.

When cooling sintered NdFeB alloy there was a deep eutectic tendency to keep the liquid out-of-order texture to room temperature because the eutectic tempera-ture of ternary NdFeB elements have big different with melting points of iron and neodymium. Neodymium presents Nd-rich phase with different forms, different uniformities and different types, as shown as single phase interface in Fig. 4.63 and Fig. 4.64, and multi-phases interface of multi-components in Fig. 4.66 be-cause of cooling velocity (quenching or cool with furnace) of sintered alloy. Owing to deep eutectic (out-of-order deep eutectic) Nd-phase in crystal interface goes deep into interior of Nd2Fe14B, Nd2(Fe, Co)14B, Nd2(Fe, Nb)14B and Nd2(Fe, Ga)14B. To obtain clean Nd-rich phase and heighten coercivity doubly sometimes need multigrade aging. This explains why the sintering program for NdFeB alloy will be two-stage aging, at 900 and 600 , respectively, and some type of alloy does not need aging at 900 . Therefore, author found in experiment that some-times quadrate degree of demagnetization curve of one stage aging is better than that of two-stage aging. Generally, the more are the elements adulterated in NdFeB alloy the more is the effort made to search after for heat treatment pro-gram and melting program. To determine these programs repetitious experiments are necessary. Determination of heat treatment temperature is related to types and states of row materials, for example, to use dysprosium or Dy-Fe alloy, use boron or B-Fe alloy, use metallic niobium or Nb-Fe alloy. Binary alloy phase diagram needs to be utilized to guide melting, sintering and heat treatment programs. Un-der the correct programs Nd-rich phase can exert verily its function in assist sin-tering, alloy densification and heightening coercivity.

The type and figure of Nd-rich phase are related to neodymium content in composition. To make the NdFeB alloy of high magnetic energy product just needs to act the above mentioned functions really without excessive neodymium and heighten proportion of magnetic phase of Nd2Fe14B as much as possible.

Neodymium in NdFeB alloy needs to supply B-rich phase besides supplying Nd in Nd2Fe14B and Nd in Nd-rich phase, that is, Nd in Nd1+�Fe4B4 (�=0.1-0.3). The B-rich phase is nonmagnetic phase, which is hard to be observed ordinarily.


Its forms are different because there are not simple multiple relation between cFe-x and cNd. It is the type of a one-dimension with incommensurate structure varying composition. This compound is similar to Nowotny phase such as MnSi2-x, etc., and is a special crystalline structure, or being called chimney-ladder structure. It belongs to paramagnetic phase so that its study has been neglected by researchers working on magnetic materials. Experiments found that B-rich phase had lattice aberrance at 322 , and the layer dislocation in B-rich phase recovered the figure as at room temperature to 500 , but at that temperature precipitates (Nd2O3) ap-peared in Nd2Fe14B and Nd-rich phases. The B-rich phase is paramagnetic phase at room temperature; its Curie temperature is 13K but it still stable at 500 . The purpose to add boron is to form ferromagnetic phase of Nd2Fe14B, excessive part is to form isolated B-rich phase, as shown in Fig. 4.73. B-rich phase can be pinning point on domain wall of reverse magnetization nucleus.

Fig. 4.73 B-rich phase in Nd15Fe77B8 alloy B Matrix phase; C B-rich phase

The structure of the deposits changed from amorphous structure to microcrys-talline as the metal ion concentration ratio increased. When the metal ion concen-tration ratio was within 0.3-0.5, the cellular size of the deposit surface decreased gradually with the metal ion concentration ratio increase. When the metal ion concentration ratio increased up to 0.6, the cellular size of the deposit surface coarsened and presented a microcrystalline morphology, indicated that there was a critical value within 0.5 and 0.6, and when it was smaller than the value, the structure presented fine amorphous cellular morphology, on the contrary, the de-posit structure began to change from amorphous to microcrystalline morphology (Yuan, Cao, Feng, et al, 2010).


4.12 Quick Quenched NdFeB Permanent Magnetic Alloy

This section introduce the following study project: (1) use cheap Nd-Pr rare earth metal compound as raw material to reduce material cost; (2) seek for temperature and program to change non-crystal to microcrystal; (3) use 1000 kV HVEM to conduct dynamic observation and achieve fruition.

4.12.1 Sample preparation technique and experimental method

Inconsumable arc furnace was used to melt alloys of (Nd, Pr)12.5Fe81B6.5, (Nd, Pr)13.5

Fe80B6.5, (Nd, Pr)14.5Fe79B6.5 and Nd13.5Fe81B5.5. Composition of raw material (wt.) is: (Nd, Pr) mix rare earth metal: Nd 94.22%, Pr 3.68%, La 0.33%, Ce 0.32%; metallic Nd: Nd 99.69%, Pr 0.1%, Sm 0.05%, Ce 0.07%, La 0.09%; Fe: 99.8%; B, used BFe alloy with B 18.7%. In melting process draw vacuum at first, then refill argon gas for protection, to ensure homogeneous turning over repeatedly and use good water-cooling condition to obtain good columnar crystal. Afterwards, use single-roller method to quick quenched alloy thin belt in high purity argon. Sur-face speed of cupper roller was 30 m/s, belt width was 2-3 mm, thickness of filmy belt was 20-30 m. Annealing was conducted at 4×10-3 Pa in vacuum heat treat-ment furnace. Quick quenched sample was identified as non-crystalline state by means of X-ray diffraction and heat treatment temperature of quick quenched was 705-710 . Comminuted magnetic powder was bond by epoxy resin and was press-molded. Observation was carried out by JEM-1000 HVEM at room tem-perature and high temperature: sample film was enclosed into JEM-1000 side-insert heating dais, voltage was 1000 kV, and vacuum degree was 266×10-5 Pa (Pan, Ping, Liu, et al, 2003; Liu, Pan, Luo, et al, 1991).

4.12.2 Measurement result of quick quenched magnet

It can be seen that using cheap (Nd, Pr) rear earth mixing metal as raw material to produce quick quenched magnet is feasible, the magnetic performance of magnet made of (Nd, Pr) was almost the same as that made of pure neodymium (Table 4.20). Thereby production cost of NdFeB magnet can be reduced considerably.

Table 4.20 Magnetic measuring result of quick quenched magnet No. Alloy Br/T iHc/kA�m�1 bHc/kA�m�1 (BH)max/kJ�m�3 1 (Nd, Pr)12.5Fe81B6.5 0.66 811 405 66 2 (Nd, Pr)13.5Fe80B6.5 0.61 1098 366 54 3 (Nd, Pr)14.5Fe79B6.5 0.58 1114 342 50 4 Nd13.9Fe80.2B5.9 0.58 1654 52 5 Nd13.7Fe80.6B5.7 0.68 848 63


4.12.3 Relationship between crystallization temperature and co-ercivity

It can be seen from Table 4.21 that crystallization is good at 700 for short time, of which mechanism can be understood as: at low temperature energy for nuclea-tion is inadequate and hereupon presents as growing up of original crystals; if at high temperature heat activation causes that difference in free enthalpy between old and new phases is below zero, and thus phase transformation occurs sponta-neously, and driving force of nucleation is generated to make uncrystallized part nucleated generally. That result in formation of micro-crystals.

Table 4.21 Relationship between crystallization temperature and coercivity

No. Crystallization temperature/ 'max 'r Hc/kA�m�1

1 200 65.8 25 1265

2 300 82.6 44.1 2056

3 500 76.2 40.0 1595

4 700 81.25 55.0 4177

5 800 87.2 50.7 2439

6 900 88.1 31.91 755

4.12.4 Microstructure at room temperature

Quick quenched NdFeB alloy and sintered NdFeB alloy have different micro-structures. B-rich phase was not observed but in Nd2Fe14B phase internal structure of crystal granules was found to be integral in quick quenched NdFeB alloy by transmission microscope, the neighboring crystal interface was also found to be integrated, thus the alloy presented Nd2Fe14B single phase with equiaxed fine granules texture.

4.12.5 The in situ and dynamic observation on the non-crystal sample transferring to micro-crystal by HVEM

Non-crystalline sample after quick quenching, i.e., the sample without micro-crysals, was selected and was prepared to be film by the method aforementioned for observation by transmission electronic microscope. The observation by elec-tronic microscope was to observe the process of precipitation and growth of crys-tals. There was not any change of the sample to be found in observation from


room temperature to 400 , and observation result indicated that the microstruc-ture was still noncrystalline from diffraction of selected area under electronic mi-croscope. But at 420 the observation found that tiny crystals formed in the area originally having no precipitates, that indicated that microcrystals (crystallite) appeared. When raising temperature from 420 to 600 the observation found that the microcrystals grew up gradually and became crystals at 600-700 refer-ring to diffraction of the original selected area, that there was no non-crystalline area indicated that the sample became crystallized completely.

4.12.6 Conclusions

The following conclusions was obtained based on studies aforementioned: 1. Result of magnetic measurement indicated that the magnetic performance of

bond magnet, made of quick quenched magnetic material with (Nd, Pr) mixing rare earth metals to prepare, was equivalent to that of quich quenched magnet made of ternary NdFeB alloy, that can reduce price of magnet considerably be-cause the price of (Nd, Pr) mixing rare earth metals is cheaper than price of pure neodymium.

2. The experiment discovered that heat treatment at 700 can heighten coer-civity of quick quenched magnet considerably.

3. Dynamic observation found by 1000 kV HVEM that microcrystals appeared in quick quenched non-crystalline sample at 420 , turned to be crystals com-pletely at 600-700 , and that there was no non-crystals area indicated the trans-fer was completed.

4. Studies indicated that quenching speed (cooling speed) in quick quenching has important effect on magnetic performance of magnetic. For alloy of composi-tion Nd13.5�14B5.5�6Fein balance quenching speed of quick quench of 20-30m/s is most appropriate. The quenching speed would be changed if alloy composition has big change.

4.13 Stability of the Rare Earth Permanent Magnetic Alloy

The stabilities of rare earth permanent magnetic alloy include temperature stabil-ity, time stability, mechanism stability, magnetic stability, radiation stability, chemical stability, etc. These stabilities are as important as the magnetic proper-ties of rare earth permanent magnetic alloy.

4.13.1 Stability on temperature

Change of outside temperature (environment temperature) causes variety in mag-netic performance of permanent magnetic alloy. The magnetic capability of the


rare earth permanent magnet will not degrade by using at or below room tempera-ture. When used at high temperature the magnetic capability will degrade or ap-pear irreversible loss. Nd15Fe77B8 is used stably at 80 or below. However, the requirement of permanent magnet being used in motors for aviation and space-flight must be working stably at an environment of 200 . High capacity rare earth permanent magnet is used in magnetic force damper, magnetic force sensor and power apparatus of high rotation speed for spaceflight works in mediums of liquid hydrogen and liquid oxygen. Accordingly, inspection at normal temperature will not satisfy the requirements because it works at lower temperature. Therefore, requirement for temperature stability has to be necessary for the rare earth perma-nent magnetic alloy (Pan, Ma, Li, 1993; Pan, Ping, Liu, et al, 2003; Pan, Zhao, Ma, 1988; Pan, Li, Li, et al, 1989; Pan, Li, 2000; Pan, Chen, Liu, et al, 1994; Tang, Feng, Luo, Pan, 1994).

The Fig. 4.74 shows that remanence B in open circuit of permanent magnetic materials vary with the change of temperature. The temperature stability of per-manent magnetic alloy can be described use the fore parameters (Zhou, Dong, 1999; Pan, 2011):

1 0 1T

0 0

( ) ( )100% 100%

( ) ( )B T B T Bh

B T B T� �

� � � � (4.9)

0 0irr

0

( ) ( )100%

( )B T B T

hB T

/ �� (4.10)

1 0rev

0

( ) ( )100%

( )B T B T

hB T

/��

/ (4.11)

1 0

0

( ) ( )100%

( )dB T B T

B T T�

/��

/ (4.12)

where hT is total loss of remanence B in open circuit (flux in open circuit); hirr is irreversible loss of flux B in open circuit; hrev is reversible loss of flux B in open circuit; � is reversible temperature coefficient at temperature T, % / .

Aging treatment is easily neglected in manufacturing process of permanent magnetic materials. The so-called aging treatment is to improve the homogeneity of permanent magnetic alloy by heating at a certain temperature. The normal ag-ing treatment method is to hold in atmosphere for 1-3 hours at 80-200 . The permanent magnetic alloy through aging treatment eliminates factors of domain structure and instable texture of the alloy.


Fig.4.74 Diagram of variation of remanence of open circuit vs temperature

4.13.1.1 Stability on temperature for alloys of Sm (Co, Cu, Fe, Zr)z and RECo5

In order to review the stability on temperature of SmCo5 and Sm (Co, Cu, Fe, Zr)7.4

alloys we measured these two alloys at high and low temperature of 1.5 K and 523 K.

Procedures of measurement are as follows:

The conclusion is derived after measurement as follows. 1. Magnetic performance of permanent alloys of SmCo5 and Sm (Co, Cu, Fe,

Zr)7.4 degrade with raising temperature, upgrade with lowering temperature, and have good reversion ability.

2. SmCo5 permanent magnetic alloy of 2:17 type Sm-Co has very good rever-sion ability at temperature range from �196 to 200 .

3. The SmCo5 permanent magnetic alloy does not have magnetic irreversible loss at low temperature.

4. The stability regularity is derived by measuring several specimens of the above mentioned first and second generation rare earth permanent magnetic alloy


at 1.5-523 K: (1) The specimen with a high intrinsic coercivity has good thermal stability; (2) The coercivity of SmCo5 permanent magnetic alloy degrade to the lowest

value at 750 by annealing from room temperature to 1000 ; (3) The SmCo5 permanent magnetic alloy appears the most severe irreversible

magnetic loss at 750 by annealing from room temperature to 1000 ; (4) Any alloy with a high intrinsic coercivity will have, if occurs, a smaller ir-

reversible magnetic loss as well. After being placed at 200-300 for a long term the irreversible flux loss is

zero for Sm (Co, Cu, Fe, Zr)7.4 alloy; if being placed at 350 the irreversible loss will arisen; after aging treatment at 400 the irreversible loss accounts for 70% in a part of specimens and the reversible loss account for 30% (Zhou, Dong, 1999).

4.13.1.2 Thermal stability of NdFeB permanent magnetic alloy

The third generation of the rare earth permanent alloy, NdFeB, have a big im-provement in comparison with the first and second generations. Nevertheless, it have a low Curie temperature of 312 only (while the Curie temperature of SmCo5 and 2:17 type Sm-Co alloys is above 700 ), the magnetic anisotropy field HA of its magnetic phase (Nd2Fe14B) is not very high as well, and that it is sensitive to temperature so that the remanence and coercivity of NdFeB magnet will degrade after being heated. The irreversible temperature coefficient of NdFeB is rather big (�Br = 0.12%/ ). Therefore, improving Curie temperatlare, lowering temperature coefficient and improving thermal stability of NdFeB mag-net is a study subject of theoretic and practical significance (Pan, Ping, Liu, et al, 2003; Pan, Zhao, Ma, 1988; Yang, Yang, 1993; Xu, Ping, Li, Ma, Pan, 1986; Liu, Pan, Luo, et al, 1991).

Author has conducted a series studies to improve thermal stability of NdFeB al-loy. The main way is alloying method. By adding dysprosium, niobium, cobalt, aluminum and gallium, in individual or combinatorial, into NdFeB alloy has got a favorable result in improving its thermal stability (Pan, Ma, Li, 1993; Pan, Ping, Liu, et al, 2003; Pan, Zhao, Ma, 1988; Pan, Li, Li, et al, 1989; Liu, Luo, Pan, et al, 1991; Zhao, Geng, 1991; Wang, Pan, et al, 1999).

The experimental method is: use neodymium, dysprosium, iron, niobium, co-balt, aluminum, gallium and boron materials with purity above 99.5%, prepare with a certain proportion after properly purifying, melt in a medium frequency induction furnace or non self-consuming electric arc furnace. Then comminute the ingot into powder of 3-4 m under protection of medium, shape up in a magnetic field above 1.2 T, sinter in a high temperature furnace of 1100-1200 , pass through aging at 900 and 500-630 respectively, and cool to room temperature. Magnetize the specimen which is in thermo-demagnetization status and measure


magnetism of the specimen (Pan, Ping, Liu, et al, 2003).

A. Adding dysprosium and adding dysprosium and aluminum together into NdFeB alloy

Addition of dysprosium proceeded from cognition as follows: in point of view from main phase of NdFeB alloy, i.e., NdFeB magnetic phase using dysprosium to substitute neodymium will obtain a much higher anisotropy field (HA), as shown in the Table 4.22.

Table 4.22 Properties of the matrix phase

Lattice constant / nm Compound

a c Density/g�m�3 Anisotropy

HA/mA�m�1 Curie temperature

Tc/K

La2Fe14B 0.8822 1.2338 7.40 about 1.59 530

CeFe14B 0.8760 1.2338 7.66 about 2.07 424

Pr2Fe14B 0.8808 1.2244 7.51 about 6.37 565

Nd2Fe14B 0.8803 1.2196 7.60 about 7.16 585

Sm2Fe14B 0.8814 1.2160 7.69 >11.94 612

Ga2Fe14B 0.8773 1.2087 7.90 about 1.91 661

Tb2Fe14B 0.8785 1.2070 7.92 >11.94 639

Dy2Fe14B 0.8768 1.2026 8.03 >11.94 592

Ho2Fe14B 0.8753 1.1990 8.11 about 5.75 570

Er2Fe14B 0.8734 1.1942 8.22 about 0.64 554

Tm2Fe14B 0.8728 1.1928 8.26 about 0.64 541

Lu2Fe14B 0.8697 1.1850 8.47 about 2.07 535

In addition, adding dysprosium into NdFeB ternary alloy can improve coerciv-

ity of magnet. Confect the material referring (Nd1-xDyx)16Fe77B7 formula and suppose x =

0.005, 0.010, 0.020, 0.030, 0.040, 0.050, 0.070, 0.090, and obtained magnetism of alloys as shown in Table 4.23.

It can be seen from the Table 4.23 that the remanence induction strength Br and maximum product of magnetic energy (BH)max decline with increase of dyspro-sium content in (Nd1-xDyx)16Fe77B7 alloy, but the coercivity is improved. The thermo-stability of alloy mentioned above will be heightened with increase x (the Curie temperature has not significant change, within the scope of tolerated error). The ternary NdFeB alloy is appropriate to be used within 80 , that the environ-ment temperature exceeding 80 will result in severe decline in magnetic capa-bility. While x = 0.3 the above mentioned alloy can be used at an environment


temperature of 100 ; when x = 0.5 the alloy can be used at an environment tem-perature of 120 ; when x = 0.9 the alloy can be used at an environment tempera-ture of 150 (because the irreversible loss is less than 3% when the specimen is toasted at relevant temperature).

Table 4.23 Magnetic performance and Tc of alloy (Nd1-xDyx)16Fe77B7

x Br/T jHc/kA�m�1 bHc/kA�m�1 (BH)max/kJ�m�3 Tc/

0.005 1.26 835.81 706.3 295.32 312

0.010 1.25 939.28 728.6 292.93 315

0.020 1.24 1058.68 735.4 290.54 315

0.030 1.23 1146.24 780.2 288.15 320

0.040 1.21 1178.08 795.6 280.19 320

0.050 1.18 1377.09 815.3 263.48 320

0.070 1.13 1464.64 852.7 237.94 320

0.090 1.10 1615.88 860.4 226.86 320

Previously experiments indicate that addition of aluminum into NdFeB alloy is

able to improve the coercivity of the alloy. Thus combined addition of a little aluminum and dysprosium is able to improve the coercivity, and that to improve thermo-stability of the NdFeB alloy. As the result the alloy becomes a quinary alloy, it can be described precisely as (Nd1-xDyx)16(Fe1-yAly)77B7. Among them selecting x = 0.020�0.040, y = 0.01�0.03 can obtain rather satisfied applicable magnet.

B. NdFeB after addition of niobium

Confect the material referring (Nd1-xNbx)16Fe77B7 formula and suppose x = 0.01, 0.02, 0.04, 0.08, 0.15, and obtained magnetism of alloys as shown in Table 4.24.

Table 4.24 Magnetic performance and Tc of alloy (Nd1-x Nbx)16Fe77B7

x Br/T jHc/kA�m�1 bHc/kA�m�1 (BH)max/kJ�m�3 Tc/

0.01 1.22 886 832 270 320

0.02 1.20 1026 916 260 330

0.04 1.17 620 570 235 327

0.08 1.10 510 442 205 310

0.15 1.08 440 410 170 126

It can be seen that as the result the remanence induction strength and maximum


product of magnetic energy (BH)max were declined monotonously by using a little niobium to substitute iron. And that the coercivity strength appears a peak value at x = 0.02. This indicates that adding niobium can improve coercivity. Addition of niobium improves the square degree of demagnetization curve, referring to the curve. And addition of niobium is avail to increase dispersion and suppress growth of crystals. This point is very good for heightening coercivity of NdFeB and improving thermo-stability of magnet. In view of Curie temperature when x = 0.01 the Curie temperature can be taken as unchanged. But as x = 0.08 there are two ferromagnetic phase of Tc1 = 110 and Tc2 =305 from thermomagnetic curve.

C. NdDyFeNbB magnet by adding dysprosium and niobium together

Addition of dysprosium and niobium in same time has advantage over individual use of them. As a result of improvement of coercivity and a nicer quadrate degree in demagnetization curve the maximum product of magnetic energy also be heightened correspondingly. Thus results in a good thermo-stability of the magnet.

D. NdFeB magnet with cobalt

The purpose of addition of cobalt is to aim at the low Curie temperature. There are many research reports about influence of cobalt on main phase of Nd2Fe14B. Cobalt belongs to the transition family. The main phase becomes Nd2(Fe1-x

Cox)14B after adding cobalt. Contribution of elements of the transition family to intermetallic compound follows Slater-Pauling relation, i.e., Tm = 2.6�x.

Magnetic property and Tc of Nd16(Fe1-xCox)77B7 is shown in Table 4.1. It can be seen from Table 4.1 that Br and (BH) max will decline monotonously by

using cobalt to substitute part of iron. For Nd16(Fe1-xCox)77B7 the cobalt raised each 8% the Curie temperature can be heightened 50 (for the main phase of Nd2Fe14B, cobalt substituting each 1% of Fe can heighten Curie temperature 10 for 2-14-1 main phase ). The higher cobalt content may not lower irreversible magnetic loss. But increasing cobalt may lower irreversible temperature coeffi-cient of remanence in open circuit (when x = 0.1, �Br = �0.07%/ ). The reason is mainly due to the formation of Nd(Co, Fe)2 magnetically soft phase. The Nd(Co, Fe)2 may become nucleating center in magnetization process. However, combined addition of cobalt, dysprosium, niobium and gallium to substitute part of iron ob-tained very satisfied result and so as to heightened thermo-stability of NdFeB magnet.

E. NdFeB magnet with cobalt, dysprosium, niobium and gallium

Magnetic performance of NdFeB magnet with cobalt, dysprosium, niobium and gallium, being added together, is listed in Table 4.3 and Table 4.25.


Table 4.25 Magnetic performance and Tc of alloy NdFeB by using Co, Nb and Ga to substitute part of Fe

Composition Br/T jHc/kA�m�1 (BH)max/kJ�m�3 Tc/

Nd12.4Dy3.0Fe72.9Co5.1B6.7 1.10 1683.5 230 380

Nd13.5Dy2.1Co16Ga1.2Fe60B7.2 1.28 971 310 500

Nd13.2Dy2.2Fe70.5Co5.2Nb1.1Ga1.1B6.7 1.03 1830.1 211 370

4.13.2 Time stability

Placed the permanent magnet at a long term at a certain temperature and meas-ured the variation of magnetic property with change of time. For SmCo5 alloy being placed for 1000 hours after heating treatment at 200�250 for 1�2 hours the decline of remained magnetic induction was less than 1%. Generally speaking without this aging experiment the remained magnetic induction would decline obviously at 200�250 for 1 hour and afterwards the variety tends to be smoothly.

Temperature of time stability is among 25�250 and the measurement can be planned at the temperature in accordance with the requirement.

4.13.3 Chemical stability

The permanent magnet works in the environment of a certain acidic or alkaline of chemical workshop or laboratory. That requires the permanent magnet being pro-vided with corrosion resistant property. Place specimen into corrosive medium and then measure its corrosive speed. The RECo5 permanent magnetic alloys are the best rare earth permanent magnetic alloys in chemical stability in comparison with the 2:17 Sm-Co type and NdFeB alloys. However, among RECo5 alloys it is the best when RE is Sm i.e., SmCo5. If RE being Pr or (CeMM)Co5 they are not as so good as SmCo5 in chemical stability. 2:17 type Sm-Co permanent magnet alloy is the best in oxidation resistance. While rare earth ferromagnet alloy is not as good as rare earth cobalt base alloy in chemical stability. NdFeB alloy uses surface coating for oxidation resistance, normally adopting nickel or zinc coating to protect permanent magnetic alloy.

4.13.4 Conclusions

Conclusion was derived through above studies as below: 1. Addition of appropriate cobalt content into NdFeB ternary alloy can height-

ens Curie temperature of the alloy from 312 to 500 . However, addition cobalt


alone can heighten the Curie temperature and lower the irreversible temperature coefficient, but can not reduce the irreversible loss. To improve the thermal stabil-ity it is more effective to add other elements together with cobalt.

2. Thermal stability of NdFeB alloy can be improved greatly through alloying method which is a brief and effective approach for the purpose. But it is very strict for adding what kind of element and how many amounts of the elements. Addition of appropriate mounts of dysprosium, niobium, gallium and/or alumi-num can uniformly heighten the coercivity of the alloy. And through heightening coercivity is an effective way to improve the thermo-stability of the alloy.

3. Using a little dysprosium to substitute neodymium, using a little niobium, aluminum, cobalt and gallium to substitute part of iron and adjusting a propor-tional relationship of amounts can obtain an applicable magnet with high coerciv-ity and high maximum product of magnetic energy. This magnet will have the thermal stability about 70-100 higher than that of ternary NdFeB alloy.

4. Any rare earth permanent magnetic alloy which possesses a high coercivity will have a better thermal stability and smaller irreversible loss. Heightening co-ercivity is an effective approach to improve stability of the rare earth permanent magnetic alloy.

5. SmCo5 permanent magnetic alloy has a very good reversion property at the temperature from �196 to 200 . It has no irreversible loss at low temperature. The irreversible loss will reach the maximum by annealing at 750 . And that 2 : 17 type Sm-Co alloy is the best in oxidation resistance.

6. The rare earth cobalt base permanent magnetic alloy has the chemical stabil-ity, time stability and thermal stability better than those of the rare earth ferrous base permanent magnetic alloy. Aging experiment avails stability of mother alloy.

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Chapter 5 Developments and Prospect of the Rare Earth Permanent-magnet Alloys

The rare earth permanent-magnet alloys are the broadly used foundational func-tional materials. Rare earth permanent-magnet alloys have achieved great pro-gress in scientific research, manufacturing and application in recent few decades, and their applications have pervaded into every region of national economy. Rare earth permanent-magnet alloy has become important material basis of new tech-nology.

China is not only rich in the storage capacity of rare earth material, but also af-fluent in output. Nowadays, improving the performance of rare earth permanent magnetic alloy and searching for new generation with outstanding performance has become one of focuses being paid most attention to this field.

The development of nanometer science and technology brings new challenges and opportunities for the rare earth permanent magnetic industry. Nanometer bril-liant exchange coupling is the permanent magnetic material which is one of the material with brightest future and the development of this material and correlation technique certainly promote the technical progress of rare earth permanent mag-netic alloys and upgrade performances of traditional products, and thus it will result in development of high and new technology products for rare earth perma-nent materials.

Two-phase composite nanocrystalline rare earth permanent magnet is a kind of promising magnetic material. Their theoretical maximum magnetic energy prod-uct can reach to 993.75kJ/m3, which as twice as that of NdFeB magnets. And the low rare earth content gives them great advance in cost efficiency. Additionally, the better corrosion resistance and higher Curie temperature broaden their appli-cation area. The development of this material will gear up the development of rare earth permanent magnets and the performance of traditional magnets.


5.1 Overseas General Development

Rare earth metals, which exist in rare earth permanent magnetic alloys as alloy elements, transform into 3d transition intermetallic compounds together with 3d transition metals. Among them rare earth cobalt base and rare earth iron base in-termetallic compounds have been successfully applied in electric machine, in-strument, petrochemical industry, automobile, automation and computer areas.

Rare earth cobalt based permanent magnetic alloy started in 1960s. Early in 1959, E. A. Nesbbit, et al (Wesbbit, et al, 1959) and in 1960 W. M. Hubbard et al. found GdCo5 alloy (Hubbard, Adams, Gilfrich, 1960) successively. In 1966, G. Hoffer, et al. found that K1=5.7�106 J/m3, 0Ms=1.06 T (Hoffer, Strnat, 1966) for YCo5. In 1967, K.J. Strnat, et al. firstly made out permanent magnetic alloy YCo5 using powder metallurgy technique, but its performance was very low, (BH)max = 9.6 kJ/m3; in succession they synthesized permanent magnetic alloy SmCo5 using Sm substituted for Y by element substitution method, and the magnetic character-istics were as follows: Br = 0.51 T bHc = 254.7 kA/m and (BH)max=40.6 kJ/m3

(Strnat, Hoffer, Olson, et al, 1967). In 1968, K. H. Buschow, et al. prepared per-manent magnetic material using powder metallurgical technique to enhancing alloy density which created new record of product of magnetic energy at that time: (BH)max=147.3 kJ/m3 and the intrinsic coercivity mHc = 1257 kA/m (Buschow, et al, 1968). The technique of permanent magnetic alloy SmCo5

was come to perfection during the year 1969 to 1972: M. G. Benz, et al. improved the performance of the alloy using the method of liquid-phase sintering in 1970 (Benz, Martin, 1970). In 1972, R. J. Chaless, et al. used reduce-diffuse (R/D) method prepared permanent magnet (Pr0.5Sm0.5)Co5 with (BH)max =207 kJ/m3 (Chaless, et al, 1972). In 1973, A. C. ��, et al. synthesized the single crystal of SmCo5 creating new record in high magnetic characters (BH)max = 254.7 J/m (32 MGs·Oe) (��, 1973). As the representative of first generation rare earth permanent magnetic alloy SmCo5 has main performance as follows.

To acquire the magnetic performance shown in the Table 5.1 the alloying com-position must be calculated as per the compound molecular formula. The compo-sition is samarium 36%-37% and cobalt 63%-64% in mass fraction. High tem-perature sintering is the key factor to form high performance alloy; the heat treatment system are holding for more than 1 h at 1120 , cooled to 900 in speed of 0.6-0.9 /min and holding for 3-5 h then quenching to room temperature. SmCo5 has the biggest magnetic anisotropy and is much easier than other rare earth permanent magnetic alloys to gain high coercivity. The main processes of producing SmCo5 using sintering process (alloy melting process) are alloy melting � milling �shaping in magnetic field � aging � processing �magnetizing. To prevent oxidization melting, milling and sintering should be carried out under the


protection of argon atmosphere.

Table 5.1 Performance of SmCo5 permanent magnetic alloy

Magnetic performance

Br/T Hc/kA·m-1 (BH)max/kJ·m-3 Tc/ Magnetizable axis

0.9-1.1 1100-1590 117-180 720 c

Anisotropy constant K of magnetic crystals Anisotropy field HA Density/g·cm�3

(9.5-11.2)×106 (9.5-11.2)×107 16716-23084 210-290

(19.3±1.4)×106 (19.3±1.4)×107 35024 440±207.18-7.92

The mechanism of coercivity of SmCo5 belongs to nucleation mechanism, that

is to say, the coercivity of the alloy is controlled by its nucleation field and the formation of reversal magnetization domain should be in upfield.

The secondary generation rare earth permanent magnetic alloys formally came into being in 1977, when T. Ojima, et al. successfully made the rare earth perma-nent magnet of the composition of Sm(Co, Cu, Fe, Zr)7.2 using powder metallur-gical technique. That rare earth permanent magnet created the highest record in product of magnetic-energy at that time with: (BH)max= 238.8 J/m3 (30MGs·Oe). Its development process spent ten years. At first in 1968, E. A. Nesbbit, et al. used cupper to substitute part of cobalt prepared intermetallic compound of Sm(Co,Cu)5 and obtained the permanent magnetic alloy with (BH)max = 31.8-55.7 J/m3 (Nesbbit, Willens, Sherwood, et al, 1968). In 1974 A. J. Perty, et al. and in 1976 A. Menth, et al. researched Sm((Co+Fe)1-x-yCuxMy)5 8.5 alloy and prepared practical applicable rare earth permanent magnetic alloy. The performances of the second generation rare earth permanent magnetic alloy Sm(Co, Cu, Fe, Zr)z

(z=7-8.5) are showed in Table 5.2.

Table 5.2 Performance of the alloy Sm(Co, Cu, Fe, Zr)z (z=7-8.5)

Magnetic performance

Br/T Hc/kA·m-1 (BH)max/kJ·m-3 Tc/ Magnetizable

axis

1.0-1.3 500-700 230-260 850 c

Average temperature coefficient of magnetic induction Operation tempera-

ture T/ K1/J·m-3 �25-100 �25-200

Density/g·cm-3

350 4.3×106 �0.025 0.03 8.4

In order to acquire high performances shown in Table 5.2, a special heat treat-

ment processes should be adopted, which include solid solution treatment and isothermal aging. The solid solution treatment is aim to acquire homogeneous monomial solid solution, and solid solution temperature is at 1130-1170 while grade aging treatment at below 850 . After solution treatment and isothermal


aging treatment at below 850 , the alloy is of exiguous cellular microtexture in the state of high coercivity. The diameter of cell is about 50 nm and cell wall thickness is about 5nm. When aging at 830 for more than 20 h the cellular tex-ture of the alloy already becomes incomplete and part of it is destroyed. The fab-rication process of alloy is to put the prepared raw material as per nominal com-position into furnace for melting at protection of argon atmosphere at first, sec-ondly quench melted alloy in water cooled mold, and the third alloy mill and then mold ingot in magnetic field; afterwards, process step aging as per above men-tioned heat treatment system. Through such multistage aging (precipitation hard-ening, producing two stage decompound) the coercivity of the alloy can be im-proved in a large magnitude, the maximum magnetic energy product of the alloy can reach 250 kJ/m3 or above.

The magnetic anisotropy constant K1 and anisotropic field HA of the secondary generation rare earth permanent magnetic alloys are both less than those of the first generation, and so is the coercivity. The coercivity of alloy depends on pin-ning field because the pinning field and initial magnetization curve enhance gradually along with increase of magnetic field and at last reaches saturation (the magnetic field must be bigger than the pinning field), its coercivity is determined by the structure of two phases, i.e., when magnetization and reverse magnetiza-tion the domain wall of Sm2Co17 is pinned by SmCo5, which enhanced the coer-civity.

Space aeronautic and aviation fields requires that the rare earth permanent magnetic materials have good performance and thermal stability at 400-500 . The 2:17 Sm-Co permanent magnetic alloy becomes preferred material for the requirement. In 2000 EEC corporate, USA, produced Sm-Co magnet with the maximum magnetic energy product of 79.6 kJ/m3 and coercivity of 95 kA/m at 500 .

In China the Research Institute of Iron & Steel of Beijing used near fast solidi-fication technique to prepare permanent magnetic alloy with the maximum mag-netic energy product of 79.89 kJ/m3 and coercivity of 95 kA/m at 500 by changing iron content to control the coercivity coefficient.

The Curie temperature and saturation magnetization intensity of the secondary generation rare earth permanent magnetic alloys are better than those of the first generation; the second generation has a smaller reversible temperature coefficient and a better thermal stability as well, so it is fit for using at higher temperature.

The third generation rare earth permanent magnetic alloy was made by M. Sa-gawa, et al. of Sumitomo Special Metals Corporation of Japan in 1983 using pow-der metallurgic method. The alloy was the Nd15Fe77B8 permanent magnetic alloy with (BH)max=286.6 kJ/m3 (36MGs·Oe), which created the highest record in magnetic-energy product at that time (Sagawa, Fujimura, Togawa, et al, 1984). It was based on the researches and studies of many scholars and scientists. Accord-


ing to early researches people found that REFe2 and other rare earth series have higher coercivity at low temperature and hard magnetic properties of amorphous material REFe2 was improved during crystallization. During 1972 to 1973, A. E. Clark et al. made compound TbFe2 into amorphous state at room temperature and annealed, as a result its coercivity improved in big magnitude to mHc = 273 kA/m (3.4 kOe) and (BH)max = 71.64 J/m3 (Clark, Belson, 1972; Clark, 1973). K. N. Koon, et al. found high coercivity in amorphous alloy La5Tb5 (Fe80B20)90 in 1981 (Koon, Das, 1981). H. H. Stadelmaier, et al. found that Gd3Fe20C phase had the similar structure with Zn22C3 in 1981. Successively, G. C. Hadjipanayis, et al. confected Pr15Fe76B6Si3 and acquired mHc=1194 kA/m (15kOe) and (BH)max = 103.5 kJ/m3 (13MGs·Oe) through fast quenching and heat treatment process (Had-jipanayis, Hazelton, Lawless, 1984). D. J. Sellmyer, et al. found that alloy hard magnetization phase was RE2Fe14B phase using X-ray analysis, which belonged to square structure compound.

Birth of the third generation rare earth permanent magnetic alloys realized the desire people expected for a long term not only because it created the highest re-cord in magnetic-energy product but also because using neodymium with a higher reserves to substitute samarium with the reserves much fewer, and using iron to substitute cobalt the strategic material. Thus permanent magnet can be used in much wider domains with a better performance and cheap material without any resource limitation. Developing, manufacturing and using the third generation rare earth permanent magnetic alloy were aroused strongly in industrial and aca-demic community, which rapidly changed the situation of research, manufacture and application of rare earth permanent magnetic alloys.

In 1985 the magnetic-energy product of NdFeB alloy reached 372.92 kJ/m3

(47.1 MGs·Oe) already. In 1988 Sumitomo Special Metals Corporation of Japan reported the highest performance of a new NdFeB material that: (BH)max = 446.24 kJ/m3 (55.78 MGs·Oe), Br = 1.514 T (15.14kGs), and mHc = 694.4 kA/m (8.6 kOe) (Wang, 2001).

In 2002 laboratory of VAC corporate in Germany produced sintered NdFeB permanent magnet with magnetic energy product reaching 451.3 kJ/m3 (56.7 MGs·Oe). Afterwards, Nippon Sumitomo corporate raised the maximum magnetic energy product to 460.1 kJ/m3 (57.8 MGs·Oe) in 2005 and produced a sintered magnet of NdFeB with the maximum magnetic energy product (BH)max = 474 kJ/m3 (59.5 MGs·Oe) in 2006, the other magnetic parameters of this sintered mag-net: Br = 1.555 T, iHc = 653 kA/m.

That we called NdFeB magnet as a king of permanent magnetic material is in a certain condition because the Curie temperature of the ternary NdFeB alloy is only 312 so that it is suited to be used in a working environment below 100 and when ambient temperature is over 100 its thermo stability becomes worse than SmCo5 and Sm-Co of 2:17 type. In other words, at higher temperature the


practicability of this kind of magnet is influenced by the following three factors: irreversible loss of magnetic flux, reversible loss of remanence Br and reversible loss of intrinsic coercivity. The irreversible loss of magnet with low coercivity represents considerable severe and even the high magnetic-energy product loses its advantage of the already obtained high magnetic-energy product after tempera-ture circulation (from low temperature to working temperature and again return-ing to room temperature) and causes big change in operating point of magnetic circuit design because of low coercivity. The design of magnetic apparatus re-quires permanent magnet to provide a constant magnetic flux and does not allow the magnetic flux to vary apparently with change of temperature. In the design of magnetic circuit of magnetic apparatuses and precise instrument generally adopts a brief interior compensation (using low � compensating magnet) to ensure a con-stant magnetic flux, while sometimes external compensation is also adopted (add-ing compensation material to out of magnet) (Pan, Chen, Liu, et al, 1994).

It was found in experiment that the Curie temperature of NdFeB mainly indi-cated the Curie temperature of Nd2Fe14B. The relevant references pointed out that Curie temperature of Nd2Fe14B took RE = Gd as the maximum and then lowered in to sides of higher and lower of the atom ordinal in sequence. Using cobalt to substitute part of iron can enhance Curie temperature, such as Nd2(Fe1-xCox)14 B when x = 0.1 the alloy is of the maximum theoretical magnetic-energy product but afterwards found addition cobalt alone lowered constant of anisotropy of Nd2Fe14B. Addition of cobalt and some other components simultaneity achieved good result, such as NdFeCoGaB alloy could raise Curie temperature of NdFeB alloy from 312 to 450-500 and Nd16Co16Fe61-xGaxB7 at x = 2 the intrinsic co-ercivity of the alloy appeared peak value. Its thermal stability improved and be-cause of adding gallium the temperature coefficient of Br and intrinsic coercivity Hc become smaller.

Furthermore, the coercivity and irreversible loss can be improved by using aluminum, niobium, molybdenum and tungsten to substituting part of iron. The Curie temperature of NdFeB can be raised from 312 to 450 by cooperative using aluminum and cobalt to substitute part of iron. Using 4%-10% of cobalt to substitute part of iron could raised the Curie temperature to 440 and if the per-cent of cobalt is raised to 16% and adding a little aluminum (such as 2%) can raise Curie temperature to 480-500 and decrease the reversible temperature coefficient significantly to �Br = 0.04%/ .

M. Toknnaga, et al researched the result using niobium to substitute part of iron in NdFeB. The result indicated that: addition of niobium in NdFeB alloy made the reversible temperature coefficient of Br be lowered by 7% and � the reversible temperature coefficient of Hc be reduced by 1% (at temperature range of 23-125 ). The granule size of NdFeBNb alloy decreased along with increasing of niobium content and that the irreversible loss declined obviously along with re-


ducing of average crystal granule size. In order to improve thermal stability generally a great lot of efforts have been

focused on enhancing coercivity. Using dysprosium to substitute part of neodym-ium may be a way to enhance coercivity. Researches indicated that although using dysprosium to substitute part of neodymium can lower Tc yet � is declined. The main reason is that dysprosium sublattice partly compensates degradation of magnetic moment of neodymium. Using a little dysprosium to substitute part of neodymium and controlling temperature system could improve Hc up to above 1600 kA/m (2�104 Oe)(Cai, Rong, 2012; Pan, Ma, Ping, et al, 1991).

To lower temperature coefficient of NdFeB using erbium and dysprosium to partly substitutes neodymium of NdFeCoB alloy can achieve good result.

Irreversible loss of (Nd0.8Dy0.2)(Fe0.835Co0.05B0.08Nb0.015Ga0.04)5.5 magnet is less than 5% after exposed at 260 . (Nd0.08Dy0.12)15Fe70.2Co5Al1.8B8 magnet has pref-erable performances: Br=1.19 T, iHc = 1630 kA/m, (BH)max = 275 kJ/m3 (Liu, Pan, Luo, et al, 1991; Pan, Chen, Liu, et al, 1994).

Great deals of corrosion resistant researches have been done on the weak point being rusty easily of NdFeB magnet under certain conditions.

In 1988, J. Jacobson, et al presented their research result about oxidation reac-tion of NdFeB, NdFeDyB and NdFeAlB alloy at different temperatures in oxida-tion environment. The thesis pointed out: when magnet was exposed in the humid atmosphere at room temperature oxides occurred at edges and corners of the magnet but dispersed over its surface discontinuously. Under dried atmosphere at 150 corrosion phenomenon also took placed and corrosion occurs at the inter-granular boundary. It was determined by X-ray analysis that the oxide in major part composed of ferric oxide and its chemical reaction equations are as follows:

Fe + 1/2O2 ==== FeO Fe + 2H2O ==== Fe(OH)2 + H2

In the humid atmosphere and at 150 the chemical reaction equations are: 2Nd + 3/2O2 ==== Nd2O3

2NdCl3 + 3H2O ==== Nd2O3 + 6HCl It can be seen that most of oxides turning to Nd2O3. That 2NdCl3 +3H2O ==

Nd2O3 + 6HCl indicates the specimen being polluted by chloride. At room tem-perature or at 150 but in humid atmosphere oxidation kinetics follows the pa-rabola rate equation that increment in weight increases along with rising of tem-perature. Whereas in drying atmosphere at 150 the weight increment is faster than that in humid atmosphere and a simple parabola rate law is not followed here.

According to the researches of A. S. Kim, et al., adding aluminum and dyspro-sium in NdFeB alloy could also slow down the rate of oxidation at room tempera-ture in the humid atmosphere.

C. N. Christodoulou, et al pointed out based on their researches that powder was oxidized at low temperature when powder was heated (in sintering process).


According to the researches of N. Imaizumi characteristics of magnet became damaged and easily corrodible after mechanical processing. But the damaged magnetic characteristics was restored after heat treatment at 600-1000 . Further detection discovered the reason for above the damage that the oxide layer of about 0.3m formed the surface after mechanical processing. Corrosion phe-nomenon did not happen to the magnet when it is exposed in the atmosphere for 100 h (in condition of 60 and 90% relative humidity).

In order to prevent magnet from corrosion by oxidizing or oxides people adopted two countermeasures: one is to add a little chromium or nickel to im-prove corrosion resistance; the other one is to add protective coating on surface of magnet.

According to report of U. C. Standiego and Osterrecher oxidation rate of Nd2Fe12Cr2B lowered from one half to one third in comparison with Nd2Fe14B.

Anticorrosion coating means coating of nickel, aluminum, chromate and epoxy resin, etc. and uses different coating according as different purpose. Not all the anticorrosion criterions are the same for different countries. That used in Japan is stricter than that in USA. It should be pointed out that the anticorrosion coating is not perfect and the metallic coating would begin to be spoilt from surface of mag-net after holding at atmosphere for 2-3 years but electroplating or chemical coat-ing is easy to have acidic or alkaline solution remained in pores of magnet. Resin coating is the better way in a severe corrosive environment.

5.2 Domestic General Development

China has abundant rare earth resource and its reserves and output are both in the top of the world, which is advantaged for the development of rare earth perma-nent magnetic alloys. The development of rare earth permanent magnetic alloys in China begins in 1980s. Early in 1970 both Beijing General Research Institute of Iron and Steel and the Beijing General Research Institute of Nonferrous Metals could supply SmCo5 in small amount and successively Baotou Research Institute of Rare Earth, Southwest Application Physics Research Institute and some other units developed PrCo5, respectively. And the HA of compound of PrCo5 reached 11,542-14,328 kA/m, its magnetic induction intensity was up to 0Ms = 1.25 T (4Ms = 12.5 kGs) and its theoretical magnetic-energy product was up to 310.4 kJ/m3. Through its performance was lower than that of SmCo5 but the reserves of Pr is 2-4 times of the reserves of Sm in rare earth mines. When developing SmCo5 people already concerned the development of PrCo5 permanent magnetic alloy. The Curie temperature of PrCo5 is close to that of SmCo5. SmCo5 has a higher magnetic anisotropy but PrCo5 has a higher Ms and the resource of Pr is abundant than that of Sm so that taking both advantages made (Sm, Pr)Co5 by using Pr to substitute Sm in powder metallurgic method. Owing to its good magnetic per-


formance and a better economy prospect (Sm, Pr)Co5 has been produced in a great deal and used widely. Successively liquid phase sintering process was adopted and its magnetic characters were enhanced in large magnitude. The mag-netic characters of alloy of composition Pr0.5Sm0.5Co5 was up to (BH)max=159.2 kJ/m3, Br = 0.89 T and iHc = 1154 kA/m after solid phase sintering and its maxi-mum magnetic energy product was up to 199 kJ/m3 but its performance at low temperature was far from as good as that of the SmCo5.

It was found from the elements periodic table that Ce is close to Pr and Ce is cheaper than Pr and SmCo5 and its resource is abundant so that (Ce, MM)Co5 was developed successively in China. In order to reduce cost further Ce(Co,Cu,Fe)5

permanent magnetic alloy was developed successfully using iron and copper to substitute part of expensive cobalt. Through the magnetic performance of the new developed alloy was as good as SmCo5 and (Pr, Sm)Co5 its still reached magnetic performance as follows: Br = 0.6-0.8 T, iHc = 860-1114 kA/m and (BH)max = 83-138 kJ/m3.

It was found in practice application that the above mentioned products could not satisfied the requirement for the application of traveling-wave tube, magnetic bearing and precision instrument because their reversible temperature coefficient of magnetic induction is on the high side. The requirement for application situa-tion of above products asked the magnetic induction reversible temperature coef-ficient of permanent magnetic alloy is less than 0.02%/ . How to lower the temperature coefficient of the above mentioned alloy? It was found through re-searches that the magnetic induction reversible temperature coefficient of RECo5 is mainly determined by the dependency relationship of magnetic moment of molecule constituted it and temperature. In RECo5 the light rare earth metals (La, Ce, Pr, Nd, Sm and Eu) atomic magnetic moment with the atomic magnetic mo-ment of Co is the ferromagnetism coupling and it has negative temperature coef-ficient. The atomic magnetic moment of heavy rare earth metals (Gd, Tb, Dy, Ho, Er, Tm, Yb and Lu) with atomic magnetic moment of Co is ferrimagnetism cou-pling, which has positive temperature coefficient within a certain temperature range. Both of them have the function of temperature compensation. Therefore, using light rare earth metals partly and heavy rare earth metals partly as RE in RECo5 alloy can obtain the permanent magnetic alloy with low temperature coef-ficient RECo5. Sm0.6Dy0.4Co5 was made with the maximum magnetic energy product 72.4 kJ/m3 and the magnetic induction reversible temperature coefficient -0.0003%/ at 22-47 through researches and that by adjusting the ratio of the light and heavy rare earth metals can also obtained an alloy with zero-temperature coefficient.

From the end of 1970s to the beginning of 1980s domestic researchers working on rare earth permanent magnetic alloy systemically developed Sm(Co, Cu, Fe, Zr)z(z = 7-8) closely following overseas development of Sm-Co 2:17 type perma-


nent magnetic alloy. Beijing University of Science and Technology, Beijing Gen-eral Research Institute for Nonferrous Metal, Institute of Electronics of China Academy of Science, etc. made many systemic researches on the performance of permanent magnetic alloy, especially the mechanism of coercivity. In 1980, D. Li group of Beijing General Research Institute of Iron & Steel successfully devel-oped Sm-Co 2:17 type permanent magnetic alloy with low temperature coefficient with the magnetic performance as follows: for Sm1.6Er0.4Co10Cu1.5Fe1.2Zr0.2 alloy average temperature coefficient within temperature range from 50 to 100 � = 0.006%/ , Br = 0.99 T, (BH)max = 179.8 kJ/m3; and for Sm1.2Er0.8Co10Cu1.5

Fe3.2Zr0.2 average temperature coefficient � = 0.000 %/ and within temperature range from 50 to 100 � = 0.002%/ , Br = 0.94 T, iHc = 413.9 kA/m, (BH)max = 143.2 kJ/m3. In 1987 J. Wang from Baotou Rare Earth Research Insti-tute systemically studied on Sm0.75Er0.25(Co, Cu, Fe, Zr)7.4 permanent magnetic alloy with magnetic performance: Br = 0.906 T, iHc = 1018.8 kA/m bHc = 636.8 kA/m, (BH)max = 141.8kJ/m3.

After long time researching, the General Research Institute for Nonferrous Metals worked out SmCo5 powder by reduce-diffusion method; Yu Chengzhou, Ying Qimin group together with Gao Qinghai, et al from Shanghai Yuelong Chemical Plant mastered the key technologies for industrialization and set up SmCo5 alloy powder factories. In 1980 to 1983, Tang Renyuan, Li Guokun, et al proclaimed application result of rare earth permanent magnetic material in electric machine, magnetism transmission, magnetism biology and electronic industry; and Sun Tianduo, Sun Daku, et al presented theoretical investigation results on coercivity mechanism of SmCo5 and Sm-Co 2:17 alloy in the international sci-ence conference, which promoted improvement in magnetic performance of rare earth permanent magnetic materials and exploited application scopes for rare earth permanent magnetic materials. In 1983, Pan Shuming subject group prom-ulgated in International Rare Earth Permanent Magnetic Material Academy Con-ference that the observation on variation of Sm2Co7 phase at high temperature at first using JEM-1000 HVEM in Beijing General Research Institute for Nonfer-rous Metals, and, together with Jin Hanmin from Jilin University, presented opin-ions on that coercivity of SmCo5 degrades at the temperature of 750 is because defects in precipitated 2:17 phase which is of a low magnetic anisotropy and so that leads to degradation of the coercivity.

The development of iron base rare earth permanent magnetic alloy began in 1980, while the broad scale researches began in 1983. Some innovations are as follows:

In 1986, Li Wei subject group, Beijing General Research Institute for Iron and Steel, proclaimed that the permanent magnetic alloy was prepared using heavy rare earth metal dysprosium to substitute part of Nd and using iron to substitute part of Co with an average coefficient of magnetic induction being lower than


�0.028%/ measured at temperature of 20-100 , and its magnetic performance being: Br = 0.95 T, iHc = 1200 kA/m, (BH)max = 160 kJ/m3. Also, they developed NdFeB alloy with zero temperature coefficient and its magnetic performance be-ing: Br = 0.79 T, iHc = 1280 kA/m, (BH)max = 114 kJ/m3. In 1988, Zhou Shouzeng subject group of University of Science and Technology Beijing proclaimed the (Nd0.5Dy0.5)15.5Fe51Co26B7.5 permanent magnetic alloy with � = 0.000%/ (i.e., the ferrous base rare earth permanent magnetic alloy with zero temperature coef-ficient), within 20-100 and the magnetic performance: Br = 0.88 T, iHc = 1233.8 kA/m, bHc = 525.4 kA/m, (BH)max = 119.4 kJ/m3.

Among high performance permanent magnetic materials the most attractive one is the high coercivity Nd-Fe-B permanent magnetic material because the rema-nence and magnetic energy product of sintered Nd-Fe-B permanent magnet have reached 93% of the theoretical value but its coercivity only reaches 12% of the theoretical value, i.e., 12% of the anisotropy field of compound Nd2Fe14B. Therefore, there is a large space for upgrading of coercivity. Accordingly, the rare earth permanent magnetic material of high coercivity and high performance without using heavy rare earth element will become a new and attractive hot sub-ject for study.

The research of ferrous-base rare earth alloy has been improved in China. In the resent more than 10 years State Intellectual Property Office promulgated in-novation patents, such as “The method of preparation of metalloid-intermetallic compound and product”, “The method of preparation of single-phase intermetallic compound and its product”, “Carbide permanent magnet” (Luo, Dong, 1999; Luo, Dong, 1998; Luo, Dong, 1997); “The manufacturing method of permanent mag-netic material”, “The manufacturing process of thorium-manganese 12 type rare earth-iron permanent magnetic material”; “The fabrication of neodymium iron permanent magnetic alloy in coprecipitation restore diffusion process”; “A method of producing Sm-Fe-N permanent magnetic alloy powder in reduce-diffusion process”; “High performance bidirectional rare earth permanent mag-netic material and its preparation method”; “A kind of high stable rare earth-iron permanent magnetism carbide and its manufacturing method”; “A method of fab-ricating hydrogenation-disproportionation-dehydrogenation-reform rare earth permanent magnetic powder”; “A kind of preparation method of permanent mag-netic powder”, etc.

Pan Wei, et al successfully developed radiation orientation magnetic ring. In recent years in domestic rare earth permanent magnetic products industry rapid coagulating technology, the popularization of hydrogenation milling technology and one-step forming instead of two-steps forming were adopted to increase product density and promote product quality of Chinese rare earth permanent magnetic alloys reaching overseas advanced level.

The production output of rare earth permanent magnetic alloys in China go


upon the top stage was not divided from the industrialization policies promoted by academician Wang Zhenxi, et al. In the twenty first century, China should take full advantage of advantaged rare earth resource and promote the rare earth per-manent magnetic alloy product to go up to the advanced status in the world.

There is a lot of equipment in developing melting technique for high perform-ance sintered Nd-Fe-B magnet in China.

General Research Institute for Nonferrous Metals is a largest-scale comprehen-sive research and development institution in nonferrous metals, the rare earth metallurgy and material, micro-electronics and photoelectron, and rare metal and noble metal materials. It has 10 national levels attached institutions such as Na-tional Engineering Research Center for Rare Earth Material, etc. National Engi-neering Research Center for Rare Earth Material is the core of research and de-velopment of General Research Institute for Nonferrous Metals, which applied 206 patents about the rare earths, including 187 invention patents, 9 overseas pat-ents. It is one of enterprises firstly developed SmCo permanent magnet in domes-tic in 1970s. Its SmCo permanent magnet was applied to the 1st artificial satellite of China and made contribution to Chinese “A-bomb, H-bomb and artificial satel-lite”. National Engineering Research Center for Rare Earth Material independ-ently developed the 1st belt throwing off furnace for rare earth NdFeB alloy which broke through the key technique for quick cooling thick belt of high per-formance NdFeB (quick solidified casting belt), obtained 4 invention patents and successfully realized mass production, that made important contribution to let Chinese NdFeB industry enter into high-end application field broke through mo-nopolization of foreign countries and obtain the 2nd grade national invention award.

Since 2003 Ningbo Konit Industry Co. Ltd. became a leader in production of high end product of the rare earth permanent magnetic alloy in application-volume coil motor (VCM) for computer hard disk driver. Their productivity reached thousands tons each year since it produced sintered Nd-Fe-B magnet and these products have entered into markets of developed countries such as Japan and Europe markets.

Beijing Sanjili Rare Earth Co. Ltd. has mastered excellent casting and solidifi-cation technique and near fast solidification squamous technique (SC), had 4 ad-vanced vacuum induction furnaces with thousands tons output of Nd-Fe-B annu-ally and became a domestic classic manufacturer.

Ningbo Ninggang Permanent Magnetic Material Co. Ltd. undertakes specially development and production of sintering 1:5 and 2:17 Sm-Co magnetic materials. It has advanced production, machining and inspection equipment and facility for annual yield of 250 t Sm-Co magnetic materials and the quality of its products reach the advanced level of the world. Ningbo Ninggang Permanent Magnetic Material Co. Ltd. independently developed 1:5 and 2:17 Sm-Co permanent mag-


netic materials with low temperature coefficient which can be widely used in aerospace, military transducers, etc. It is the largest enterprise of the world in scale of production Sm-Co permanent magnet.

As one of fast-growing rare earth magnet manufactures, Feller Magtech spe-cialized in NdFeB permanent motors used in oil pumps, water pumps in coalfield and direct driven wind power generators.

5.3 Development Survey of Preparation Technology

For many years technique of powder metallurgy method, alloy melt rapid solidifying method, diffusion reduction method, etc. to prepare NdFeB permanent magnetic material achieved quiet great progress. The process flow diagrams of several methods are shown in Fig. 5.1.

Alloy melt rapid solidifying method: USA is famous for manufacturing NdFeB by this method. GM Corporation used this method to produce NdFeB permanent magnetic materials. IG. Technologies, Ovonic Synthetic Materials Company, Har-rison Dickson, etc. also used this method. For example, the powder of MQ3/4 of GM Corporation exported to Japan was produced using this method. The main steps of this method are mainly as follows: melting neodymium, iron and boron in high-frequency electric furnace, afterwards ejecting to water-cooled copperplate for quenching, thus obtained the alloy strip with a thickness approximate 20m, and after this grinding the alloy trip to fine powder. According to process flow in Fig. 5.1, processing and further producing goods as requested.

Fig.5.1 Flow chart of manufacture of NdFeB permanent magnet


China began to make broad researches on this method since 1983, created its own features with bold innovations, and obtained independent patents of China. The NdFeB compound factories have been built independently in China, which can provide magnet of any shape for customers.

Powder metallurgy method: currently 80%-90% of NdFeB permanent magnetic alloy has been manufactured by this method. Compared to alloy melt rapid solidi-fied method, this method needs less investment and get higher magnetic energy product. Its process flow diagram is shown in Fig. 5.1. That is raw material , alloy smelting(medium-frequency induction furnace) � milling � forming in magnetic field � sintering � aging (heat treatment) � magnetization � anisot-ropic NdFeB permanent magnetic alloy. Because neodymium is easily to be oxi-dized, the key technology is protecting against oxidation. Vacuum pumping and inert gas protecting is necessary for alloy smelting and sintering. Chinese has got-ten several intellectual properties and patent rights by bold innovation in technol-ogy research.

The sintering process plays a decisive role in the manufacturing process of high-performance rare earth iron-boron base materials. The purpose of sintering is to make sure that the migration of atoms at high temperatures occurs between the powder particles bonded to the alloy performance quantitative and qualitative transformation, bond strength between the powder alloys to meet the require-ments. Sintering process of diffusion, flow, re-crystallization, creep, reply materi-alized job is almost the same time (Pan, 2011).

Diffusion reduction method: in Symposium of International Rare Earth Perma-nent magnetic Materials and Application Convention opened in May 1985, C. Herget of Goldschmidt Co. in the former West Germanic reported the technology and theory of producing NdFeB permanent magnetic alloy with rare earth oxide in laboratory.

1200 C2 3 40 60 vacuum

15 77 8

(15/2)Nd O + 71(2/3)Fe + (4/30)Fe B + (45/2)Ca

Nd Fe B + (45/2)CaO

1,

Goldschmidt Co. had produced permanent magnet with the maximum magnetic energy product of 167 kJ/m3 but with a low intrinsic coercivity. By using dyspro-sium to substitute part of neodymium and using aluminum to substitute part of iron the intrinsic coercivity was increased sharply. The intrinsic coercivity of NdDyFeB alloy was 708 kA/m and the maximum magnetic energy product was 237.2 kJ/m3. The advantage of this method is low cost. Neodymia reacts with reducer and be reverted to neodymium in this process. When inert atmosphere is heated up to 1200 neodymium is inter-diffused with iron and boron to become NdFeB permanent magnetic alloy. The coercivity of magnet made by this method is lower than those made by powder metallurgy method and alloy melt rapid so-lidified method. This method is still in development.


Hot pressing and hot rolling method: main procedures this method is that the Pr(Nd)FeCuB alloy ingot � enclose with metal � heating � hot machining, and then the permanent magnetic alloy with magnetic anisotropy is obtained through hot pressing and hot rolling. Some results have been acquired in China using this method and the permanent magnetic alloy with maximum magnetic energy pro-duct (BH)max 240 kJ/m3 have been produced.

NdFeB permanent magnetic alloys could be produced by spark plasma sinter-ing technique. Spark Plasma Sintering (briefly called as SPS) is use plasma to be conductive bearer inside sample to form electric current, moving directionally under the effect of electric field (plasma swathed on particle surface), merging together at particles connecting node (local current density increase), and forming high sintering points due to high temperature and thus caused substance transfers by evaporation, diffusion, etc. Sintering points increase gradually and distribute equably so that turns into comprehensive sintering status. Its remanence Br and the maximum magnetic energy product (BH)max appear at 780 , intrinsic coercivity iHc turn up at about 800 heat treatment temperature is 950-1000 . Its magnetic properties can reach: Br = 1.15 T iHc = 1200 kA/m (BH)max = 262 kJ/m3. Compared to traditional powder metallurgy technology, the advantages of SPS are, (1) low sintering temperature, (2) ordinary materials sintering densifica-tion only need 3-5min, so has high sintering speed, (3) simple operation and short process flow, and (4) effectively inhibiting grain growth of NdFeB permanent magnetic alloy during sintering process (Xiao, Yue, Wang, et al, 2002).

As the world’s largest rare earth magnets supplier and market place, Chinese industry has a big demand for milling equipment, calibration service of magnetic properties and machinery processing equipment. Taiyuan Shengkaiyuan Perma-nent Magnet Equipment Corporation developed special equipment for crushing of NdFeB permanent magnet materials the rotary hydrogenated pulverizing fur-nace. That is an advanced pulverization technique and can effectively control shape and size of powder. Its main technical index reaches international ad-vanced level and can replace imported equipment. At present more than 120 equipments are used in China.

Beijing Xindake Electric Technology Co. Ltd. undertakes production of milling equipment, including coarse crusher, middle crusher and jet mill with nitrogen gas. Among them the QLMR-T series, whole sealing close circular loop jet mill, occupies 70% of domestic jet mill market. It has advantages of good sealing, cen-tralized distribution of powder size and less impurities.

China National Institute of Metrology (NIM) has developed a serials of com-mercial hysteresisgraphs, named NIM2000, for determination of magnetic proper-ties of hard magnetic, including hysteresisgraph for measuring hard ferrite and rare earth magnets, hysteresisgraph at high temperature(up to 500 ), and mag-netic field scanner for multi-pole magnets, etc. Nowadays, due to the traceable


accuracy, reliability, and price only a fraction of that of the developed countries, NIM2000 serials equipped nearly 90% NdFeB manufacturers in the world.

EDTA volume method was used to determine total rare earth contents in Dy-Fe alloy. Iron interference was eliminated by the addition of fluoride ions. Recovery in 99%-101% and RSD less than 1% were obtained. This method is accurate (Sun, Zhang, Zhang, et al, 2010).

Zhaoqing Dingchen Permanent Magnet Equipment Co. Ltd. is the specialty company producing machinery equipment to process permanent magnets. It has a special equipment development tram and obtained several patents. Its process technology is in leading position in China and reaches to advanced level of the world. It has various processing machines, including: digital controlled three working position inner circle slicer, auto-feed multi-position centreless grinder, digital controlled tile (C shape) grinder, vertical auto-feed dual work position arc grinder, auto-chamfer grinder, etc. Thus, its equipment is widely used in domestic market for manufacturing the rare earth permanent magnets and imported to part of overseas market, including Germany, Malaysia, India, etc.

Manufacture of bonded permanent magnetic alloys: in recent years the applica-tion fields of bonded permanent magnetic alloys has been enlarged continually in China. Among them the output growth rate of bonding permanent magnetic alloys with highly accurate size and complicated shape (bonded NdFeB permanent mag-netic alloy and bonded SmCo permanent magnetic alloy) grows fastest.

The commercialization of bond NdFeB permanent magnetic alloys started at 1980s. GM Corporation exploited raw material to make bonded NdFeB perma-nent magnetic alloy, and produced rapid quenching NdFeB magnetic powder in 1987. Bonded NdFeB permanent magnetic alloys have formed series products in past more than ten years. The price of raw material for bonded NdFeB permanent magnetic alloys is much cheaper than that of bonded SmCo permanent magnetic alloys. And magnetic properties of bonded NdFeB permanent magnetic alloys are higher, machinability is better, and material utilization rate is higher. So bonded NdFeB permanent magnetic alloys take most part of bonded rare earth permanent magnetic alloys. The yield of bonded rare earth permanent magnetic alloys is 3,538 ton in 2004 all over the world, where yield in Japan is 565 ton, in China is 1,350 ton.

The manufacturing method of bonded permanent magnetic alloys uses powder bonding method. Bonded NdFeB permanent magnetic alloys could be manufac-tured by compression forming, injection forming, extrusion forming, and calen-dering forming. The operation procedures of compression forming method are: mixing � forming � solidification � machining � coating � inspection. Although this method has many operation procedures, the production efficiency and material utilization are higher. The maximum magnetic energy product of bonded NdFeB permanent magnetic alloy produced by Aiwa Steel Group in Japan


was 119-159 kJ/m3. The maximum magnetic energy product and density of differ-ent bonded permanent magnetic alloys are shown in Table 5.3. The comparation of performances between anisotropic and isotropic bonded NdFeB permanent magnetic alloy is given in Table 5.4.

Table 5.3 Magnetic energy product and density of permanent magnetic alloy with

different bonding methods

(BH)max � Types of magnetic materials

kJ/m3 MGs·Oe g/cm3

Anisotropy SmFeN (injection moulding) 46-106 5.8-13.3 3.4-4.8

Anisotropy SmFeN ferrite (injection moulding) 32-56 4-7 3.8-4.4

Anisotropy SmFeN (compression forming) 56 7 6.2

Anisotropy SmFeN (injection moulding) 32 4 5.1

Anisotropy SmFeN (roll forming) 26-36 3.2-4.5 4.5-5.5 Anisotropy NdFeB (d-HDDR injection moulding of magnetic

powder) 119-159 15-20 4.8-5.2

Anisotropy NdFeB (MQP-B compression forming) 62-90 7.8-11.3 5.5-6.3

Anisotropy NdFeB (MQP-B injection moulding) 32-62 4-8 4.2-5.6

Anisotropy ferrite (injection moulding) 12-23 1.5-2.9 2.9-3.8

Anisotropy ferrite (roll forming) 10 1.2 3.5

Table 5.4 Comparison of magnetic performance of the permanent magnetic alloy between anisotropy and isotropy bonding

Anisotropy Isotropy Features

MF20 MF18 MF15 NEO10 NEO6

Forming process Compression Compression Injection moulding Compression Injection

moulding (BH)max/kJ·m-3(MGs·Oe) 159(20) 143(18) 119(15) 80(10) 48(6)

Br/T(kGs) 0.98(9.8) 0.93(9.3)

iHc/kA·m-1(kOe) 1,034(13) 1,512(19)

�/%· -1 0.13 0.13 0.13 0.11 0.11

�/%· -1 0.50 0.45 0.50 0.40 0.40

Tc/ 310 310 310 360 360 Ratio of price to mag-netic energy product 6.0 8.3 8.0 10.0 16.7

As shown in Table 5.4 anisotropic bonded NdFeB permanent magnetic alloy

has higher performance. Because of the development of bonded rare earth perma-nent magnetic alloy technology people could choose different bonded permanent magnetic alloys according to various demands.

In comparison with bonded NdFeB permanent magnetic alloy, sintering NdFeB permanent magnetic alloy has been developing in a high speed in recent years in China. The annual output of sintering NdFeB permanent magnetic alloy reached


83,000 ton in China in 2010. In the last few years the performance of sintering NdFeB permanent magnetic alloy has been improved remarkably as merchandise. Table 5.5 gives the maximum grades of sintering NdFeB permanent magnetic alloy produced by Sumitomo Corporation in Japan. The performances of sintering NdFeB permanent magnetic alloy produced by some Chinese companies reached international advanced standards with maximum magnetic energy product (BH)max = 376-408 kJ/m3 (47-51 MGs·Oe)(Luo, 2003).

Table 5.5 The highest trademarks of sintered NdFeB permanent magnetic alloy

manufactured in Japan Sumitomo Metal Corporate

Trademark (BH)max /kJ·m-3(MGs·Oe) Br/T iHc

/kA·m-1bHc

/kA·m-1 �

/%· -1�

/%· -1 �

/g·cm-3

NEOMAX-50 376-408(47-51) 1.39-1.45 880 840-1,040 0.11 0.59 7.5

NEOMAX-48DH 360-392(45-49) 1.36-1.42 1,120 1,024-1,104 0.11 0.58 7.5

NEOMAX-46H 344-376(43-47) 1.33-1.39 1,280 1,000-1,072 0.11 0.58 7.5

NEOMAX-39SH 288-320(36-40) 1.23-1.29 1,680 928-1,008 0.11 0.55 7.5

NEOMAX-38VH 288-320(36-40) 1.22-1.28 2,000 936-1,000 0.10 0.49 7.55

NEOMAX-35VH 264-296(33-37) 1.17-1.23 2,240 904-968 0.10 0.48 7.55

NEOMAX-32BH 240-272(30-34) 1.11-1.19 2,400 840-920 0.09 0.45 7.6

For many years the outstanding evolutions of NdFeB permanent magnetic

technology express in undermentioned three aspects: (1) aiming at low Curie temperature and big temperature coefficient of NdFeB alloy, the means to im-prove thermal stability of NdFeB magnet have been considered; (2) in order to overcome easy-rusting of iron base alloy in certain condition anticorrosion re-search have been carried out; (3) respond to requirement of users researches for different materials technique have been developed.

The newly developed heat-flow-transmutation procedures based on heat-pressing technique is an effective way to produce high performance nanocrystal Nd-Fe-B magnet. Process of heat-flow-transmutation is used to precipitate homo-geneous composite magnetic soft and hard phases of nanometer level from non-crystalline matrix and keeps coherence in interface. That meeting requirement for control microstructure of nanocrystal composite magnet as well as realizing mag-netic anisotropy synchronously provides a way for heat-flow transmutation to produce high performance anisotropic nanometer composite magnet.

5.4 Application and Expectation

The aim of developing new material is for application, but it takes a long time for a new material from exploit to use. And that NdFeB, the king of current perma-


nent magnetic materials, occupied permanent magnet market in a speed without precedent since it was born. Up till now through many years commercializing development it has proved that NdFeB is a very important permanent magnetic material with wide of application fields and large potential. The committee of the european communities has analysed the market of NdFeB and figured out in this report that NdFeB alloy should not only be the substitutor of all permanent mag-netic materials for matching existed devices, but also has been widely used in the new market of superseding electromagnetic and non-electromagnetic designed devices. Table 5.6 is the quantity demanded of NdFeB alloy based on market forecast. The actual quantity demanded is far more than the amount listed in the table. Table 5.7 show the application, distributing field and its variation of NdFeB in China, respectively (Pan, Li, 2000; Pan, Ping, Liu, et al, 2003; Luo, 2003; Pan, 2001).

Table 5.6 Demands for NdFeB alloy of industrial developed countries as

per application fields from 1990 to 1995

Market share/% Application Total/t·a-1

Civil Industry Medicine Office Communi-cation Log Test Transportation

Electric motor 960 10 65 15 10

Audio transducer 400 85 5 10

MRI 400 5 90 5 Magnetic sus-

pending 60 5 95

Breaker 120 35 45 10 10

Gripper 100 10 50 20 10 10

Printer 20 100

Bearing/coupling 35 80 5 10 5

Engine 1 10 70 20

Separator 7 14 85 1 Wave beam controller 7 10 50 40

Vacuum tube for microwave 1 80 10 10

Transistor 1 80 20

Switch/relay 5 60 20 10 10

Transfer sensor 15 40 60

Transducer 1 20 20 10 50

Total/t·a-1 2133 453 778 380 260 1 116 25 120 The 1990s, NdFeB/% 100 21 36 18 12 0 5 1 6

The 1990s, SmCo/% 100 25 20 15 10 10 10 5 5


Table 5.7 Application distribution and the variation of NdFeB alloy in China since 1988 (t)

Year 1988 1995 1998 1999 2000

Speaker 35(47%) 530(42%) 1,010(31%) 1,130(27%) 1,220(22%)

Dewax unit 15(20%) 130(10%) 230(7%) 250(6%) 270(5%)

Magnetic separator 12(16%) 180(14%) 390(12%) 420(10%) 440(8%)

Motor/breaker 6(8%) 160(13%) 460(14%) 630(15%) 890(16%)

Magnetic coupling unit 7(9%) 90(7%) 200(6%) 210(5%) 220(4%)

CD/DVD 40(3%) 590(18%) 1,010(24%) 1,660(30%)

Communication 60(5%) 260(8%) 420(10%) 670(12%)

Others (6%) (4%) (3%) (3%)

Total 75 1,260 3,260 4,200 5,550

The motors made up with neodymium, iron and boron have the advantages of

high efficiency, big unit power, light, small volume, etc. NdFeB permanent mag-netic motors and SmCo permanent magnetic motors do not need electric excita-tion, then they do not have excitation coil and iron core. Thus the volume and quality of these motors can be reduced by more than 30%. And that NdFeB per-manent magnet motors do not have excitation loss so that its efficiency is higher than common motors. Magnet itself does not glow thus in condition of same tem-perature rising NdFeB permanent magnetic motor has large allowable output power, so it has high efficiency. The 5.5 kW synchromotor used in colliery refit-ted by Northwest Technical University using rare earth permanent magnetic synchronous motor. Afterwards, its efficiency was improved from 74% to 83%, cos� was enhanced from 0.78 to 0.86. Extension of this motor will save a great deal of energy sources. General Motor Company succeeded in manufacturing starting motor of automobile with NdFeB permanent magnetic alloy and launched into small batch production in 1986. Benz Corporation also carried out trial-manufacturing successfully. Using permanent magnetic alloys to produce motors began in 1984 in China. In 1985, Northwest Technical University and cooperated with other organization successfully manufactured 1.5 kW startup motor for automobile with NdFeB permanent magnetic alloy. In 10W cyathiform armature motor AlNiCo permanent magnetic alloy was replaced by NdFeB permanent magnetic alloy. The weight of permanent magnet alloy reduced from 141g to 45g and the cost of permanent magnet cut down from 28 Yuan to 13 Yuan but the power was enlarged from 4W to 10W. The manufacturing and research in NdFeB permanent magnetic motor developed rapidly in recent years. Since 1987 NdFeB permanent magnetic alloy had been successfully applied in synchronous motor, servo-motor, DC motor and other two types of motors. With the enhancement of thermal stability of NdFeB permanent magnetic alloy, it will develop much faster


in the future. NdFeB permanent magnetic alloy have been successfully applied in MRI. Sub-

sidiary companies of Siemens in Germany and China have already sold instru-ments made with NdFeB permanent magnetic alloy.

The application and exploitation of NdFeB material have been developing con-tinuously. Examples of its application based on function show as follows.

Transforming mechanical energy to electrical energy: such as generator, pickup device, geophone, microphone, transmit speed instrument.

Transforming electrical energy into mechanical energy: small motors such as DC motor, trumpets in audio devices, moving-coil instrument, and electromag-netic pump.

Mechanical energy transformation: using attraction of magnet, magnetic cupule, magnetic force transport, magnetic separator, magnetic transmission, magnetic adsorption fixture, piston pump, etc. Using the attractive force and repulsive force between magnets: in magnetic axletree, magnetic puddler, magnetic suspension, magnetic clutch, etc. And using the induced eddy current function between per-manent magnet and conductor: wattmeter, velocimeter, and electric meter, eddy current driving device, etc.

The utilization of magnetic sources: NMR devices, magnetic field generator, dry reed relay, etc.

Deflexion of charged particle in magnetic field: magnetron, traveling-wave tube, electron gun controller, prionotron, electromagnetic flow meter, etc.

Magnetic biology: water magnetizer, magnetotherapy device. The application in automobile: one modern saloon car has more than 80 parts

using permanent magnet. The parts using neodymium-iron-boron permanent magnetic alloy are shown as follows: starter motor, cooling fan, portfire, front light, water meter, driving control device, glass rain scrape, flush pump, motor used in cover sunshine equipment, tape transmitting motor, horn, imitating fuzzy motor, motor used to lock door, retropack, motor used to open and close windows, motor used in fuel pump, motor used to control seats, antenna motor, air condi-tioner, crankshaft transport angle sensor, throttle sensors, velometer, etc.

NdFeB permanent magnetic alloy firstly used in automobile was starter motor, in the condition of coequal output torque, weight of motor which is made with NdFeB can be lightened 50%, volume can be reduced 30%-40%, and efficiency can be improved obviously. There are 60-80 motors in an automobile. If all mo-tors use NdFeB permanent magnetic motor, more than 5,000 ton bonded and sin-tering NdFeB permanent magnetic alloy would be needed (in demand for 50 mil-lion cars).

Application of NdFeB permanent magnetic alloy in voice coil motor ade-quately incarnates its advantages in its performance. Voice coil motor is an actua-tor of read-write head of disk drive. Because the trend of miniaturization and even


micromation of computer voice coil motor is required to be smaller, lighter, and have the magnetic capability as follows, Br=1.38T, iHc 1200kA/m, (BH)max

= 368kJ/m3. The used amount taking the weight of NdFeB permanent magnetic alloy in one voice coil motor record as 15g, and it needs thousands tons NdFeB permanent magnetic alloy to satisfying the requirement.

The biggest amount of NdFeB permanent magnetic alloy is used to manufac-ture permanent magnet motor all over the world. People most need high effi-ciency motor. Thus motor manufacturers make great efforts to enhance efficiency of motors. Motors with different structures have different efficiency (power coef-ficient). The highest efficiency of DC motor is 85%-89%, induction motor is 94%-95%, and permanent magnet motor is 95%-97%. Therefore, permanent magnetic motor will take the place of induction motor gradually. Bond NdFeB permanent magnetic alloy takes up the absolute advantage in multipolar spindle motor of hard disk and floppy drive. The required amount is more than a hundred million (Luo, 2002). Recently, according to the requirement of customer, we re-placed NdFeB permanent magnets with Ferrite, AlNiCo and SmCo permanent magnets by using combined magnetic circuit design methods to ensure the qualified magnetic properties of the components (Pan, 2011).

Great progress has been achieved in field of applying and developing NdFeB in China. The products with profit over 1 million have: antiwax apparatus, fuel economizer, MRI, magnetic separator, wattmeter, sensor, shock absorber, water mangetizer, magneto therapy apparatus, etc. Apparatuses such as magnetic field generator, magnetron, magnetic stirring, magnetic bearing, magnetic chuck, loud-speaker, etc. have been successfully exploited.

Electric bicycle which is widely used for transportation in middle-small cities get rapid development in recent years. Taking 2002 for example, one million elec-tric bicycles had been produced and nearly 600 ton NdFeB permanent magnetic alloy had been used. Forecasting that in 2007 electric bicycle and power assistant vehicle will need 5,700 ton and 2,385 ton NdFeB permanent magnetic alloy, re-spectively, electric vehicle will need 2,000 ton; total weight will need more than 10,000 ton in one year.

The third generation rare earth permanent magnetic alloy have come out for twenty years, its magnetic energy product has promoted from 286.6 kJ/m3 (36 MGs·Oe) at first to present 413.92 kJ/m3 (52 MGs·Oe) for commercial use. Its yield has increased from a few hundred kilogram to tens of thousands ton (Seen in Table 5.8). The yield of NdFeB sintering permanent magnetic alloy all over the world reached 34,510 ton in 2004. The total yield of bond and sintering perma-nent magnetic alloy is 28,860 ton in China, accounted for 75.26% of world ship-ment (Luo, 2003b). And the prospect of permanent magnetic industry is very good.


Table 5.8 Yield of sintered and bond NdFeB permanent magnetic alloy from 1985 to 2008 (t)

China Japan USA Europe Sum Year

Sintered Bond Sintered Bond Sintered Bond Sintered Bond Sintered Bond

1985 10 45 10 7 72

1986 20 97 1 22 4 16 155 5

1987 33 147 20 35 19 30 1 245 40

1988 75 320 100 75 60 50 5 765 165

1989 110 460 152 120 70 75 13 765 255

1990 180 910 250 230 40 110 20 1,430 310

1991 340 1,100 350 300 45 120 20 1,860 415

1992 490 1,200 460 350 60 130 15 2,170 535

1993 740 18 1,435 520 430 70 185 20 2,790 610

1994 1,230 40 1,555 710 645 90 250 40 3,680 840

1995 1,820 70 2,100 810 520 120 410 120 4,950 1,120

1996 2,100 100 2,600 890 640 140 510 200 5,850 1,330

1997 2,550 200 3,800 920 750 160 580 240 7,680 1,520

1998 3,260 300 4,500 1,180 710 180 630 280 9,100 1,840

1999 4,200 480 6,400 1,230 810 200 680 320 10,990 2,230

2000 5,600 700 5,100 700 950 400 750 350 13,700 2,900

2001 6,500 800 5,600 591 610 460 640 399 12,850 3,500

2002 8,800 1,140 6,200 500 280 460 580 350 15,260 3,660

2003 18,460 1,300 6,400 540 100 280 460 330 25,220 4,000

2004 22,910 1,350 6,700 565 210 300 345 30,710 3,836

2005 30,160 1,900 8,500 450 39,610 4,280

2006 38,200 2,800 10,500 1,080 49,800 5,070

2007 45,100 3,200 11,800 1,210 58,110 5,280

2008 52,400 4,200 13,000 1,100 63,580 6,000

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ZL.97104333.7 (in Chinese) Proceeding 7th Int’l Workshop on REPM, 1983

Appendix

Appendix 1 The Structure of Outer Electrons for Rare Earths

The structure of outer electrons for rare earths is listed in Table A.1.

Table A.1 The structure of outer electrons for rare earths

M N O P Ordinal number of

atom

Element symbol 3s 3p 3d 4s 4p 4d 4f 5s 5p 5d 6s

21 Sc 2 6 1 2

39 Y 2 6 10 2 6 1 2

57 La 2 6 10 2 6 10 2 6 1 2

58 Ce 2 6 10 2 6 10 1 2 6 1 2

59 Pr 2 6 10 2 6 10 3 2 6 2

60 Nd 2 6 10 2 6 10 4 2 6 2

61 Pm 2 6 10 2 6 10 5 2 6 2

62 Sm 2 6 10 2 6 10 6 2 6 2

63 Eu 2 6 10 2 6 10 7 2 6 2

64 Gd 2 6 10 2 6 10 7 2 6 1 2

65 Tb 2 6 10 2 6 10 9 2 6 2

66 Dy 2 6 10 2 6 10 10 2 6 2

67 Ho 2 6 10 2 6 10 11 2 6 2

68 Er 2 6 10 2 6 10 12 2 6 2

69 Tm 2 6 10 2 6 10 13 2 6 2

70 Yb 2 6 10 2 6 10 14 2 6 2

71 Lu 2 6 10 2 6 10 14 2 6 1 2


Appendix 2 Atomic and Ionic Radius of Rare Earths

Atomic and ionic radius of rare earths is listed in Table A.2.

Table A.2 Atomic and ionic radius of rare earths

Atom ordinal number Element symbol Radius of atom/nm Trivalent ion radius/nm

57 La 0.1877 0.1061

58 Ce 0.1825 0.1034

59 Pr 0.1828 0.1013

60 Nd 0.1821 0.0995

61 Pm (0.1810) 0.0979

62 Sm 0.1802 0.0964

63 Eu 0.2042 0.0950

64 Gd 0.1802 0.0938

65 Tb 0.1782 0.0923

66 Dy 0.1773 0.0908

67 Ho 0.1766 0.0894

68 Er 0.1757 0.0881

69 Tm 0.1746 0.0869

70 Yb 0.1940 0.0859

71 Lu 0.1734 0.0848

39 Y 0.1801 0.0880

21 Sc 0.1641 0.0680

Appendix

Appendix

Appendix 4 Fundamental Physical Constants

Fundamental physical constants is listed in Table A.4.

Table A.4 Fundamental physical constants

Physical quantity Symbol Value

Avogadro constant NA 6.023×1023mol 1

Boltzmann constant k 1.381×10 23J/K 8.62×10 5eV/K

Planck constant h 6.626×10 34J^s

Gas constant R 8.314J/(mol^K)

Faraday constant F 9.649×104C/mol

Vacuum dielectric constant �0 8.854×10 12F/m

Bohr magneton B 9.27×10 24A^m2

Speed of light in vacuum c 3×108m/s

Magnetic conductivity 0 4�×10 7H/m

Electric charge of electron e 1.602×10 19C

Mass constant of the atom mu 1.661×10 27kg

Mass of electron me 9.109×10 31kg

Mass of proton mp 1.673×10 27kg

Mass of neutron ma 1.675×10 27kg


Appendix 5 Conversion of magnetic quantity between SI and Gaussian units

Conversion of magnetic quantity between SI and Gaussian units is listed in Table A.5.

Table A.5 Conversion of magnetic quantity between SI and Gaussian units

Physical quantity Unit name and symbol of SI

Unit name and symbol of CGS

Conversion factor (value of CGS is obtained by using this

factor to time value of SI)

Length, l ( L ) Meter, m Centimetre, cm 102

Mass, m Kilogram, kg Gram, g 103

Force, F Newton, N Dyne, dyn 105

Moment, M N·m Dyn·cm 107

Work, W ( A ) Joule, J Erg, erg 107

Power, P Watt, W Erg/s 107

Pressure, p N/m2 Pascal, Pa Dyn/cm2 10

Density, � kg/m3 g/cm3 10-3

Current, I Ampere, A emu 10-1

Voltage, V Voltage, V emu 108

Inductance, L Henry, H emu 109

Resistance, R Ohm, _ emu 109

Magnetic field, H A/m Oersted, Oe 4�×10-3

Flux, � Weber, Wb Maxwell, Mx 108

Density of flux (magnetic induction), B

Wb/m2

Tesla, T Gauss, Gs 104

Magnetic polarization, J Wb/ m2 Gs 104/4�

Appendix

Continued Table A.5

Physical quantity Unit name and symbol of SI

Unit name and symbol of CGS

Conversion factor (value of CGS is obtained by using this

factor to time value of SI) Magnetization intensity, M

(�) A/m

Am2/kg Gs

emu/g 10-3

1

Magnetic pole strength, m Wb emu 108/4�

Magnetic dipole moment, jm Wb·m Magnetic torque 1010/4�

Magnetic moment, Mm A·m2 Magnetic torque 103

Magnetic potential, �m A·m2 Magnetic torque 103

Magnetometive force, Vm Ampere·turns, ATS Oe·cm 4�×10-1

Magnetization rate (comparative), � 1/4�

Magnetic conductivity (comparative), 1

Magnetic conductivity in vacuum, 0

4�×10-7 H/m 107/4�

Demagnetization factor (N= H/M) 4�

Magnetic resistance, Rm A/Wb Oe·cm/Mx 4�×10-9

Magnetic conductance, A Wb/A Mx/(Oe·cm) 109/4�

Density of energy, E Magnetic anisotropy

constant, K J/m3 erg/cm3 10

Gyromagnetic ratio m/(A·s) J/(Oe·s) 103/4�

Product of magnetic energy, (BH)m

J/m3

kJ/m3 Gs·Oe

MGs·Oe 4�×10

4�×10-2

Absolute magnetic conductivity, 0 ( ) H/m Gs/Oe 107/4�

Index

Composition design 133 Dynamic cross 1, 23, 25, 46, 162, 178, 179 Electron micrograph 41, 42, 48, 50, 61-64, 71, 72, 99, 100, 102-105, 111-113,

141-143, 146-149, 151, 152, 166-168, 179-181, 186-193, 197-200, 208-211 Generation 1-3, 27, 95, 129, 220, 231-235, 252 Hysteresis loop 65 77-82 In situ and dynamic observation 22, 32, 35, 41, 42, 52-54, 61-64, 66, 101, 111,

118, 129, 147, 151, 184, 195, 196, 198, 199, 201, 204-206, 216 Magnetic parameter (performance) 1, 7, 8, 23, 56, 74, 75, 82, 83, 85-87, 105,

133, 135, 139, 150, 155, 158, 161, 162, 170, 203, 215, 222, 224, 239-241 anisotropy 1-4, 8, 27-31, 33, 35, 38, 50-56, 58, 66, 67, 74, 115, 119, 120,

131, 136, 137, 140, 146, 157, 159, 163, 173, 176, 178-180, 184, 195, 198, 201, 208, 220, 221, 228, 233, 234, 238, 241, 245, 247

coercivity 1-4, 7-9, 21, 22, 24, 25, 27-40, 46, 51-55, 59, 65-70, 73-75, 80-84, 86-90, 95-101, 105-107, 110-118, 119-124, 126, 129-133, 135-141, 144-146, 154-156, 158, 160-163, 165, 166, 168-171, 173-175, 178-180, 182-184, 187, 188, 194, 196-202, 203-205, 208, 210, 212, 213, 215-217, 220-225, 232-241, 244, 245, 247, 248, 252

Curie temperature (Tc) 1-3, 8, 9, 14, 17, 22, 35, 36, 74, 96, 98, 121, 129-136, 138-140, 143-145, 147, 150, 153, 161-165, 168, 175, 180-183, 192, 193, 195, 201, 202, 204, 220-223, 213, 220-225, 231, 233-236, 238, 241, 243, 247, 248

demagnetization curve 2, 69, 83-86, 89, 90, 98, 106, 120-123, 136, 137, 158, 170, 171, 175, 202, 213, 223

magnetic energy product ((BH)max) 1, 2, 4, 8, 9, 56, 83, 86, 87, 90, 96, 98, 123, 124, 126, 129-131, 135, 136, 139-141, 144, 154-158, 163, 168, 170, 175, 178, 195, 201, 205, 215, 221-225, 231-241, 244, 245, 247, 248, 252

magnetic field (H) 1, 2, 7, 78-86, 97, 115, 162, 175, 183, 234 magnetization intensity (4M) 1, 7, 8, 51, 56, 78-82, 86, 87, 106, 121, 133,

161, 183, 234, 238 magnetic remanence (Br) 1, 2, 4, 7, 40, 56, 75, 83, 86, 87, 98, 123, 124, 126,

129, 131, 135-137, 139-141, 150, 154, 155, 158, 161, 163, 170, 171, 175, 178, 201, 205, 215, 218-220, 222-224, 232-241, 245, 247, 248, 252


temprerature coefficient(�) 2, 87, 96, 97, 117, 130, 131, 144, 220, 223, 225, 233, 236, 239-241

Mössbauer 33, 119, 121, 131, 133-135, 137-139, 144-147, 150, 153, 162-164, 175-177, 180, 183

Phase 1, 3, 4, 8-11, 14, 16, 19-21, 70, 132, 146, 185, 202, 223 B-rich phase 130, 132, 134, 139, 142, 143, 145, 147-156, 159, 163, 165,

172, 184, 187, 189-193, 195, 196, 201, 203, 205, 213, 214, 216 Ga-rich phase 143, 144, 166, 167, 180, 187, 188, 212 Laves phase 130, 132, 139, 144, 165, 179, 180, 183, 187, 196, 203 Matrix phase 20, 31, 32, 43-50, 52, 53, 64-67, 73, 74, 100, 107, 114, 139,

142, 143, 146-152, 159, 160, 166, 167, 173, 180, 182, 183, 188, 190-193, 197-201, 204-211, 213, 214, 221

Nd-rich phase 24, 130, 132, 134, 138, 139, 142-148, 150, 151, 155, 159-161, 163, 165, 171-174, 176-178, 181, 184-192, 195, 196, 199-201, 203-207, 209-214

Nowotny phase 154, 156, 214 RE2Co17 phase (2: 17 phase) 3, 9, 22, 28-34, 37, 38, 42-53, 58, 61-67, 70,

72-74, 95, 97-101, 110, 111, 115, 116, 119-121, 132, 146, 159-161, 220, 224, 225, 234, 239, 240

RE2Co7 phase (2: 7 phase) 9, 22, 28, 29, 31-34, 37, 46-50, 58, 61, 66, 70, 72-74, 240

RE2Fe14B phase (Nd2Fe14B phase etc.) 3, 6, 7, 23, 24, 129-132, 137-141, 143, 145-151, 154, 156, 158-161, 164, 167, 168, 173, 174, 176, 178-184, 186-199, 201, 203-205, 208, 209, 213, 214, 216, 220-221, 223, 235, 236, 241-253

SmCo5 phase (1: 5 phase) 3, 7, 9, 22, 25, 29, 31-33, 37, 44-48, 50-53, 58, 59, 61-67, 70, 72-74, 82, 90, 95, 97, 99, 105, 107, 110, 111, 115, 116, 119, 121, 126, 220, 224, 225, 232, 238-240

Phase transformation 1, 8-17, 19, 27, 30-32, 35, 42, 44-46, 48-51, 53-55, 64, 65, 68, 70, 95, 101, 110, 113-117, 129, 184, 189, 190, 193, 195, 196 belt 155, 162, 185, 188, 190-192, 195, 205-210, 215, 244 boundary 3-4, 18, 24-25, 28, 30, 40, 44-46, 50, 53, 64, 73, 94, 115, 130,

141-144, 148-149, 151, 154-156, 160-161, 166, 171, 173-174, 177-178, 180-183, 185-190, 194, 199-200, 204-205, 207, 209, 237

crystal 1, 4-7, 10, 17-18, 20, 22-24, 28, 30, 33, 35, 37-38, 40, 43-44, 46, 48-50, 52-54, 58, 66, 73, 95, 100, 111, 113, 118-121, 129, 141-144, 146-148, 150-152, 154, 156, 159-166, 169, 171, 173, 176, 178, 180, 183-191, 193-196, 198-201, 203-213, 215-217, 223, 232-234, 237

film (filmy) 17, 27, 31, 39, 41, 55, 61, 64, 98, 101, 118, 145, 147, 150-151, 155, 157, 162, 166, 185, 189-192, 195-196, 205-210, 215, 217

growth (growing up) 8, 9, 20, 21, 31, 42, 44-46, 49, 73, 101, 107, 109, 117,

Index

132, 152, 178, 179, 208, 216, interface 20, 21, 31, 34, 35, 44, 46, 110, 116, 166, 195, 199, 200, 205-207,

209, 212, 213, 216 nanocrystal 1, 3, 25, 174, 175 new phase 8, 19-21, 32, 43-46, 48-50, 55, 110, 116, 152, 176, 177 nucleation 22, 27-31, 33, 34, 42, 43, 45, 46, 48-52, 66, 73, 110, 114, 115,

116, 154, 163, 174, 178, 182, 184, 216 precipatation (precipitated phase) 28, 30, 33, 34, 46, 48, 53, 62-65, 67, 71-

73, 81, 82, 96, 97, 100, 101, 105, 107, 109, 111, 117, 118, 133, 146, 148, 150, 184, 190, 193-195, 199, 216

strip 31, 103-105, 109, 111, 141, 179-180, 192, 243 Rare earth permenent magmetic alloys 1, 2, 4, 8, 22-24, 27, 29, 82, 97, 129,

218, 231-239, 252 SmCo5 (1: 5 type) 1, 2, 4, 22, 25, 27, 28, 30-39, 41, 42, 50-90, 219, 220,

224, 232, 233, 235, 238-240, 242 Sm2Co17 (2: 17 type) 1, 3, 4, 25, 65, 90, 95, 96, 99, 100, 106, 107, 109, 111,

112, 121, 219, 224, 225, 234, 235, 239, 240, 242 Sm(Co, Cu, Fe, Zr)7.4 (Sm(Co, Cu, Fe, Zr)z) 95-100, 102-106, 111-118,

120-126, 220, 233, 240 rare earth iron permanent-magnet alloys (RE-Fe-B etc.) 1, 3, 22, 24, 129-

225, 234-238, 241-253 Stability 4, 89, 122, 130, 133, 217-225

chemical stability 224, 225 temperature stability 130, 217, 218 thermal stability 131, 133, 162, 183, 193, 195, 199, 204, 220, 225, 234, 236 time stability 217, 224, 225

Substitution 22 129-133, 135-141, 143-145 232 Temperature 11-15, 19, 21, 24, 27, 31, 33, 36-38, 40, 43, 48, 50-61, 63, 65, 66-

70, 73, 74, 78, 80, 81, 83-90, 98, 105-109, 115-119, 122, 126, 141, 153, 155, 156, 161, 184, 185, 190-192, 194, 195-200, 204, 205, 214-225, 233-240 aging temperature 97, 99, 104, 115, 121, 132, 144, 180, 200, 203, 213 annealing temperature 36, 37, 40, 56, 58-60, 101, 106, 185, 188 high temperature 1, 25, 28, 30, 31, 35, 42, 48, 51, 53, 59, 82, 96, 98, 101,

114, 121, 122, 129, 131, 149, 152, 172, 174, 184, 189, 190, 193, 195, 196, 200, 218, 220, 245

room temperature 27, 31, 36, 38, 41, 42, 53-55, 57, 58, 65, 68, 71, 73, 75, 76, 81, 83, 86, 97-101, 109, 111, 113, 117, 118, 126, 134, 146-148, 151-157, 161-163, 166-168, 175-177, 179-181, 183, 186-191, 194, 196, 197, 199, 220, 208, 211, 213, 214, 216-218, 220, 225, 232, 233, 235-237

tempering temperature 43

rare earth permanent-magnet alloysâ€™ high temperature phase transformation: in situ and...

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