High-Energy Ball Milling

High-energy ball milling

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High-energy ball milling Mechanochemical processing of nanopowders Edited by Małgorzata Sopicka-Lizer

Oxford

Cambridge

© Woodhead Publishing Limited, 2010

New Delhi

Published by Woodhead Publishing Limited, Abington Hall, Granta Park, Great Abington, Cambridge CB21 6AH, UK www.woodheadpublishing.com Woodhead Publishing India Private Limited, G-2, Vardaan House, 7/28 Ansari Road, Daryaganj, New Delhi – 110002, India www.woodheadpublishingindia.com Published in North America by CRC Press LLC, 6000 Broken Sound Parkway, NW, Suite 300, Boca Raton, FL 33487, USA First published 2010, Woodhead Publishing Limited and CRC Press LLC © Woodhead Publishing Limited, 2010 The authors have asserted their moral rights. This book contains information obtained from authentic and highly regarded sources. Reprinted material is quoted with permission, and sources are indicated. Reasonable efforts have been made to publish reliable data and information, but the authors and the publishers cannot assume responsibility for the validity of all materials. Neither the authors nor the publishers, nor anyone else associated with this publication, shall be liable for any loss, damage or liability directly or indirectly caused or alleged to be caused by this book. Neither this book nor any part may be reproduced or transmitted in any form or by any means, electronic or mechanical, including photocopying, microfilming and recording, or by any information storage or retrieval system, without permission in writing from Woodhead Publishing Limited. The consent of Woodhead Publishing Limited does not extend to copying for general distribution, for promotion, for creating new works, or for resale. Specific permission must be obtained in writing from Woodhead Publishing Limited for such copying. Trademark notice: Product or corporate names may be trademarks or registered trademarks, and are used only for identification and explanation, without intent to infringe. British Library Cataloguing in Publication Data A catalogue record for this book is available from the British Library. Library of Congress Cataloging in Publication Data A catalog record for this book is available from the Library of Congress. Woodhead Publishing ISBN 978-1-84569-531-6 (book) Woodhead Publishing ISBN 978-1-84569-944-4 (e-book) CRC Press ISBN 978-1-4398-2974-5 CRC Press order number: N10169 The publishers’ policy is to use permanent paper from mills that operate a sustainable forestry policy, and which has been manufactured from pulp which is processed using acid-free and elemental chlorine-free practices. Furthermore, the publishers ensure that the text paper and cover board used have met acceptable environmental accreditation standards. Typeset by Toppan Best-set Premedia Limited Printed by TJ International Limited, Padstow, Cornwall, UK

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Contents

1

Contributor contact details

xi

Introduction to mechanochemical processing M. Sopicka-Lizer, Silesian University of Technology, Poland

1

Part I Basic science of mechanochemistry 2

2.1 2.2 2.3 2.4 2.5 2.6 2.7 2.8 3

3.1 3.2 3.3 3.4 3.5

Mechanism and kinetics of mechanochemical processes F. Kh. Urakaev, Institute of Geology and Mineralogy SB RAS, Russia Introduction General issues Calculation of t–P–T conditions at point of contact of milling tools and treated particles Mechanism of mechanochemical processes Kinetics of mechanochemical processes Use of kinetic equations Conclusions References Kinetic behaviour of mechanically induced structural and chemical transformations G. Mulas, Università degli Studi di Sassari, Italy; and F. Delogu, Università degli Studi di Cagliari, Italy Introduction Mechanochemistry Outline of experimental methods Degradation of chlorinated aromatics Hydrogen absorption on Mg2Ni/Ni composite

7

9

9 15 16 22 27 32 40 40

45

45 46 48 50 53 v

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vi

Contents

3.6 3.7 3.8

Catalytic hydrogenation of carbon monoxide Future trends References

4

Materials design through mechanochemical processing M. Senna, Keio University, Japan Introduction Benefits of a mechanochemical route for materials design Theoretical aspects of modern mechanochemistry Composite oxides in electromagnetic applications Organic synthesis and utilization of spontaneous chemical reactions Conclusions and future trends References

4.1 4.2 4.3 4.4 4.5 4.6 4.7 5

5.1 5.2 5.3 5.4 5.5 5.6 5.7

Kinetic processes and mechanisms of mechanical alloying F. Delogu, Università degli Studi di Cagliari, Italy; and G. Mulas, Università degli Studi di Sassari, Italy Introduction Fundamentals of mechanical alloying processes in ball mills A phenomenological model of mechanical alloying kinetics Collecting and analysing experimental data on the kinetics of mechanical alloying reactions Mechanisms of chemical mixing on the atomic scale Future trends References

Part II Mechanochemical treatment of different materials 6

6.1 6.2 6.3

Mechanochemical synthesis of complex ceramic oxides T. Rojac and M. Kosec, Jozˇef Stefan Institute, Slovenia Introduction Mechanisms and kinetics of mechanochemical reactions in complex oxide systems Mechanochemical synthesis of complex oxides with various properties

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57 59 60

63 63 64 67 68 80 84 85

92

92 93 95 98 103 106 108

111

113 113 114 126

Contents 6.4 6.5

Future trends References

7

Production of intermetallic compound powders by a mechanochemical approach: solid–liquid reaction ball milling D. Chen, Y. Jiang, J. Cai and Z. Chen, Hunan University, People’s Republic of China; and P. Huang, Central South University, People’s Republic of China Introduction Experimental equipment and methods As-milled products obtained from solid–liquid reaction ball milling Reaction mechanism of solid–liquid reaction ball milling Conclusions Acknowledgement References

7.1 7.2 7.3 7.4 7.5 7.6 7.7 8

8.1 8.2 8.3 8.4 8.5 9

9.1 9.2 9.3 9.4 9.5 9.6

Mechanochemical processing of non-oxide systems with highly covalent bonds M. Sopicka-Lizer, Silesian University of Technology, Poland Introduction Manufacturing non-oxide powders by mechanochemical processing Mechanochemical processing of reactive systems Conclusion References Mechanochemical synthesis of metallic–ceramic composite powders K. Wieczorek-Ciurowa, Cracow University of Technology, Poland Introduction Composite powder formation: bottom-up and top-down techniques Monitoring mechanochemical processes Examples of applied high-energy milling in the synthesis of selected metallic–ceramic composite powders Copper-based composite powders with Al2O3 Nickel-based composite powders with Al2O3

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149

149 151 152 158 164 164 164

167

167 169 180 188 188

193

193 194 203 205 207 214

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Contents

9.7

Other possible variants of the synthesis of metal matrix– ceramic composites in Cu–Al–O and Ni–Al–O elemental systems using mechanical treatment ex situ and in situ Conclusions Acknowledgements References

9.8 9.9 9.10 10

10.1 10.2 10.3 10.4 10.5 10.6

Mechanochemical synthesis of organic compounds and rapidly soluble materials A.V. Dushkin, Institute of Solid State Chemistry and Mechanochemistry SB RAS, Russia Introduction Solid state mechanochemical synthesis of low-molecularweight organic compounds Effect of solid particles aggregation on reactivity in mechanochemical reactions Synthesis of rapidly soluble materials for pharmaceutical applications Future trends References

Part III Mechanochemical processes in metal powder systems and other applications 11

11.1 11.2 11.3 11.4 11.5 11.6 11.7 12

12.1

Mechanochemical plating and surface modification using ultrasonic vibration S.V. Komarov, Nippon Light Metal Co. Ltd, Japan; S.E. Romankov, N. Hayashi and E. Kasai, Tohoku University, Japan Introduction Key concepts Surface modification and coating treatment of as-obtained substrates Surface treatment of precoated substrates Production of nanostructured coatings Conclusions and future trends References Mechanochemically activated powders as precursors for spark plasma sintering (SPS) processes R. Orrù, R. Licheri, A.M. Locci and G. Cao, Università degli Studi di Cagliari, Italy Introduction

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217 219 220 220 224

224 225 229 235 244 245

249

251

251 253 255 260 267 272 273 275

275

Contents 12.2 12.3 12.4 12.5 12.6

Experimental materials and methods Results and discussion Conclusions and future trends Acknowledgements References

13

Synthesis of titanium dioxide-based, visible-light induced photocatalysts by mechanochemical doping S. Yin, Q. Zhang, F. Saito and T. Sato, Tohoku University, Japan Introduction Titanium dioxide and its photocatalytic applications Synthesis of anion-doped, titanium-dioxide, visiblelight induced photocatalysts by mechanochemical doping Future trends Acknowledgement References

13.1 13.2 13.3 13.4 13.5 13.6 14

14.1 14.2 14.3 14.4 14.5 14.6 14.7 14.8 15

15.1 15.2 15.3 15.4 15.5 15.6

Soft mechanochemical synthesis of materials for lithium-ion batteries: principles and applications N.V. Kosova, Institute of Solid State Chemistry and Mechanochemistry SB RAS, Russia Introduction: principles of soft mechanochemical synthesis of solid inorganic compounds Reactions of LiOH with anhydrous oxides Reactions of LiOH with solid hydroxides Reactions of LiOH and Li2CO3 with crystal hydrates and acidic salts Conclusions Future trends Acknowledgements References Materials for lithium-ion batteries by mechanochemical methods L.C. Yang, Q.T. Qu, Y. Shi and Y.P. Wu, Fudan University, China; T. van Ree, University of Venda, South Africa Introduction Lithium-ion batteries Mechanochemical preparation of cathode materials Mechanochemical preparation of anode materials Solid electrolytes from mechanochemical methods Conclusions

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ix 277 282 296 297 297

304

304 305 310 325 326 326

331

331 333 344 348 354 357 357 357

361

361 362 365 374 397 400

x

Contents

15.7 15.8

Acknowledgement References

401 402

Index

409

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Contributor contact details

(* = main contact)

Chapter 3

Editor and Chapter 1

G. Mulas* Dipartimento di Chimica Università degli Studi di Sassari via Vienna 2 07100 Sassari Italy E-mail: [email protected]

M. Sopicka-Lizer Faculty of Materials Engineering and Metallurgy Akademicka 2A Silesian University of Technology 44-100 Gliwice Poland E-mail: malgorzata.sopicka-lizer@ polsl.pl

Chapter 2 F. Kh. Urakaev Institute of Geology and Mineralogy SB RAS Acad. Koptyug Av., 3 Novosibirsk, 630090 Russia E-mail: [email protected]; [email protected]

F. Delogu Dipartimento di Ingegneria Chimica e Materiali Università degli Studi di Cagliari piazza d’Armi 09123 Cagliari Italy E-mail: [email protected]

Chapter 4 M. Senna Faculty of Science and Technology Keio University 3-14-1, Hiyoshi Yokohama 223-8522 Japan E-mail: [email protected]

xi © Woodhead Publishing Limited, 2010

xii

Contributor contact details

Chapter 5 F. Delogu* Dipartimento di Ingegneria Chimica e Materiali Università degli Studi di Cagliari piazza d’Armi 09123 Cagliari Italy E-mail: [email protected]

P. Huang School of Material Science and Engineering Central South University Changsha 410082 People’s Republic of China

Chapter 8

G. Mulas Dipartimento di Chimica Università degli Studi di Sassari via Vienna 2 07100 Sassari Italy E-mail: [email protected]

M. Sopicka-Lizer Faculty of Materials Engineering and Metallurgy Akademicka 2A Silesian University of Technology 44-100 Gliwice Poland E-mail: malgorzata.sopicka-lizer@ polsl.pl

Chapter 6

Chapter 9

T. Rojac* and M. Kosec Electronic Ceramics Department Jozˇef Stefan Institute Jamova cesta 39 1000 Ljubljana Slovenia E-mail: [email protected]

K. Wieczorek-Ciurowa Faculty of Chemical Engineering and Technology Cracow University of Technology 24 Warszawska 31-155 Cracow Poland E-mail: [email protected]

Chapter 7 D. Chen*, Y. Jiang, J. Cai and Z. Chen School of Material Science and Engineering Hunan University Changsha 410082 People’s Republic of China E-mail: [email protected]

Chapter 10 A.V. Dushkin Institute of Solid State Chemistry and Mechanochemistry SB RAS Kutateladze 18 Novosibirsk 630128 Russia E-mail: [email protected]

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Contributor contact details

xiii

Chapter 11

Chapter 13

S.V. Komarov* Nippon Light Metal Co. Ltd Nikkei Research & Development Center Kambara 161 Shimizu-ku Shizuoka-shi Japan E-mail: sergey-komarov@nikkeikin. co.jp

S. Yin*, Q. Zhang, F. Saito and T. Sato Institute of Multidisciplinary Research for Advanced Materials Tohoku University Katahira 2-1-1 Aoba-ku Sendai Japan E-mail: [email protected]

S.E. Romankov, N. Hayashi and E. Kasai Tohoku University, IMRAM Katahira 2-1-1 Aoba-ku Sendai-shi Japan

Chapter 12 R. Orrù* Dipartimento di Ingegneria Chimica e Materiali Università degli Studi di Cagliari piazza d’Armi 09123 Cagliari Italy E-mail: [email protected] R. Licheri, A.M. Locci and G. Cao* Dipartimento di Ingegneria Chimica e Materiali Università degli Studi di Cagliari piazza d’Armi 09123 Cagliari Italy E-mail: [email protected]

Chapter 14 N.V. Kosova Institute of Solid State Chemistry and Mechanochemistry SB RAS Kutateladze 18 Novosibirsk 630128 Russia E-mail: [email protected]

Chapter 15 L.C. Yang, Q.T. Qu, Y. Shi and Y.P. Wu* New Energy and Materials Laboratory (NEML) Department of Chemistry and Shanghai Key Laboratory of Molecular Catalysis and Innovative Materials Fudan University Shanghai 200433 China E-mail: [email protected] T. van Ree Department of Chemistry University of Venda Thohoyandou 0950 South Africa

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1 Introduction to mechanochemical processing M A Ł G O R Z ATA S O P I C K A - L I Z E R, Silesian University of Technology, Poland

Abstract: This chapter discusses the development of mechanochemical processing in various systems throughout its short history and shows the most impressive industrial application of the technique. A brief description of the complex action that takes place during powder mechanochemical processing is also included. Key words: crystal disordering, mechanochemical, metastable equilibrium, process feature.

Mechanochemical processing (MCP) can be defined as ‘a powder processing technique involving deformation, fracturing and cold welding of the particles during repeated collisions with a ball during high-energy milling’. Using mechanical energy to grind down various materials dates back to the beginning of human history and the application of flints to make a fire can be seen as an example of a mechanochemical treatment of materials. However, throughout the centuries, mechanical milling only involved diminution of particles without changing their structure and/or properties. The effect of chemical reactions initiated by mechanical action was found at the end of 19th century when Carey Lea first reported that the halides of gold, silver, platinum and mercury decomposed to halogen and metal during fine grinding in a mortar (Carrey Lea, 1893). His paper was published just after Ostwald introduced the term mechanochemistry in 1891. Heinicke’s much later definition (Heinicke, 1984) that ‘mechanochemistry is a branch of chemistry which is concerned with chemical and physicochemical transformation of substances in all states of aggregation produced by the effect of mechanical energy’ has been widely accepted. Since then the main applications of mechanochemical treatment of materials were found initially within the field of extractive metallurgy, where the process was used as a pretreatment step prior to leaching and extraction in order to increase the solubility of the minerals in question. The application of this rather simple technique to manufacturing advanced materials was driven by the industrial necessity to develop an alloy combining oxidedispersion strengthening (ODS) with γ′-precipitation hardening in a nickelbase superalloy intended for gas turbine applications. Benjamin’s work in the late 1960s showed that mechanochemical alloying (MA) must be used 1 © Woodhead Publishing Limited, 2010

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High-energy ball milling

when ordinary dispersion of oxide particles in liquid metal is not possible (Benjamin, 1970). The majority of the further work on development and production of ODS superalloys for application in the aerospace industry was done in the INCO laboratories in the USA (Ivanov and Suryanarayana, 2000). This is presently the major user of the MA process in the commercial production of nickel or iron-based high-temperature alloys which can be used at operating temperatures higher than 1300°C in carburizing or sulphidizing environments (Benjamin, 1970). Another spectacular example of commercial application where mechanochemically alloyed powders of Mg and Fe were used as heaters of MRE (meal, ready-to-eat), comes from their intensive use during the Desert Storm Operation by the USA in 1996 (Ivanov and Suryanarayana, 2000). On the other side of the world, in the former USSR, intensive studies on the application of mechanochemical processes in mineral and waste largescale processing resulted in several practical applications including treatment of tungsten-containing ore which was introduced at the Chirchik Plant or production of amalgam for children’s dentistry which was introduced at the Nikolsky Plant (Boldyrev, 2002). Development of unique grinding equipment in a variety of scales broadened the range of practical applications and new types of materials are being researched all over the world. Since 2000, more than 1500 scientific papers in the field of mechanochemistry have been published every year (Ivanov and Suryanarayana, 2000) and excellent reviews and monographs are available (Suryanarayana, 1999, 2004, 2008; Tkacova, 1989; Gutman, 1998; Arzt and Schultz, 1989; Barbadillo, 1993; Lai and Lu, 1998; Avvakumov, 1986). However, great developments in the technique itself and the accompanying milling devices have brought about new applications of MCP, broadening the range of the new materials and opening up new perspectives. There has been considerable interest in updating and summarizing existing knowledge of MCP. Mechanochemical processes use mechanical energy to activate chemical reactions and structural changes as well as particle size reduction. Under the action of cyclic loading, after breaking crystal bonds, MCP engages powder particles into a non-equilibrium state with a relaxation time of 10−7–10−3 s (Meyer and Meier, 1968). Comparison with other far from equilibrium processes (Table 1; Froes et al., 1995) shows that departure from equilibrium in MCP is faster than in rapid solidification. In addition, some long-lived defects with a lifetime of 10−3–106 s can be generated because of the solid imperfections and even if relaxed, the residual disorder remains in the activated material (Meyer and Meier, 1968). The mechanism of particle failure changes with the particle size and the structure of the particles undergoing grinding. A change in the accumulated energy relaxation from fracture to plastic deformation results in a dramatic

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Introduction to mechanochemical processing

3

Table 1.1 Departure from equilibrium achieved in various processes (Froes et al., 1995)

Process

Maximum departure from equilibrium (kJ/Na)*

Solid-state quench Quench from liquid (rapid solidification) Condensation from vapour Irradiation/ion implantation Mechanical cold work Mechanical alloying

16 24 160 30 11 30

Na, Avogadro's number. * Assuming relaxation owing to kinetic effects.

increase in the strain followed by an extreme dislocation flow. Accordingly, the elastic strain energy transforms into elastic energy in the lattice defects and into structural disordering or it can be relaxed by the fracture of brittle material or crystallographic lattice rearrangement in polymorphic transformation. If more than one component is under the action of ball/powder particles collisions, then relaxation occurs by mechanical alloying, decomposition or synthesis of a new chemical compound. Thus stress, deformation and fracture initiate changes in the solids while the type and capacity of the changes involved remain a function of material’s properties (lattice bonds and crystal structure, elastic modulus, particle surface properties) and stress conditions in a milling device (the magnitude and direction of acting forces, stress rate and frequency of loading). It appears that the increase in energy of the milled powder caused by the increased volume fraction of grain boundaries and lattice disordering raises the free energy above the level of the amorphous state. A unique feature of a mechanochemically activated mixture of powders is formation of numerous reaction couples which increase with decreasing particle size and regenerate through repeated particle fracture and welding events. What is even more important, the reaction product phase does not separate the reactants as happens in ordinary chemical reactors since continuous product phase removal takes place during ball/powder particles collisions. Lack of a diffusion barrier in the reacting couples as well as formation of various defects acting as the fast diffusion paths overcomes the problem of the diffusion as the rate-controlling process. Consequently, the reaction can proceed with acceptable kinetics without the necessity of raising the reaction temperature. Therefore, several solid phase chemical reactions or alloying reactions occur at ambient temperature during MCP or they could easily proceed for the duration of the subsequent thermal treatment if necessary.

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High-energy ball milling

In a summary of MCP’s features, the application of this technique to the manufacture of advanced materials provides an opportunity to make products with unique characteristics, which may be difficult to produce in any other way. These products have the following benefits: •

•

•

•

•

•

exceptionally high reactivity of the resultant crystalline powder particles, besides substantial size reduction, because of their higher free energy accumulated in the defective crystal lattice and increased grain boundary volume; significant decrease in crystalline grains at the nanometer level or transformation to an amorphous structure as a result of growing crystal cell disordering when the mechanochemical treatment is lengthened; extension of solid solubility over the range of the equilibrium owing to the formation of a metastable equilibrium between the terminal solid solution and an amorphous phase; formation of micro(nano)-structured dense composite particles with a homogeneous distribution of (nano)dispresoids if the various elastic properties of the mechanically activated particles are employed; nucleation of various composites and complex oxides at ambient temperature by formation of hetero-bridging bonds as a result of intimate mixing under controlled shear stressing in the dry solid process and short range atomic transfer across the boundaries of solids, which no other method, like photo-, magneto- or plasma chemical processes, could achieve; synthesis at room temperature of a variety of equilibrium solid compounds without transition phases and/or without the need to use solvents because numerous reaction couples are formed and there is continuous removal of diffusion barriers.

This book presents the application of MCP to a variety of systems and the latest advances, challenges and future trends are described. Each chapter has been written by an internationally well-established and experienced researcher. The book starts by dealing with kinetics problems that arise during the action of MCP; generalized equations are constructed followed by the presentation of experimental research on kinetics in the chosen systems. Then the effects of mechanochemical treatment in a variety of systems and special application of the activated novel materials are covered. Finally, the drawbacks of the method (contamination, production cost) are discussed.

References arzt e. and schultz l. (eds) (1989). New Materials by Mechanical Alloying Techniques, DGM Informationgesellschaft, Oberursel, Germany.

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Introduction to mechanochemical processing

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avvakumov e.g. (1986). Mechanical Methods of Activation of Chemical Processes, Nauka, Novosibirsk. de barbadillo j.j., froes f.h. and schwarz r. (eds) (1993). Mechanical Alloying for Structural Applications, ASM International, Materials Park, OH. benjamin j.s. (1970). ‘Dispersion strengthened super alloys by mechanical alloying’, Metall Trans V, 1(10), 2943–51. boldyrev v.v. (2002). ‘A historic view on the development of mechanochemistry in Siberia’, Chem Sust Develop, 10, 3–10. carrey lea m. (1893). ‘On endothermic reactions effected by mechanical force: Part I’, Am J Sci III, xlvi, 241. froes f.h., suryanarayana c., russel k.c. and ward-close c.m. (1995). Proceedings of the International Conference on Novel Techniques in Synthetics and Processing of Advanced Materials, Singh J. and Copley S.M. (eds), TMS, Warrendale PA, 1. gutman e. (1998). Mechanochemistry of Materials, Cambridge International Science Publishers, Cambridge, UK. heinicke g. (1984). Tribochemistry, Academy Verlag, Berlin. ivanov e. and suryanarayana c. (2000). ‘Materials and process design through mechanochemical routes’, J Mater Synth Process, 8(3–4), 235–44. lai m.o. and lu l. (1998). Mechanical Alloying, Kluwer Academic, Boston, Massachusetts. meyer k. and meier w. (1968). Kristal Technik, 3, 399. suryanarayana c. (1999). Non-equilibrium Processing of Materials, volume 2, Pergamon Materials Series, Oxford. suryanarayana c. (ed) (2004). Mechanical Alloying and Milling, Marcel Dekker, New York. suryanarayana c. (2008). ‘Recent developments in mechanical alloying’, Rev Adv Mater Sci, 18, 203–11. tkacova k. (1989). Mechanical Activation of Minerals, Elsevier, Amsterdam.

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2 Mechanism and kinetics of mechanochemical processes F. K h. U R A K A E V, Institute of Geology and Mineralogy SB RAS, Russia

Abstract: Parameters of the impact interaction between the milling tools and the material to be treated have been calculated for a series of mechanochemical reactors (comminuting devices, ball mills) using models based on non-linear elastoplastic theory for collision of solids. Possible mechanisms in the formation of nanocrystal particles and of chemical reactions taking place during crystallization have been studied. Generalized equations have been deduced for the kinetics rate constants of mechanochemical reactions and for the mechanical activation of substances in ball mills. Various examples of the application of the generalized kinetic equation to the calculations of ab initio rate constants of specific mechanochemical processes in ball mills are presented. The obtained theoretical values have been compared with experimental results. Key words: mechanochemistry, ball mills, mechanical activation, mechanochemical reactions, kinetics, theory, numerical modelling.

2.1

Introduction

The popularity and rapid development of investigations into mechanochemistry and mechanical activation (MA) (Urakaev and Boldyrev, 1999a,b,c, 2000a,b; Suryanarayana, 2001; Takacs, 2002; Koch, 2003; Miani and Maurigh, 2004; Butyagin and Streletskii, 2005; Kajdas, 2005; Butyagin, 2006; Urakaev and Shevchenko, 2007; Korchagin and Lyakhov, 2008; Urakaev, 2007, 2008, 2009) render actual character to the problem of the creation of models to describe the performance of devices used as mechanochemical reactors to predict the results of mechanical treatment of solids and to evaluate numerically the kinetics of mechanochemical processes (MP). In this context it is helpful to discuss the approaches to this problem recently proposed by Vasil’ev and Lomaeva (2003) and Butyagin (2003). In these works, the authors considered the athermic diffusion and deformation mechanisms of formation of cementite (Fe3C) upon the MA of an iron and carbon powder mixture. Note that the role of local heating caused by the impact–friction interactions of milling tools with particles of a treated 9 © Woodhead Publishing Limited, 2010

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material in the course of MA was completely ignored. This contradicts not only classical works, for example, by the Bowden-Tabor school (Bowden and Persson, 1961; Bowden and Tabor, 1974, 1986; Vick et al., 2000; Vick and Furey, 2001, 2003), but also all the practice of human history, which has aimed either to decrease heating of materials when they are treated mechanically or to use released heat, including primitive ways of making fire. The modelling of MP associated with arising short pressure or temperature pulses or the so-called t–P–T conditions (Urakaev and Boldyrev, 2000a) is built upon a rigorous thermodynamic basis. In my opinion, this basis can be complemented by invoking diffusion and deformation processes, but under no circumstances can be replaced by them. In the present work, on the basis of the models constructed for the interaction of milling tools with the substance under treatment, t–P–T conditions for different MP in ball mills are calculated. Urakaev and Boldyrev (2000b) show that the time factor (t) of the momentum of mechanical action upon the substance, along with traditional pressure and temperature characteristics, is especially important and the determining factor in the specific character of MP providing their difference from thermochemical and other methods used to perform solid state reactions. Known t–P–T conditions at the impact–friction point of contact of the particles to be treated allow traditional thermodynamic and physicochemical approaches to be used to consider the mechanism and to calculate the kinetics of MP in ball mills. Ball milling is characterized by a relative velocity W of the action of milling tools at the conjunction of particles under treatment (Fig. 2.1). For example, this case is realized in ball planetary mills (Fig. 2.2) characterized by the kinematic coefficient κ = ω/ω1 which is the ratio of the angular velocity of mill vial rotation ω1 to the angular velocity of the carrier rotation ω (scalar |ω| = 2πω, where ω is the frequency of the mill rotation); another characteristic of these mills is geometric, Γ = D/D1 which is the ratio of the distance D between the rotation axes to the radius D1 of the mill vial; Equation [2.1] follows (Urakaev, 2004, 2009): |W| = 2πωl1 [(κ + 1)2 + Γ2 − 2Γ(κ − 1) cos ϕ + (Γ + 1)2]0.5; Wn = |W| sin ϕ; Wt = |W| cos ϕ

[2.1]

where cos ϕ = −(1 + κ)/Γ determines the condition of a ball leaving the wall and where Wn and Wt are the normal and tangential components of the relative velocity, respectively. Now we shall characterize in brief the major direction of fundamental investigations into the theory of impact on solids. The most widely known is the elementary Newton’s theory of collisions based on the introduction of the collision coefficient at impact (Goldsmith, 1960). The wave theory of impact was first developed by Saint-Venant, Bernoulli, Poisson and Navier (Poisson, 1817; Kilchevsky, 1976; Kubenko, 1999). However, the methods

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© Woodhead Publishing Limited, 2010

ω

A1

Thermocouples

ω1

A Motor

ω

Vials

ω1

Water

Water

Differential thermocouple battery

Planetary mill type AGO-2

Test object

Capsule

2.1 Modern vibratory (including SPEX 8000) and planetary (AGO-2) mills and schemes for temperature, intensity measurement and calculations used for planetary (Chattopadhyay et al., 2001; Pustov et al., 2001; Mio et al., 2002) and vibratory (Delogu et al., 2000) mills.

T

D

D1

SPEX 8000 or vibratory mill

12

High-energy ball milling

ω1 D1

D

ω

W ϕ

2.2 Scheme for calculation of a planetary ball mill (AGO-2).

they used resulted in very slow convergence of the rows. A diametrically opposite theory was worked out by Hertz (1895). He assumed that the development of the impact process involves only small regions inside the bodies adjacent to the contact surface while other parts are not deformed during collision and move as absolutely rigid bodies. As a result, the problem of the collision of two elastic bodies was reduced to the problem of the collision of two material points, an elastic element being present between them (a spring or springs, Fig. 2.3). In the case of non-central collision of two elastic bodies, Routh’s ‘ξ-hypothesis’ was accepted (Routh, 1897; Maw et al., 1976; Keller, 1986; Stronge, 1990; Stoimenov, 1992). Accordingly, the connection between the tangential (It) and normal (In) momentums during the impact is formulated similarly to the Coulomb law of friction (Painlevé, 1895) It ≤ ξIn, where ξ is the dynamic friction coefficient (Fig. 2.4). The inequality sign is related to the cases when It is so small that slipping does not take place. If slipping occurs, the equality sign should be accepted. The essence of our approach to calculations in ball milling is best illustrated by the game of billiards. For example, to hit a ball into the right top billiard pocket under the Hertz theory it is necessary to strike the ball exactly at its centre (Fig. 2.4a). Clearly, the possibility of getting the balls

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Mechanism and kinetics of mechanochemical processes Wt

ft

W

Wn

Xi

i Ri,ci,λi,ρi,θi Li = 1/Ri

0

fn

fij

εj

2rij

fn εi

Wt +

Lj = 1/Rj Rj,cj,λj,ρj,θj j

Xj

t=0

13

εij = εi + εj

ξ ij

Wt tij = 2rij

Wn tij = 2εij 0 ≤ t ≤ tij

2.3 Kinematics and dynamics of impact–friction contact of two elastic bodies (or particles) i and j, selected arbitrarily from the lined layer, in the region πr 2δ of impact action produced by a ball, see also Fig. 2.4b and Fig. 2.5.

(a)

(b)

In

It

δ∗

In

I

R′

θ′

I

It

D >> R′~R˝ >> δ∗~ δ >> Ri

1

2Ri

W Wt

I t ≤ ξ In

δ 3

W

Wt

L = 1/D ~ 0; θ wall

Wn

Wn θ˝

R˝ 2

2.4 Game in billiards, (a) illustrates Hertz theory and Routh’s ‘ξ-hypothesis’ (Coriolis, 1835; Delassus and Peres, 1923; Peres, 1923, 1924; Horák, 1948; Johnson, 1985; Persson, 2000; Popov et al., 2002, Schmitz et al., 2003; Ismail and Stronge, 2008) application to calculation of ball milling. (b) Impact–friction contact of lined (1, 2) or unlined (2, see Fig. 2.5) balls.

into the correct position and of having a player with the ability to hit the ball directly at its centre is rare. Therefore in ball mills, as in billiards, noncentral interactions prevail as shown by the direction of the ball in the top middle billiard pocket. In this case the normal interactions of balls are described by Hertz theory and the tangents are described by the Routh ξ-hypothesis. When the interactions are equal friction comes into force and the time taken by the frictional interaction is identical to the time taken by the normal interaction of the balls in Hertz theory. The same is true for impacts of a ball with a billiard table wall.

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14

High-energy ball milling

Moving on to mechanochemistry, it is sufficient to cover balls and billiard table walls with a treated material (Fig. 2.4b) according to the known property of self-lining of a surface by milling tools. The accuracy of the calculations of the parameters of collision by this procedure will depend on how precisely we manage to measure the altered elastic properties of self-lined balls and walls. The synthesis of the wave theory and Hertz theory is usually performed in two directions. One of them (Goldsmith, 1960; Harris and Piersol, 1961; Panovko, 1977; Brach, 1989; Zegzhda, 1997) implies the expedient fulfillment of the boundary conditions at the edge of colliding bodies. The second (Timoshenko, 1955; Brennan, 1957–58; Horák and Pacáková, 1961; Panovko, 1985) is based on the combination of the theory of transverse vibrations of a rod (a body) and Hertz’s theory of contact interaction. This is a modern approach (approaches) involving Hertz’s theory and wave theory, but like the theory presented by Timoshenko (1955), is only connected with the studies performed by Saint-Venant (Kilchevsky, 1976). It is curious that the experimental confirmation of these theories is only found in the cases of rather small relative velocities of collision (W ≤ 10 m s−1, billiards) or rather high ones (W ≥ 1 km s−1, shells), when the energy released is sufficient to change the aggregate state of colliding bodies. In this case, the problem is reduced to the solution of hydrodynamic equations of movement for a plastically deformed solid (liquid) in another liquid. Classical Newton’s theory does not take account of the physical properties of colliding bodies. Because of this, its application in the calculations of comminuting devices (mechanochemical reactors) gives the desired results only in the calculation of the kinematics and dynamics of the movement of particles under treatment in comminuting devices in combination with Hertz’s theory. On the other hand, the application of modern concepts of collisions in calculations for ball mills seems not to be feasible practically owing to the complexity of the mechanochemical processes occurring in them. This is possible only in the framework of non-linear elastic or elastoplastic Hertz’s theory with minimum (hypothetical) application of wave theory. The impact of balls with each other and of balls with a wall have been studied in detail and within 10% of the difference between the experimental measurements and the calculated data can be coordinated with Hertz’s theory if the relative velocity W of collision does not exceed 10 m s−1. Relative velocities of interaction of milling tools in ball mills are approximately 10 m s−1 also, therefore the Hertz–Routh theory is the best available for modeling MP. Therefore, its main statements and the results which will be used to construct the model and to calculate the comminuting devices are listed below.

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Mechanism and kinetics of mechanochemical processes

2.2

15

General issues

Let us consider the collision of two solids, i and j (Fig. 2.3), without any connection to ball mills. In the framework of classical non-linear elastic Hertz theory, the connection between the impact force fij and the total deformation εij = εi + εj of the colliding bodies is accepted: fij = Bεij3/2. The coefficient B depends on the properties of materials of which the bodies are made and on the surface curvatures of the bodies Lk = 1/Rk, where Rk is the curvature radius at the contact point. The deviations in the calculations according to this theory increase with an increase in the relative normal velocity Wn of collisions, owing to the appearance of plastic deformations. More precise values (up to Wn = 30 m s−1) are obtained using the semi-empirical non-linear elastoplastic model (Batuev et al., 1969) in which it is accepted that fij = bεijn where b and n are constants determined experimentally. The results of calculations performed according to the non-linear elastic theory which is true up to Wn = 10 m s−1 are presented. The final results of calculations of the most important parameters have been presented in the most convenient form by Dinnik (1952) and Urakaev (2004) as follows: •

time tij of interaction of the two solids tij = (π/4)(10π)0.4{Wn−1(θi + θj)2(Li + Lj)[ρiρj/(ρiLj3 + ρjLi3)]2}0.2 [2.2]

•

maximum force fij of the interaction fij = (2/3)(10π)0.6{Wn6(θi + θj)−2(Li + Lj)−1 [ρiρj/(ρiLj3 + ρjLi3)]3}0.2

•

maximum stress σij in the centre of the contact area σij = 4(10/π4)0.2[Wn2(θi + θj)−4(Li + Lj)3ρiρj/(ρiLj3 + ρjLi3)]0.2

•

[2.5a]

maximum surface sij of the contact area, sij = πrij2 = (π/4)[(3/2) fij (θi + θj)/(Li + Lj)]2/3

•

[2.4b]

maximum radius rij of the contact area, rij = (1/2)[(3/2)fij(θi + θj)/(Li + Lj)]1/3

•

[2.4a]

mean mechanical stress <σij> at the contact point <σij> = 2σij/3 = (8/3π)(10π)0.2[Wn2(θi + θj)−4(Li + Lj)3 ρiρj/(ρiLj3 + ρjLi3)]0.2

•

[2.3]

[2.5b]

maximum total deformation εij of the bodies, εij = εi + εj = (1/4)[(9/4) fij2(θi + θj)2(Li + Lj)]1/3

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[2.6a]

16

High-energy ball milling εiθj = εjθi

and from Equation [2.6a] εi = εijθi/(θi + θj);

εj = εijθj/(θi + θj)

[2.6b]

Here ρi and ρj are the densities, θi = 4(1 − νi2)/Ei and θj = 4(1 − νj2)/Ej are the values of compliance coefficients, νi and νj are the Poisson coefficients, Ei and Ej are the values of Young’s modules for the solids i and j.

2.3

Calculation of t–P–T conditions at point of contact of milling tools and treated particles

When building up a calculation model of ball mills, the following facts were used (Urakaev and Boldyrev, 1999a, 2000a; Urakaev, 2004). 1

2

3

4

5

6

One ball (N = 1; hypothetical choice) is present in the mill vial to comminute the material (the procedure for other numbers of balls N > 1 will be described below). The velocity W of the impact action of balls, with radius R, upon the treated material is determined by a set of dimensional and dimensionless parameters describing the rotation (vibration) frequencies ωk and the linear dimensions of the mill Dm so that |W| = W(N, ωk, Dm) and Dm >> R, for example, see Equation [2.1]. At the initial time interval of treatment, a quasi-equilibrium size distribution of the particles to be ground is achieved which is described by the mean radii Ri (i = 1, 2, . . . , j is the number of components of the material to be treated). For example (Schneider, 1968), in case i = 1 in the treatment of model materials (sodium chloride or quartz): R1(NaCl) = 16 × 10−5 cm and R1(SiO2) = 3 × 10−5 cm. A phenomenon involving self-lining (a layer of compact powder will be formed on the surface of milling tools) occurs (Kobayashi, 1995; Belyaev et al., 1998; Revesz and Takacs, 2007; Takacs and Torosyan, 2007), the thickness of the lining layer being δ (Fig. 2.4b and Fig. 2.5). In this case, Ri << δ∼δ* << R. This material can be characterized by its eigen mechanical properties (E, ν and others; see Anderson, 1965). A ‘hindered’ impact action of the ball load on the particles of the material to be treated is performed either during a collision of an unlined or a lined ball with the near-wall lined layer of particles, or during the collision of lined balls with each other. The parameters of these collisions are described by Hertz’s theory. The impact of the interaction of the particles being treated occurs as a result of stress wave propagation, in the region πr2δ of the impact of the balls (Fig. 2.5) which is determined by the conditions listed in (5). The parameters of these collisions are also described by Hertz’s theory. Only

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Mechanism and kinetics of mechanochemical processes (a) ∗

δ 1

2Ri

R

Ri << δ∗~ δ << R << Dm <σ> Wt r ρ

Wn

W

2' 2r

(b) 2Ri

δ∗

1 R

L

θ

Wt ft ξ

δ θ

17

Wn 2'

2 W

πr 2δ

Wall

2.5 Impact of milling tools on the layer of particles lining the near-wall being treated in a ball mill: (a) lined (1) or unlined (2′) ball with lining the inner wall; (b) lined (1) or unlined (2′) ball with the lined ball (2). The diagram also shows the friction (ft, ξ) contact of lined or non-lined balls with the layer of treated material lining the wall.

<σ>

V 2.6 Stress wave propagation in the region πr2δ of impact of the milling tools (between the balls and between a ball and the wall).

the first pass of the stress wave is taken into account (Fig. 2.6). This is connected with the fact (Love, 1927) that approximately equal velocities of compression wave propagation in the ball ν = [E(2 − 5ν + 3ν2)/ρ(2 − 5ν + 2ν2)]0.5 ≈ (E/ρ)0.5 and in the layers of particles being treated determine their periods of vibration (differing by approximately ∼R/δ, ∼R/Ri and ∼δ/Ri times). It is evident that under this assumption the times of impact interaction according to point 5 above (see Fig. 2.4b and Fig. 2.5) will be much longer than the time of impact interaction of particles being treated (Fig. 2.3, Fig. 2.6 and Fig. 2.7).

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18

High-energy ball milling

V*

R1

s12 R2

1

2

d* V = V1 + V2

2.7 Impact–friction contact of two particles 1 and 2 (in calculating the kinetics of MP). V and V* (= sd*) are volumes of particles near the contact area, s is the contact area and d is the thickness.

Conditions 3 and 4 are sufficient for calculations to be correct, but are not obligatory in principle. In the most general case, it is possible to choose the material surrounding the ball so that it will not be lined by the substance to be treated. Then, if we use the Equations [2.2]–[2.6], which describe the impact of the interaction of a ball with the layer of particles lining the inner wall (Fig. 2.5), it should be remembered that from the condition described in point 2 above, Li = 1/Ri >> 1/Dm = Lj = L, the inner wall layer can be treated in practical calculations of the collision with a ball with a density ρ = ρi, compliance θ = θi and radius R = 1/Li as a plane with curvature Lj = L = 0 and compliance θ = θj = θ. Using Equations [2.2]–[2.6] for the impact interaction of lined balls with a plane inner wall layer of the treated material, in accordance with point 4 above, one should accept the following values of the radius, density and compliance of the lined (*) balls: R* = 1/ Li ≈ R, ρ* = ρi ≈ ρ and θ* = θj = θ. For the impact of lined balls with each other, one should accept Li = Lj = L* = 1/R* ≈ 1/R, ρi = ρj = ρ* ≈ ρ and θi = θj = θ* ≈ θ in Equations [2.2]–[2.6]. The resulting equations are presented in Tables 2.1 and 2.2 where it is recommended that θ (Anderson, 1965) be described using the arithmetic means of compliances θi of the particles of different components of the treated mixture that make up the layer for compliance of the lined near-wall layers. For example, for a single-component substance, θ = θ1, for a twocomponent mixture, θ ≈ (θ1 + θ2)/2, and so on. Both during collision in milling tools (W) and under impulse pressure or stress wave propagation (<σ> = P) at the conjunction of particles under treatment in the ball mills, the impact velocity W or force fij that acts at the point of contact between the milling tools and the particles can be resolved into normal (Wn, fn) and tangential (Wt, ft) constituents (Figs 2.3–2.5). As a result, during a characteristic time interval τ* of the action of the momen-

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s12* = 1.8 (L1 + L2)−2ρ0.4θ−1.6(θ1 + θ2)2W0.8 r12* = 0.75 (L1 + L2)−1ρ0.2θ−0.8(θ1 + θ2)W0.4 f12* = 2.3 (L1 + L2)−2ρ0.6θ−2.4(θ1 + θ2)2W1.2 ε12* = 0.57 (L1 + L2)−1ρ0.4θ−1.6(θ1 + θ2)2W0.8 wn* = 0.28 [(ρ1L23 + ρ2L13)/ρ1ρ2(L1 + L2)3]0.5 × ρ0.5θ−2(θ1 + θ2)2W → (1300 cm s−1) t12* = 4.0[ρ1ρ2(L1 + L2)/(ρ1L23 + ρ2L13)]0.5 × ρ−0.1θ0.4W−0.2 → (1.4 × 10−9 s) wt* = 0.38 [(ρ1L23 + ρ2L13)/ρ1ρ2(L1 + L2)3]0.5 × ρ0.3θ−1.2(θ1 + θ2)W0.6 → (610 cm s−1) σn* = 1.5 ρ0.2θ−0.8W0.4 → (5.1 × 1010 dyn cm−2) σt* = 1.7 ρ0.12θ−0.48(θ1 + θ2)−0.4W0.24 σt* = 9.6 × 1010 dyn cm−2; ε12* = 6.8 × 10−7 cm; f12* = 1.4 dyn; r12* = 3.2 × 10−6 cm; s12* = 3.2 × 10−11 cm2.

t* = 3.6 Rρ0.4θ0.4W−0.2 → (1.54 × 10−5 s) f* = 2.3 R2ρ0.6θ−0.4W1.2 → (6.2 × 107 dyn) σ* = 1.9 ρ0.2θ−0.8W0.4 → (6.6 × 1010 dyn cm−2) <σ∗> = 1.3 ρ0.2θ−0.8W0.4 → (4.4 × 1010 dyn cm−2) r* = 0.75 Rρ0.2θ0.2W0.4 → (2.1 × 10−2 cm) s* = 1.8 R2ρ0.4θ0.4W0.8 → (1.4 × 10−3 cm2) ε* = 1.1 Rρ0.4θ0.4W0.8 → (4.36 × 10−3 cm)

Ball*–ball*

When studying the kinetics of reactions in mixtures of solids, we did not observe the lining of the balls made of WCo8 alloy. So, numerical values for different ΔT(x, t) sections are presented for the collisions of the unlined ball with the near-wall layer: ΔT(x, t) ≈ 1.0 ξ12W(c1c2λ1λ2ρ1ρ2)−0.25(θ + θ)−2(θ1 + θ2)[ρ(ρ1L23 + ρ2L13)/ρ1ρ2(L1 + L2)3]0.5 × {t0.5iErfc[x/2(a1a2)0.25t 0.5] − (t − t12)0.5iErfc[x/2(a1a2)0.25(t − t12)0.5]}; ΔT(x = 0; t = t12) ≈ 1300 K; ΔT(x = 0.5ε12; t = t12) ≈ 1100 K.

a

s12 = 3.1 (L1 + L2)−2ρ0.4(θ + θ)−1.6(θ1 + θ2)2W0.8 r12 = 1.0 (L1 + L2)−1ρ0.2(θ + θ)−0.8(θ1 + θ2)W0.4 f12 = 5.3 (L1 + L2)−2ρ0.6(θ + θ)−2.4(θ1 + θ2)2W1.2 ε12 = 1.0 (L1 + L2)−1ρ0.4(θ + θ)−1.6(θ1 + θ2)2W0.8 wn = 0.57 [(ρ1L23 + ρ2L13)/ρ1ρ2(L1 + L2)3]0.5 × ρ0.5(θ + θ)−2(θ1 + θ2)2W → (1600 cm s−1) t12 = 3.5 [ρ1ρ2(L1 + L2)/(ρ1L23 + ρ2L13)]0.5 × ρ−0.1(θ + θ)0.4W−0.2 → (1.0 × 10−9 s) wt = 0.57 [(ρ1L23 + ρ2L13)/ρ1ρ2(L1 + L2)3]0.5 × ρ0.3(θ + θ)−1.2(θ1 + θ2)W0.6 → (690 cm s−1) σn = 2.0 ρ0.2(θ + θ)−0.8W0.4 → (5.6 × 1010 dyn cm−2) σt = 2.0 ρ0.12(θ + θ)−0.48(θ1 + θ2)−0.4W0.24 σt = 10 × 1010 dyn cm−2; ε12 = 8.9 × 10−7 cm; f12 = 1.8 dyn; r12 = 3.5 × 10−6 cm; s12 = 3.8 × 10−11 cm2.

s12* = 1.0 (L1 + L2)−2ρ0.4θ−1.6(θ1 + θ2)2W0.8 r12* = 0.57 (L1 + L2)−1ρ0.2θ−0.8(θ1 + θ2)W0.4 f12* = 1.0 (L1 + L2)−2ρ0.6θ−2.4(θ1 + θ2)2W1.2 ε12*= 0.33 (L1 + L2)−1ρ0.4θ−1.6(θ1 + θ2)2W0.8 wn*= 0.14 [(ρ1L23 + ρ2L13)/ρ1ρ2(L1 + L2)3]0.5 × ρ0.5θ−2(θ1 + θ2)2W → (630 cm s−1) t12*= 4.6 [ρ1ρ2(L1 + L2)/(ρ1L23 + ρ2L13)]0.5 × ρ−0.1θ0.4W−0.2 → (1.6 × 10−9 s) wt* = 0.25 [(ρ1L23 + ρ2L13)/ρ1ρ2(L1 + L2)3]0.5 × ρ0.3θ−1.2(θ1 + θ2)W0.6 → (400 cm s−1) σn* = 1.2 ρ0.2θ−0.8W0.4 → (4.1 × 1010 dyn cm−2) σt* = 1.45 ρ0.12θ−0.48(θ1 + θ2)−0.4W0.24 σt* = 8.2 × 1010 dyn cm−2; ε12* = 3.9 × 10−7 cm; f12* = 0.6 dyn; r12* = 2.4 × 10−6 cm; s12* = 1.8 × 10−11 cm2.

t* = 4.1 Rρ0.4θ0.4W−0.2 → (1.76 × 10−5 s) f* = 4.0 R2ρ0.6θ−0.4W1.2 → (11 × 107 dyn) σ* = 1.5 ρ0.2θ−0.8W0.4 → (5.1 × 1010 dyn cm−2) <σ∗> = 1.0 ρ0.2θ−0.8W0.4 → (3.4 × 1010 dyn cm−2) r* = 1.1 Rρ0.2θ0.2W0.4 → (3.2 × 10−2 cm) s* = 4.1 R2ρ0.4θ0.4W0.8 → (3.2 × 10−3 cm2) ε* = 1.3 Rρ0.4θ0.4W0.8 → (5.15 × 10−3 cm)

t = 3.1 Rρ0.4(θ + θ)0.4W−0.2 → (1.46 × 10−5 s) f = 5.3 R2ρ0.6(θ + θ)−0.4W1.2 → (13 × 107 dyn) σ = 2.5 ρ0.2(θ + θ)−0.8W0.4 → (7.2 × 1010 dyn cm−2) <σ> = 1.7 ρ0.2(θ + θ)−0.8W0.4 → (4.8 × 1010 dyn cm−2) r = 1.0 Rρ0.2(θ + θ)0.2W0.4 → (2.9 × 10−2 cm) s = 3.1 R2ρ0.4(θ + θ)0.4W0.8 → (2.7 × 10−3 cm2) ε = 1.0 Rρ0.4(θ + θ)0.4W0.8 → (4.35 × 10−3 cm)

Particle interaction in near-wall lined layers

Ball*–flat layer

Ball–flat layer

Table 2.1 Calculating expressions. Numerical values for the treatment of an equimolar mixture of calcite and quartz in a KhK 871 mill (ω = 11 s−1, Wn = W = 9.2 m s−1, balls from WCo8 alloy are given in parentheses). An asterisk (*) corresponds to the lined ball loada

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s1* = 1.8 R12ρ0.4θ10.4W0.8 → (9.5 × 10−10 cm2) r1* = 0.75 R1ρ0.2θ10.2W0.4 → (1.7 × 10−5 cm) f1* = 2.3 R12ρ0.6θ1−0.4W1.2 → (17 dyn) ε1* = 1.1 R1ρ0.4θ10.4W0.8 → (3.6 × 10−6 cm) wn* = 0.57 (ρ/ρ1)0.5W → (910 cm s−1) t1* = 4.0 R1ρ10.5ρ−0.1θ10.4W−0.2 → (8.1 × 10−9 s) wt* = 0.38 ρ1−0.5ρ0.3θ1−0.2θ1W0.6 → (4300 cm s−1) σn* = 1.5 ρ0.2θ1−0.8W0.4 → (2.1 × 1010 dyn cm−2) σt* = 1.3 ρ0.12θ1−0.88W0.2 → (4.0 × 1010 dyn cm−2) 2q1* = ξ1 <σ*> wt → (6.3 × 1013 erg cm−2 s−1).

t* = 3.6 Rρ0.4θ10.4W−0.2 → (1.7 × 10−5 s) f * = 2.3 R2ρ0.6θ1−0.4W1.2 → (2.6 × 107 dyn) σ* = 1.9 ρ0.2θ1−0.8W0.4 → (2.7 × 1010 dyn cm−2) <σ∗> = 1.3 ρ0.2θ1−0.8W0.4 → (1.8 × 1010 dyn cm−2) r* = 0.75 Rρ0.2θ10.2W0.4 → (2.1 × 10−2 cm) s* = 1.8 R2ρ0.4θ10.4W0.8 → (1.4 × 10−3 cm2) ε* = 1.1 Rρ0.4θ10.4W0.8 → (4.5 × 10−3 cm)

Ball*–ball*

In the experiment on the treatment of NaCl particles, balls are lined only insignificantly, so the numerical values for different sections of temperature pulses are presented for the impact interaction of an unlined ball with the flat near-wall layer (T0 = 375 K): ΔT(x, t) ≈ 0.96 ξ1Wρ0.5(c1λ1ρ1)−0.5(θ + θ1)−2θ1{t0.5iErfc[x/2(a1t)0.5] − (t − t1)0.5iErfc[x/2(a1(t − t12)0.5]}. At the rise of ΔT: Tm − T0 = ΔTm = ΔT(x = 0; t = tm) = ΔT(x = xm; t = t1); tm ≈ 2.5 × 10−9 s; iErfc[xm/2(a1t)0.5] = 0.5642(tm/t1)0.5 and xm ≈ 8.7 × 10−6 cm. At the drop of ΔT: ΔTm = ΔT(x = 0; t = tm′); tm′ = (t1 + tm)2/4tm ≈ 1.0 × 10−8 s. ΔT(x = 0; t = t1) ≈ 1220 K; ΔT(x = 0.5ε1; t = t1) ≈ 1050 K; Δtm = tm′ − tm ≈ 8.0 × 10−9 s.

a

s1 = 3.1 R12ρ0.4(θ + θ1)−1.6θ12W0.8 → (1.3 × 10−10 cm2) r1 = 1.0 R1ρ0.2(θ + θ1)−0.8θ1W0.4 → (2.0 × 10−5 cm) f1 = 5.3 R12ρ0.6(θ + θ1)−2.4θ12W1.2 → (27 dyn) ε1 = 2.0 R1ρ0.4(θ + θ1)−1.6θ12W0.8 → (5.1 × 10−6 cm) wn = 1.1 (ρ/ρ1)0.5(θ + θ1)−2θ12W → (1000 cm s−1) t1 = 3.5 R1ρ10.5ρ−0.1(θ + θ1)0.4W−0.2 → (7.6 × 10−9 s) wt = 0.57 ρ1−0.5ρ0.3(θ + θ1)−1.2θ1W0.6 → (5400 cm s−1) σn = 2.0 ρ0.2(θ + θ1)−0.8W0.4 → (2.5 × 1010 dyn cm−2) σt = 1.5ρ0.12(θ + θ1)−0.48θ1−0.4W0.2→ 4.3 × 1010 dyn cm−2 2q1 = ξ1 <σ> wt → (9.1 × 1013 erg cm−2 s−1).

s1* = 1.0 R12ρ0.4θ10.4W0.8 → (5.3 × 10−10 cm2) r1* = 0.57 R1ρ0.2θ10.2W0.4 → (1.3 × 10−5 cm) f1* = 1.0 R12ρ0.6θ1−0.4W1.2 → (7.4 dyn) ε1* = 0.66 R1ρ0.4θ10.4W0.8 → (2.2 × 10−6 cm) wn* = 0.28 (ρ/ρ1)0.5W → (450 cm s−1) t1* = 4.6 R1ρ10.5ρ−0.1θ10.4W−0.2 → (9.3 × 10−9 s) wt* = 0.25 ρ1−0.5ρ0.3θ1−0.2θ1W0.6 → (2800 cm s−1) σn* = 1.2 ρ0.2θ1−0.8W0.4 → (1.7 × 1010 dyn cm−2) σt* = 1.1 ρ0.12θ1−0.88W0.2 → (3.4 × 1010 dyn cm−2) 2q1* = ξ1 <σ*> wt → (3.1 × 1013 erg cm−2 s−1).

t* = 4.1 Rρ0.4θ10.4W−0.2 → (1.9 × 10−5 s) f* = 4.0 R2ρ0.6θ1−0.4W1.2 → (4.4 × 107 dyn) σ* = 1.5 ρ0.2θ1−0.8W0.4 → (2.1 × 1010 dyn cm−2) <σ∗> = 1.0 ρ0.2θ1−0.8W0.4 → (1.4 × 1010 dyn cm−2) r* = 1.1 Rρ0.2θ10.2W0.4 → (3.2 × 10−2 cm) s* = 4.1 R2ρ0.4θ10.4W0.8 → (3.3 × 10−3 cm2) ε* = 1.3 Rρ0.4θ10.4W0.8 → (5.2 × 10−3 cm)

t = 3.1 Rρ0.4(θ1 + θ2)0.4W−0.2 → (1.6 × 10−5 s) f = 5.3 R2ρ0.6(θ1 + θ2)−0.4W1.2 → (5.5 × 107 dyn) σ = 2.5 ρ0.2(θ1 + θ2)−0.8W0.4 → (3.1 × 1010 dyn cm−2) <σ> = 1.7 ρ0.2(θ1 + θ2)−0.8W0.4 → (2.1 × 1010 dyn cm−2) r = 1.0 Rρ0.2(θ1 + θ2)0.2W0.4 → (2.9 × 10−2 cm) s = 3.1 R2ρ0.4(θ1 + θ2)0.4W0.8 → (2.6 × 10−3 cm2) ε = 1.0 Rρ0.4(θ1 + θ2)0.4W0.8 → (4.2 × 10−3 cm)

Particle interaction in near-wall lined layers

Ball*–flat layer

Ball–flat layer

Table 2.2 Calculating expressions. Numerical values for the treatment of a model single-component material (sodium chloride) in a EI 2 × 150 mill (ω = 10 s−1, Wn = W = 8.4 m s−1, steel balls are given in parentheses). An asterisk (*) corresponds to the lined ball loada

Mechanism and kinetics of mechanochemical processes

21

tum, deformation of milling tools occurs with a shift which leads to an increase in temperature T at the point of contact (Fig. 2.4b and Fig. 2.5). For example, during impact–friction contact of the unlined ball with the lined layer wall (Fig. 2.5), the following expressions were obtained: •

time t of interaction of the milling tools (see Table 2.1, column 1, line 1) t = 3.1 Rρ0.4(θ + θ)0.4Wn−0.2

•

mean mechanical stress <σ> or pressure P at the contact point <σ> = P = 1.7 ρ0.2(θ + θ)−0.8Wn0.4

•

[2.7]

[2.8]

density of heat source 2Q and temperature impulse ΔT(X, t) in the vicinity of contact area X 2Q = ξPWt; ΔT(X, t) = 2Q (cc*λλ*ρρ*)−0.25 × {t0.5 iErfc[X/2(aa*)0.25t0.5] − (t − t)0.5 iErfc[X/2(aa*)0.25(t − t)0.5]}; ΔT(max) = ΔT(0, t); iErfc [0] = 0.5642

[2.9]

Here ξ is the dynamic friction coefficient, a = λ/cρ and a* = λ*/c*ρ*, λ and λ*, c and c* are thermal diffusivities, thermal conductivities and specific heat capacities of the ball and treated material in a ball mill, respectively. Thus, we have determined t–P–T conditions for the impact–friction contact of milling tools in a ball mills. During the collision of milling tools under the impact force f in a cylindrical volume πr2δ with the end face area s = πr2 (or a diameter of 2r) and a height δ (Fig. 2.5), a mean pressure impulse P = <σ> = f/s is realized, see Equation [2.8]. Because of this, when considering the interaction of the particles to be treated, selected from among those present in the given volume, mean values of these parameters are used. In this case, for a mean force characteristic of the interaction between any selected particles i and j in this volume (Fig. 2.3), fij = sij (f/s) = sij P is obtained. Independently, (see Equation [2.5a]) the expression for the impact contact area of the selected particles sij follows from the Hertz’s theory: sij = (π/4)[(3/2)Psij (θi + θi))/(Li + Lj)]2/3

[2.10]

Having determined the impact contact area of the selected particles sij from the Equation [2.10], fij, rij can be determined; the corresponding overall impact deformation (approach) of the particles εij can be found from Equation [2.6]. In order to determine separately the characteristic time (tij) and the normal constituent of the relative velocity (wn) describing the interaction of the selected particles i and j (Fig. 2.3), evident equation tij = 2εij/wn is used. An independent consequence of the Hertz’s theory, see Equation [2.2], gives the expression:

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2εij/wn = (π/4)(10π)0.4{wn−1(θi + θi)2(Li + Lj)[ρiρj/(ρiLj3 + ρjLi3)]2}0.2

[2.11]

Having determined wn, tij, the tangential constituent of the relative velocity of particle interaction wt = 2rij/tij can be determined. Applying the Hertz’s theory, see Equation [2.4], the corresponding normal (σn) and shift (σt) maximum stresses which are developed at the impact contact of the selected particles i and j are calculated. For the temperature impulse ΔT(x, t) over the averaged distance x and time t in the vicinity of the impact−friction contact of the particles i and j (Fig. 2.3), the equation of thermal conductivity is solved ∂T(x, t)/∂t = (aiaj)0.5∂2T(x, t)/∂x2 with the boundary conditions (t ≥ 0; T(x, 0) = T0; λij [∂T(∞, t)/∂x] + qij = 0; 0 ≤ x ≤ ∞; T(∞, t) = T0; ∂T(∞, t)/∂x = 0) for semi-limited bodies with the heat source of the density 2qij = ξijσnwt symmetric on the sliding plane (Lykov, 1965, 1967): ΔT(x, t) = 2qij(cicjλiλjρiρj)−0.25{t0.5iErfc[x/2(aiaj)0.25t0.5] − (t − tij)0.5iErfc[x/2(aiaj)0.25(t − tij)0.5]}

[2.12]

Here ξij is the dynamic friction coefficient, ai and aj, λi and λj, ci and cj are thermal diffusivities, thermal conductivities and specific heat capacities of the particles, respectively (Fig. 2.3). Numerical values of the calculated parameters are presented in Tables 2.1 and 2.2 as examples of the conditions under which the mechanochemical reaction CaCO3 + SiO2 = CaSiO3 + CO2 proceeds and for the conditions under which a single-component model material (NaCl) is activated. The results are presented for different variants of the interaction between the milling bodies. The heat capacity and heat conductivity of the materials are listed for the temperature T = (T0 + Tm)/2 where Tm is the melting point of the compounds to be treated and T0 = 375 K is the background temperature in the operating mill. Thus, we have determined t–P–T conditions for the impact–friction contact of treated particles 1 and 2 (Fig. 2.7) in a ball mill. For the case NaCl treatment, both for time t and for distance x (up to xm > ε1/2), temperature impulse exceeds the melting point of NaCl, ΔTm ≈ 700 K, see points ΔT(x, t) in Fig. 2.8: ΔTm = ΔT(0, tm) = ΔT(0, tm′) = ΔT(xm, t1). However, contact melting of particles requires special consideration since the time during which the temperature impulse exists at a level above melting point is relatively small.

2.4

Mechanism of mechanochemical processes

The quantitative aspect of this problem includes numerical estimates of heat and mass transfer in the local regions of the particles under given t–P–T conditions aimed at the description of the possibilities and mechanisms of different physicochemical processes, namely, phase transitions and chemical reactions. Local processes of melting and the subsequent inevi-

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Mechanism and kinetics of mechanochemical processes

23

x (cm) 10 1400

–4

10–5

10–6

0 1400

ε1/2

t1 = 7.6×10–9 s

1200

1200

xm

ΔT (0, t)

ΔT = ΔTm

800

800

ΔT (x, t 1)

1000

1000

600

600 tm

400

400

t′m

200

200

0

0 0

1

2

3 4 5 t (in t1 units)

6

7

2.8 Graph of the temperature impulse ΔT(x, t) at impact–friction contact of NaCl particles (R1 ≈ 1.63 × 10−4 cm, taking account of their contact melting) during treatment in a planetary mill EI 2 × 150 at ω = 10 s−1.

table solidification (crystallization or amorphization) of particles of substances treated in ball mills have consequences which are important for mechanochemistry generally, for example the effect of short-term melting is the formation of nanocrystals and amorphous nanoparticles. If we demonstrate the possibility of local contact melting of the treated particles, then the subsequent process of solidification will take place under very unusual conditions (Urakaev and Boldyrev, 1999b, 2000b). It has been shown by Landau (1950) that the propagation rate of the melting zone surface m′ is determined by a single parameter: h = π0.5c1ΔTm/2Hm

[2.13] −1

where Hm(NaCl) = 4.8 × 10 erg g is latent heat of fusion. The rate of fusion m′(t, h) increases with time and approaches to the stationary rate of fusion: 9

ms′ = q/ρ1(Hm + c1ΔTm)

[2.14]

where q is the density of the heat source at the boundary between the crystal and melt. Monograms of the ratio m′/ms′ = f(t, h) = f[y(t), h] have been presented as a function of a dimensionless parameter y(t) = (t/tm) − 1 for different

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h values, where tm is time required to achieve the melting point of the initial crystal. In the case under consideration, from Table 2.2 and Fig. 2.8 it is easy to obtain the required parameters tm ≈ 2.5 × 10−9 s and tm′ ≈ 1.0 × 10−8 s is the time of ceasing fusion as determined by the condition ΔT(0; tm, tm′) − ΔTm = 0; y = (tm′/tm) − 1 ≈ 3 and h ≈ 1.3. Let us define q for the fused zone existing at the impact-friction contact of NaCl particles treated in the mill. The solution of this problem has been presented in (Schlichting, 1951; Bowden and Persson, 1961). A laminar flow of viscous film of a liquid between two parallel plane walls moving at a relative rate of wt is considered; in our case from Table 2.2, wt = 5400 cm s−1, if the viscosity of the melt is μ(NaCl, T = Tm) and the distance between the walls is 2d*(t), the density of heat source 2q (in the case of symmetry along the middle plane) will be: q = μwt2/d*

[2.15]

Using Equations [2.14] and [2.15], a graphic solution m′/ms′ = f(t) and the identity d*(t) = m′(t − tm), the following equations were obtained: m′ = wt[μf(t)/(t − tm)ρ1(Hm + c1ΔTm)]0.5 d* = wt[μf(t)(t − tm)/ρ1(Hm + c1ΔTm)]0.5

[2.16]

q = wt[μρ1(Hm + c1ΔTm)/f(t)(t − tm)]0.5 Since f(t = tm′) = f(3, 1.3) ≈ 0.55, it is easy to determine from Equation [2.16] the numerical values of the required parameters for fusion at t = tm′ before the reverse process (i.e. crystallization) starts, in the vicinity of the impact–friction contact of NaCl particles during MA: m′(tm′) ≈ 33 cm s−1; 2d*(tm′) ≈ 7 × 10−7 cm; 2q(tm′) ≈ 5 × 1012 erg cm−2 s−1, when the size of NaCl particles in contact R1 ≈ 1.63 × 10−4 cm. It is important to compare 2q with 2q1 ≈ 9.1 × 1013 erg cm−2 s−1 for ‘dry friction’ between NaCl particles, see Table 2.2. When both walls are kept at a constant temperature T(x = 0) = T(x = 2d*) = Tm, the temperature of the melt ΔTm follows the square law (the distance from one wall to another being x and a maximum is achieved at the medium plane of shift, x = d*): ΔTm′ = T − Tm = μwt2x[1 − (x/2d*)]/2λd*

[2.17]

Since for NaCl melt, λ(Tm = 1074 K) = 2.5 × 105 erg cm−1 K−1 s−1, a maximum temperature increase at x = d* in the middle zone of the melt is ΔTm′(x = d*) = μwt2/4λ ≈ 0.6 K. This indicates that local heating is limited above the melting point of particles treated in ball mills when the possibility of fusion is taken into account (Fig. 2.8). Proof of possible contact melting during MA has been obtained by Bowden and Persson (1961) and by Urakaev and Avvakumov (1978) studying of the rate exchange reaction

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Mechanism and kinetics of mechanochemical processes (a)

25

(b)

0.8 0.8 α

α

0.6 0.4

0.4

0.2 0

0.6

0.2 0

4

8

12

16

τ / 60 = 8 s 0

τ / 60 (s)

2

4

6

8

10 12

ω (s–1)

2.9 Degree α of transformation for the reaction NaNO3 + KCl = KNO3 + NaCl depending on (Urakaev and Avvakumov, 1978): time τ (planetary mill EI 2 × 150 for ω = 7.1 s−1; ×, usual ball rotating mill) and frequency ω of rotation (for MA τ = 8 min in a planetary mill EI 2 × 150).

NaNO3 + KCl = KNO3 + NaCl executed in a ball planetary mill. It is noticeable (Fig. 2.9) that the reaction occurs at a high rate only when conditions for contact fusion at ω ≥ 5 s−1 have been achieved. The dependence of the thickness of crystallized layer d and overcooling ΔTc at the boundary between the phases on crystallization time tc are described by Equations [2.18] and [2.19] (Lyubov and Roytburd, 1962; Heertijes and Bouter, 1965): tc = ρ1Hm[d2 + (2λd/ρ1HmA)]/2λΔTs ≈ d/AΔTc

[2.18]

ΔTs/ΔTc = [1 + (2ρ1Hm A2ΔTstc/λ)]0.5 ≈ 1

[2.19]

where ΔTs = Tm − Ts (ΔTs is the substrate temperature). One can see that at the start of the crystallization process at small tc and d values, their interdependence is expressed in a linear form and ΔTs ≈ ΔTc. The value of the crystallization factor is A ≈ 0.1 cm K−1 s−1 for rapidly crystallized substances. Maximum overcooling is possible for the existing melt of NaCl at ΔTc* ≈ 170 K (Ubellohde, 1965). If this value is taken as ΔTc for crystallization in the region near the contacts between NaCl particles treated in a mill, the time for crystallization of a layer with a thickness of d = 2d* will be tc ≈ 4.1 × 10−8 s (see Equation [2.18]). This assumption can be shown to be correct both by the graphic representation shown in Fig. 2.8 or by Equation [2.12]. Even taking account of the released crystallization heat 2qc = 2d*ρ1Hm/tc ≈ 3.5 × 1011 erg cm−2 s−1 and 2qc < 2q < 2q1, overcooling of the melt by ΔTc = ΔTc* ≈ 170 K in the zone of contact between NaCl particles is achieved within the time interval Δt* ≈ 5 × 10−9 s which is much less than tc. Since crystallization and solidification are generally not identical processes (cf.

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Chalmers, 1964; Strickland-Constable, 1968; Flemings, 1974), the very high cooling rate obtained for the fused contact zone of the treated particles explains the widely known process involving amorphization during their MA of substances that are difficult to crystallize. However, for any compound that undergoes crystallization, it is important to know the size of critical nuclei rcr at which new crystals are formed in the fused zone: 2rcr = 2σsTm/ρ1HmΔTc, where 2rcr is the diameter of a sphere or a side of a cube of the new phase (a crystal that has just formed) and σs is the surface tension at the boundary between the melt and the crystal (for NaCl, σs = 23–34 erg cm−2; Gorskii and Mikulich, 1964). In order to meet the demand rcr < d*, it is necessary for ΔTc > 17 K. As we have shown above, this ΔTc can easily be achieved in the fused zone of impact– friction contact between NaCl particles. It is clear that the size of new formed crystalline nuclei of NaCl on the surface of treated particles will remain in the order rcr ≈ 10−7 cm. So, we have shown that during impact– friction contact of NaCl particles treated in comminuting devices, the formation of nanocrystal particles of about 10−7 cm occurs owing to local short-time fusion-crystallization processes (∼10−7 s). It is known that the properties of nanocrystals can be substantially different from the properties of macrocrystals. Because of this, MA of the substances in comminuting devices can be considered to be the result of the formation of a nanocrystal or X-ray amorphous state of the particles during treatment at the submicrometre level. More generally, this is the result of quenching non-equilibrium defects and structures (dislocations, twinning and intergrain boundaries, nanocrystal or nanoamorphous submicrometre states) during crystallization (solidification) of the melt during rapid cooling (to be more exact, during the short pressure and temperature impulses). Contact fusion during impact–friction contact of particles of different types leads to a sharp increase in diffusion coefficients. As a consequence, the crystallization of the melted layer should also result in the formation of the nanocrystal products of different mechanochemical reactions, in accordance with thermodynamic and kinetic conditions in the fused zone. The proposed mechanism for the mechanochemical reaction can be used directly in other types of MP involving contact melting (i.e. mechanodestruction of individual compounds except those decomposing in the solid state below their melting temperature like Ag2C2O4 and many other compounds) (Brown et al., 1980). It can be used for the majority of reactions that occur during mechanical alloying of metals (Lü and Lai, 1998; Bhadeshia, 2000; El-Eskandarany, 2001) except for ones that fuse with difficulty. For example, the proposed mechanism provides an explanation for the very low decomposition rate of sodium nitrate and many other compounds in the mill because it is identical to the thermal decomposition at

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Mechanism and kinetics of mechanochemical processes

27

their melting point (temperature impulse at contact between the particles is limited by the contact temperature at which the particles melt, as shown earlier). Thermal decomposition of sodium nitrate occurs at a noticeable rate only during fusion at temperatures much above its melting point Tm = 580 K. The kinetics of this type of process that occurs in comminuting devices, as well as the kinetics of processes with different mechanisms from the one described (solid state reactions including mechanical alloying of metals) will be considered in detail in the next section.

2.5

Kinetics of mechanochemical processes

The interaction of the particles of a compound that is being treated occurs in a coupled manner in the lining layer. Because of this, in order to build up a calculative kinetic model of ball milling, we shall select arbitrarily a pair of particles in contact, designated as 1 and 2 in Fig. 2.7. In order to obtain thermodynamic and kinetic parameters describing MP, it is sufficient to obtain the dependence of pressure σ and temperature T in some region x near the contact surface s12 of the particles 1 and 2 at the time t of impact– friction contact between these particles. In the general case for ball mills, the degree α of MP is a function of the frequencies ωk, the number of milling balls N, the ratio R/Dm of the ball size R to the linear dimension of the mill vial Dm, the properties of ball material and the compounds to be treated X and the time τ of MA; α(ωk, R/Dm, N, X, τ). Since any definite process in the mill is performed at some definite ωk, N, R/Dm, X values which are independent of τ, it is possible to separate the variables: α(ωk, R/Dm, N, X, τ) = K(ωk, R/Dm, N, X)α(τ) = K α(τ)

[2.20]

where K is the rate constant of the MP and α(τ) is described by the character of changes in the morphology of particles treated as a result of MP. The periodicity of action of loaded balls on particles of the substance being treated is one of the major characteristics of a ball mill of any type. In the general case, the impact of balls in the mill will be performed with a period of 1/ω′ or frequency ω′ = ω(ωk). Let us introduce a cycle of MA in ball mills: one period of impact of a ball load on the particles being treated is called the cycle of MA. The lined layers of the material being treated, which are formed at the mill vial walls and at the ball surface (Fig. 2.4a and Fig. 2.5), are composed of the contacting pairs of particles 1 and 2 (Fig. 2.3 and Fig. 2.7). Because of this, in order to describe the kinetics of the MP in general, it is sufficient to follow the changes in α that occur during the whole mill operation time τ for one pair of particles selected arbitrarily (Fig. 2.7). As a result of the impact interaction of particles 1 and 2 in a volume V*(σ, ΔT, τ) = d*(σ, ΔT)s12(τ) adjacent to the impact contact surface s12, conditions providing the possibility of MP arise. Let Φ*(U, σ, ΔT) describe the fraction

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of reacted material in volume V* as a result of a single interaction (here U is the activation energy of the process). Let us assume that for only one ball present in the mill vial (Ν = 1), at some moment in time τ = τ′ in its operation, a selected pair of particles 1 and 2 has moved into the region of the impact of the ball πr2δ (Fig. 2.5). In this case, the degree to which MP runs will be α*(N = 1) = g(k1, k2)Φ*V*(τ′)/ V(τ′), where g(k1, k2) takes account of the identical impact–friction interactions of the selected particles 1 and 2, not only with each other, but also with the surrounding neighbours, see the regions marked in black on the circles in Fig. 2.7. Now, let us assume that the number of balls N > 1 and the particles 1 and 2 have moved into the region of the ball impact during a cycle. In this case α*(N > 1) = ζ(N)α*(N = 1), where ζ(N) is a dimensionless coefficient depicting the collective action of the ball load, ζ(N = 1) = 1 and ζ(N > 1) ∼ N (Urakaev and Boldyrev, 1999c, 2000a). Consequently, for N > 1, the α value of the MP within a single cycle of ball load action on the lined layers of particles being treated can be written as follows: α*(τ = τ′) = ζ(N)g(k1, k2)Φ*V*(τ = τ′)/V(τ = τ′)

[2.21]

Here V(τ = τ′) = V1 = V2 is the overall volume of the selected pair of particles 1 and 2 (Fig. 2.7) at some moment in time τ = τ′ of mill operation. As far as the function g(k1, k2) is concerned, it is not difficult to state that g(k1, k2) = k1 + k2 ≈ 24 for the treatment of an individual compound and g(k1, k2) = 2k1k2/(k1 + k2) ≈ 12 for the treatment of a mixture of two compounds (in further numerical estimations we will consider g = 10). The overall fictitious degree αf at any moment in time τ during the MA will be equal to αf(τ) = α*Ψτ, since these particles are repeatedly involved in collisions with balls with a definite probability, Ψ = η(N, R/Dm)ψω′. The geometric probability ψ is the ratio of the impact contact area (s) of an unlined ball or a lined ball with the lined near-wall layer or with another lined ball to the ‘dead’ surface of the lined layer covered by the colliding ball and therefore not approachable by other balls. For example, in the case of collision of unlined balls with a flat near-wall layer, it follows from Equation [2.5a] and Fig. 2.5 that: ψ = s/(2R)2 = 2−4(10π)0.4ρ0.4(θ + θ)0.4Wn0.8

[2.22]

where R, W, ρ and θ are the radius, velocity, density and compliance of the ball and θ is the compliance of the lined layer of material being treated. Now we shall define the η(N, R/Dm) function which is responsible for the influence of the number of balls on the probability Ψ that repeated impact interactions of the particles 1 and 2 will occur. One can note that for the case of a single ball (N = 1), the probability Ψ = ω′(s/Πv) = [π(2R)2/Πv]ψω′, so the minimum value is η(N, R/Dm) = η(1, R/Dm) = π(2R)2/Πv ∼ R2/D2, where Πv ≈ 4πD2 is the overall area of lined walls accessible to the impact

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Mechanism and kinetics of mechanochemical processes

29

of a ball. Ideally, the maximum value η = 1 is achieved for a minimum number of balls Nmin ≈ D2/R2 providing a uniform treatment of the whole area Πv within unit time ω′. So, it has been shown that η(N, R/Dm) ≈ NR2/ D2 ∼ NR2. The product ζη ≈ (NR/D)2 ∼ (NR)2 can be estimated to be ∼1 with an accuracy of an order of magnitude, that is, ζη ≈ 0.1–10 (Urakaev and Boldyrev, 2000a). So, we have replaced the current time (τ) which has no physical sense by the real time (rpc) τrpc = τpcΨτ of the physicochemical (pc) interaction of the selected particles 1 and 2 with each other during MA. Here Ψ is the frequency of pulse interaction of the particles 1 and 2 with each other or 1/Ψ is the period at which the pulses follow each other and τpc is the time of the MP under consideration (Fig. 2.10), both during the impact–friction contact t12 of the particles and after the contact, under the conditions of σ and ΔT relaxation at the contacts. To sum up, for the transformation degree αf(τ), according to Equations [2.20]–[2.22], and at g(k1, k2)ζη ≈ 10: αf(τ) = α*Ψτ ≈ g(k1, k2)ζηψω′Φ*(V*/V)τ = K′τ

[2.23]

Equation [2.23] contains the value V*/V which is dependent on τ. The existence of this dependence is based on the completeness of the mechanochemical reaction which means the gradual disappearance of the initial particles (decrease of their size down to zero) during the formation of the final products of the reaction. Because of this, the assumption concerning the existence of some limiting quasi-equilibrium size of the treated particles R1 and R2 (see Section 2.3) can be true only for MA of the substance that

Δt(2)+Δt(1) Δt(1) = t′m – tm trpc (s)

= ~8×10–9 s

1/Ψ = ~ 20 s ΔT2

ΔT1 τ = τ0 τ (s)

2.10 Illustration of the dependence of real time (fixed) (trpc) of the activation process during contact between the particles 1 and 2 at the current time (τ) of the treatment in the mill vial, for melting NaCl particles as an example (τ0 is the time within which the first temperature impulse arises; 1/Ψ for ω = 10 s−1 is the period of the following ΔTi pulses).

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involves changes in properties without changing the composition of the compounds. So, Equation [2.23] with its linear law for the degree of MA αf(τ) which changes with the time of mechanical treatment τ can be used to describe the kinetics of MA (K′, s−1). In order to determine the time dependence V*(τ)/V(τ), the existence of some function describing the changes of the specific surface S(τ) of the treated particles during the MP is assumed. Then, for spherical particles 1 and 2, assuming (L1 + L2)/2 = L(τ) = 1/R12, we find the mean curvature of the particles at the contact point L(τ) = ρ12S(τ)/3. On the other hand V(τ) = 8π/3L3(τ), and for V*(σ, ΔT, τ) = d*(σ, ΔT) s12(τ), since σ and ΔT are independent of L(τ), using Equation [2.7] for s12: V*(τ) = 2−4π(10π)0.4ρ0.4Wn0.8(θ + θ)−1.6(θ1 + θ2)2d*L−2(τ) = C1d*L−2(τ) [2.24] So, the equation which we wanted to obtain will be: V*(τ)/V(τ) = 8πC1d*L(τ)/3 = 8πC1d*S(τ)ρ12 = C1C2d*τ = Cd*τ

[2.25]

where C1 and C2 = C2′ρ12/8π are constants, and θ1, θ2 and R12 = 2R1R2/(R1 + R2), ρ12 = 2ρ1ρ2/(ρ1 + ρ2) are, respectively, compliances and the mean radius and density of the particles 1 and 2 during treatment. Usually (Fig. 2.11 and Fig. 2.12) a linear law for the changes in the specific surface area S of the reagents is used in practical applications, so it is accepted in Equation [2.25] that S(τ) = C2′τ. In order to take account of the influence of the final solid products on the kinetics of MP, we employ a universal principle (Barret, 1973, p.152) which says that the reaction rate is equal to the fictitious specific rate multiplied by the size of the effective uniform region within which the reaction is localized. This leads to the following correlation between the real and fictitious transformation degree: dα(τ) = (1 − α)pdαf(τ). Integrating this equation for different values of p (for explanation, see Delmon, 1969: 281, 414) and using Equations [2.23] and [2.25] makes Equations [2.20] and [2.21] more specific and leads to the following relations: Kτ2 = α

at p = 0

[2.26]

Kτ = −ln (1 − α)

at p = 1

[2.27]

Kτ2 = α/(1 − α)

at p = 2

[2.28]

2

Here Equation [2.26] should correspond to the kinetics of mechanochemical reactions without the formation of solid final products, for example: NH4NO3 = N2O + 2H2O, C + 2S = CS2, Ni + 6CO = Ni(CO)6 and possibly the kinetics of the MA (K, s−2). The case in Equation [2.27] is realized when the degree to which the new solid products formed during © Woodhead Publishing Limited, 2010

Mechanism and kinetics of mechanochemical processes

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τ2 × 10–6 (s2) 1.0

12 14

0 1 2 3 4 5 6 7 8 9 10

16

18

20

22

24

–lg (1–α)

0.1 0.2

α/(1–α)

0.8

0.3 0.5

0.6

0.6 0.7 0.8

0.4

0.9

WCo8 ρ = 15.6 g cm–3 R = 0.2 cm N = 383 ω = 9 s–1

5 0.2

0.0 S0 = 0.035 0

1

2

3

4

5

6 –3

τ × 10

7

8

–lg (1 – α); α/(1 – α) ([2.27]) ([2.28])

0.4

α

S (m2 g–1)

15

10

0

1.0 1.1 1.2 9

10

11

12

(s)

2.11 Change in specific surface (S, m2 g−1) and conformity to the Equations [2.27] and [2.28] of the rate of mechanochemical decomposition of silver oxalate, Ag2C2O4 = 2Ag + 2CO2, in a planetary mill KhK 871 (Urakaev et al., 1982).

S – S0 (m2 g–1)

40

WO3

30

WCo8 ρ = 15.6 g cm–3 R = 0.2 cm N = 228 ω = 14.3 s–1

20

BaCO3 10

0 0

4

8 τ / 60 (s)

12

2.12 Change in specific surface (S, m2 g−1) of barium carbonate (BaCO3) and tungsten oxide (WO3) in a planetary mill KhK 871 (Kopylov et al., 1978).

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the MA of a substance affect the kinetics is proportional to the transformation degree α of the reaction, while Equation [2.28] is realized for nonlinear dependence of the kinetics upon the solid products formed. For rate constant K accompanied by changes in particle size, see Equations [2.26]– [2.28], using Equations [2.22]–[2.25]: K = 2−11(10π)0.8g(k1, k2)ζηρ0.8(θ + θ)−1.2(θ1 + θ2)2ωWn1.6ρ12Cd*Φ*

[2.29]

Numerical estimates of rate constants (K′, K) for MA and mechanochemical reactions in ball mills according to the Equations [2.21]–[2.29] will be performed in the next section.

2.6

Use of kinetic equations

Let us transform Equation [2.29] into a form which can be used to estimate K numerically. One can see from Tables 2.1 and 2.2 and Fig. 2.8 that ΔT(x, t12) is relatively constant at the depth x of the treated particles up to a distance x = x* = ε12/2. For example, in the case of NaCl treatment, the temperature impulse ΔT(x, t1) at the impact–friction contact point of particles can exceed their melting point ΔTm up to a distance 2xm* > ε1/2. Consequently, at some approximation, one can accept d*(σ, ΔT) = 2x* ≈ ε12, that is, the thickness of the reaction zone is equal to the total impact deformation of the particles 1 and 2 under consideration. Then, using the result obtained for ε12 in Equation [2.6] and [2.6a], the rate constant from Equation [2.29] becomes: K = 2−14(10π)1.2g(k1, k2)ζηρ1.2(θ + θ)−2.8(θ1 + θ2)4ωWn2.4ρ12R12CΦ*

[2.30]

A dimensionless parameter Φ*(U, σ, ΔT) can be determined by the mechanism of the chemical reaction under the prevailing conditions provided that calculated values of σ and ΔT impulses in the region of impact– friction contact of particles 1 and 2 are known. According to the definition, Φ*[U, σ(x, t) ΔT(x, t)] = n*[σ(x, t),ΔT(x, t)]/n0*, where n* is the number of reacted molecules and n0* is the total number of molecules (atoms) in the volume under consideration V* in which the pressure σ(x, t) and temperature ΔT(x, t) impulses are realized. The rate of the chemical reaction leading to the final products is d2n*/dxdt = n0*Kr[U, σ(x, t), ΔT(x, t)], where Kr is the reaction rate constant. It seems impossible to integrate the equation in this form, so, as a first approximation, under the conditions at which the chemical reaction occurs at the point of contact of the particles 1 and 2, the mean impulses can be σ(x, t) = σ and ΔT = 0.5[ΔT(0, t12) + ΔT(ε12/2, t12)] within the time interval. This means that we accept τrpc = t12. In this case, it follows that: Φ* = t12Kr = t12Kr0exp[−(U − χ12<σ>)/k(T0 + ΔT)]

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[2.31]

Mechanism and kinetics of mechanochemical processes

33

where Kr0 is a pre-exponential factor, χ12 ∼ 10−23 cm3 (the structural factor, see Regel et al., 1974: 130) and k = 1.38 × 10−16 erg K−1 (the Boltzmann factor) are constants and T0 is the background temperature in the mill. Now we can present a more specific equation for K taking account of t12 from Equations [2.2] and [2.11]: K = 2−16(π/2)0.25(10π)2.6g(k1, k2)ζηρ1.1(θ + θ)−2.4 (θ1 + θ2)4ωWn2.2ρ121.5R122CKr

[2.32]

For Kr depending on the conditions of MA, only the most important parameters, that is, ρ and W of a ball, see Equation [2.1], at the next assumption about the known function Kr in Equation [2.31] are considered. The activation energy U of thermal solid phase processes controlled by diffusion and chemical transformations are usually much higher than χ12σ. If U >> χ12σ and ΔT >> T0, according to Equations [2.1], [2.12] and [2.32] it follows that: Wn = Aω, Kr ≈ Kr0exp(−U/kΔT) = Kr0exp(−Λ/ω0.9ρ0.45)

[2.33]

Λ ≈ 3(π/2) 2 (10π) iErfc[0](θ + θ) (c1c2λ1λ2ρ1ρ2ρ12) U/kA0.9R120.5ξ12(θ1 + θ2)

[2.34]

K ∼ (NR) Θ ρ ω exp(−Λ/ω ρ

[2.35]

1.625 1.5

2

1.6 1.1

0.2

3.2

1.8

0.9 0.45

)

0.25

where generalized compliance Θ1.6 = (θ + θ)−2.4(θ1 + θ2)4 depicts the dependence of K on the mechanical properties (compliances) of the material of both the balls, θ or θ = (θ1 + θ2)/2 for lined balls, and the particles being treated (θ1, θ2). Specific cases of application of the Equations [2.20]–[2.35] to various mechanochemical reactions (Figs 2.9, 2.11 and 2.13–2.17) occurring in ball mills are considered in the next sections.

2.6.1 Mechanical activation of substances Let us assume that the mechanism of MA of NaCl particles involves contact melting followed by hardening which is accompanied by the formation of non-equilibrium defects. The time interval within which ΔT ≥ ΔTm for NaCl particles is tpc = Δtm = t′m − tm ≈ 8 × 10−9 s (Figs 2.8 and 2.10). The possibility that NaCl particles will melt under these conditions and the determination of the thickness of fused layer d* ≈ 7 × 10−7 cm were presented above. Returning to the Equation [2.23] for the conditions under which NaCl particles are treated in the mill EI 2 × 150 at g(k1, k2) ≈ 10, ζη ≈ 1, R = 0.2 cm, ω = 10 s−1, Equation [2.22] or Table 2.2 are also used to find Ψ = Ψ = ψω ≈ 53 × 10−3 s−1 and 1/Ψ ≈ 20 s. This means that, in reality, NaCl particles with a size R1 ≈ 1.6 × 10−4 cm are subjected to the impact of balls about once every 20 s. The coefficient providing the transition to real time of treatment will be trpc/τ = tpcΨ = ΔtmΨ ≈ 4.2 × 10−10 (Fig. 2.10). To estimate

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High-energy ball milling ω = 15.3 s–1 ω = 12.4 s–1 ω = 11.0, 9.7 s–1 ω = 9.0, 8.8 s–1

1.00

ω = 8.2 s–1

0.80

ω = 7.1 s–1 WCo8 ρ = 15.6 g cm–3 R = 0.2 cm N = 228

α

0.60 0.40

ω = 5.0 s–1

0.20 0.00

0

20

40

60

80 τ / 60 (s)

100

120

140

2.13 Degree α of transformation for the reaction BaCO3 + WO3 = BaWO4 + CO2 in a planetary mill KhK 871 depending on frequency ω of rotation (Avvakumov and Urakaev, 1982; Avvakumov et al., 1983).

WCo8 ρ = 15.6 g cm–3

1.00

Steel ρ = 7.9 g cm–3

Al2O3 ρ = 4.0 g cm–3

0.80

α

0.60

0.40

0.20

0.00 0

5

10

15

20

25

30

35

40

τ / 60 (s)

2.14 Degree α of transformation for the reaction BaCO3 + WO3 = BaWO4 + CO2 in a planetary mill KhK 871 depending on the density ρ of balls; R = 0.2 cm, N = 228, ω = 12.4 s−1 (Avvakumov and Urakaev, 1982).

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Mechanism and kinetics of mechanochemical processes 1.00

35

R = 0.4 cm R = 0.25 cm

0.80

R = 0.15 cm

0.60 α

R = 0.11 cm R = 0.075 cm

0.40

0.20

0.00 0

5

10

15

20

25

30

τ / 60 (s)

2.15 Degree α of transformation for the reaction BaCO3 + WO3 = BaWO4 + CO2 in a mill KhK 871 depending on the radius R of balls; for steel ρ = 7.9 g cm−3, m = 209 g, ω = 12.4 s−1 (Avvakumov and Urakaev, 1982; Avvakumov et al., 1983). N = 467

1.00

N = 137

N = 262

0.80

N = 68

α

0.60

0.40

0.20

0.00 0

10

20

30

τ / 60 (s)

2.16 Degree α of transformation for the reaction BaCO3 + WO3 = BaWO4 + CO2 in a mill KhK 871 depending on the number N of balls; for steel ρ = 7.9 g cm−3, R = 0.2 cm, ω = 12.4 s−1 (Avvakumov and Urakaev, 1982).

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High-energy ball milling WCo8 ρ = 15.6 g cm–3; R = 0.2 cm N = 228 ω = 12.4 s–1

1.00

BaCO3 β = 2.03 Mbar–1

0.80

SrCO3 β = 1.75 Mbar–1 MgCO3 β = 1.07 Mbar–1

CaCO3 β = 1.37 Mbar–1

α

0.60

0.40

0.20

0.00 0

1

2

3

4

5

6

7

8

9

10

τ / 60 (s)

2.17 Degree α of transformation for the reactions MeCO3 + WO3 = MeWO4 + CO2, where Me = Ba, Sr, Ca and Mg, in a mill KhK 871 depending on the compressibility β (Clark, 1966; Avvakumov and Urakaev, 1982) of the alkaline earth’s carbonates [β = 1/Ω (modulus of volume elasticity) ∼ 1/E (Young’s modulus) ∼ θ (compliance) ∼ Θ (generalized compliance)]. In τ (τ, s) 9

8

7

6

5

4

3

2

1

3

–6 –5

2

–4

1b

–3

3b

0

–2

2b

–1

–1

–2

2a

3a

0

1a

–3

1

–4

2

–5

3

–6

0

1

2

3

4 5 6 In τ (τ, s)

In [α /(1-α)]

1 In [–In (1-α)]

0

7

8

9

2.18 Experimental kinetic curves processed at the logarithmic coordinates of Equations [2.27] and [2.28]. The straight lines (a) correspond to Equation [2.27] and lines (b) to Equation [2.28]; 1 is related to the reaction NaNO3 + KCl = KNO3 + NaCl in a planetary mill EI 2 × 150 at ω = 10 s−1; 2 and 3 are related to the reactions BaCO3 + WO3 = BaWO4 + CO2 and Ag2C2O4 = 2Ag + CO2 in a planetary mill KhK 871 at ω = 11 s−1 and ω = 9 s−1, respectively. © Woodhead Publishing Limited, 2010

Mechanism and kinetics of mechanochemical processes

37

K′ in Equation [2.23], Φ* ≡ 1 as an identity should be taken into account since the accepted mechanism of NaCl particle activation is based on a process involving fusion and quenching. Let us determine (using Table 2.2) the ratio V*/V = 3d*s1/8πR13 ≈ 2.5 × 10−5; so the rate constant will be K′ ≈ 1.3 × 10−5 s−1. Because of this, the treatment time for complete fusion and quenching of all NaCl particles in a mill (αf = 1) will be, according to Equation [2.23], τ = 1/K′ ≈ 7.5 × 104 s or ∼21 h. Now we shall assume that for the same mechanism of NaCl activation (Φ* ≡ 1) the size of particles being treated changes. In this case, the MA kinetics will be described by Equation [2.26] with the constant from Equation [2.29]. According to Equation [2.29], K(ω = 10 s−1) ≈ 1.5 × 10−7 s−2 will be obtained. According to Equation [2.26], the corresponding time of MA (τ) necessary for NaCl particles to remelt completely (α = 1) will be τ = (1/K)0.5 ≈ 2600 s or only ∼0.7 h. One can see that a substantial difference is observed at the poles.

2.6.2 Reactions controlled by diffusion Now let us consider the kinetics of the reactions in a system of ionic salts (Fig. 2.9). It has been shown above that the rate of this exchange reaction is limited by mutual diffusion of ions in the fused layer. So, in Equation [2.29] we have d*Φ* ≈ 2.4 × 10−6 cm. Using relation [2.29], it is easy to obtain the calculated value of the rate constant for the reaction under consideration from Equations [2.26]–[2.28]: K(ω = 10 s−1) ≈ 2.2 × 10−6 s−2. Figure 2.18 presents the results of the treatment of kinetic curves at the coordinates determined by Equations [2.27] and [2.28] for a series of MA reactions. One can see that the experimental points agree well with the straight lines at the coordinates of Equation [2.27] within the whole range of α change. It can be concluded that the degree of ‘screening’ of a MP by the solid products of the reaction is proportional to the volume of the transformed substance. For the reaction being considered, an experimental value of the rate constant at the coordinates of Equation [2.27] is K1a ≈ 8.3 × 10−5 s−2 which is about 40 times higher than the calculated value. This difference will be discussed below when considering the dependence of K on the conditions of MA. An example of MP controlled by diffusion is a reaction BaCO3 + WO3 = BaWO4 + CO2. Kinetic curves α(τ) for this reaction were obtained by measuring the pressure of gaseous CO2 formed (see Figs 2.13–2.17). The thermal mechanism of this reaction involves the diffusion of tungsten into decomposing carbonate, the reaction rate being determined by the diffusion coefficient of tungsten. In order to estimate the rate constant of this MP, we shall use Equation [2.30] and the parameters calculated for a related reaction, see Table 2.1. It is shown that the key value Φ* ≈ 1. The constant for

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High-energy ball milling

the MA reaction under consideration will be estimated using Equation [2.30]: K(ω = 11 s−1) ≈ 0.6 × 10−6 s−2. The rate constant determined experimentally at the coordinates of Equation [2.27], see Fig. 2.18, is K2a ≈ 2.1 × 10−5 s−2 which exceeds the calculated value by a factor of 35. It is easy to note (Fig. 2.13) that when ω of mill rotation is decreased, owing to the decrease in the temperature pulse at the point of contact of particles, synthesis of BaWO4 should be transferred into the region controlled by the kinetics (see below).

2.6.3 Reactions controlled by the kinetics Decomposition of silver oxalate can be used as an example of MP controlled by kinetics (Fig. 2.11). Kinetic curves α(τ) for thermal and MA decomposition according to the scheme Ag2C2O4 = 2Ag + CO2 were also obtained by measuring the pressure of gaseous CO2 formed. The rate constants Kr and K of silver oxalate decomposition, under the pressure conditions and the temperature impulses that are realized at the point of impact–friction contact of the particles being treated can be estimated using Equations [2.31] and [2.32]. Using the calculation procedure described above when deducing Equation [2.31], we obtain Kr ≈ 2.0 × 10−7 s−1. This gives, see Equation [2.32], K(ω = 9 s−1) ≈ 4.0 × 10−9 s−2. The experimental value is K2a ≈ 3.5 × 10−8 s−2 (see Figs 2.11 and 2.18) which is about 10 times higher than the calculated value.

2.6.4 Effect of mechanochemical activation conditions on the kinetics In order to establish the dependence of K on the conditions of MA, see Equation [2.35], a complex investigation of the kinetics of MA reactions of the type BaCO3 + WO3 = BaWO4 + CO2 has been performed (Figs 2.13–2.17). Experimentally determined rate constants of the studied reactions, depending on the treatment conditions, should be straight lines at the coordinates shown in Fig. 2.19. It shows a result of this treatment which is in satisfactory agreement with the above considerations and assumptions. Using the slopes of the straight lines (1) and (2) in Fig. 2.19, <[ΔK(N) + ΔK(R)]/ΔN2 ΔR2> ≈ 3.5 × 10−12 s cm−2, one can determine the value ζη ≈ 3.5 × 10−12 N2R2/K(N = 1; R) ≈ 15 where N = 137 and R = 0.2 cm can be determined, corresponding to the identical conditions of the experiment when studying the dependence of K on the number and size of balls where K(N = 1; R = 0.2 cm) ≈ 1.8 × 10−10 s−2. Using the slopes of the straight lines (3) and (5) we obtain the values lg K(ω = 1; ρ = 15.6) ≈ −17.8, lg K(ρ = 1; ω = 12.4) ≈ −6.65 and in Equations

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Mechanism and kinetics of mechanochemical processes

39

(5) ← 3.2 lg ω – lg e/ω0.9 3.5

3.0

2.5

2.0

10 –3

(4)

6

–5

(3)

4

(1)

2 0

–4 lg K

K × 105 (s–2)

(2)

(5)

8

–6 –7

(3)

2

4

6

8

10 12 14

Θ1.6 × 1015

(4) 0.2 0.4 0.6 0.8 1.0 1.2 1.4

1.1 lg ρ – lg e/ρ0.45

(2)

N 2 × 10–4

(1)

4 2

8 4

6

12 8

16

20

10 12 14

R2 × 100 (cm2)

2.19 Experimental values of the rate constants K (s−2) for BaWO4 mechanochemical synthesis compared to the theoretical dependence [2.35] of K on the conditions of mechanical treatment in a mill KhK 871 (see also text): (1, 䊉) K − R 2, R in cm; (2, 䊊) K − N 2; (3, +) K − Θ1.6, Θ in cm2 dyn−1; (4, ⵧ) lg K − [(1.1 lg ρ − (lg e / ρ0.45)], ρ in g cm−3; (5, 䊏) lg K − [(3.2 lg ω − (lg e / ω0.9)], ω in s−1.

[2.32] and [2.33], Λ ≈ ω0.9ρ0.45[lg K(ω = 1) − lg K(ρ = 1) + 3.2 lg ω − 1.1 lg ρ]/ (ρ0.45 − ω0.9)lg e ≈ 110 g0.45 cm−1.35 s−0.9, where ρ = 15.6 g cm−3 and ω = 12.4 s−1 correspond to identical K when studying the dependence of K on the ball density and on the mill rotation frequency. It is interesting to compare the experimental value of Λ with the theoretical one (Λ = 84) for a related MA synthesis of calcium silicate (for parameters, see Table 2.1): Λ ≈ 103 g0.45 cm−1.35 s−0.9. In the case under consideration, the two values agree to an accuracy within an order of magnitude. It is important to note that for low ω, the diffusion regime of MA synthesis changes to the kinetic regime, see the fracture in curve (5) in Fig. 2.19. The final results of the comparison between the experimental and calculated rate constants of the MA reactions in planetary mills, taking account of the value ζη ≈ 15 determined experimentally, will be listed below. For the reaction NaNO3 + KCl = KNO3 + NaCl, K1a = 8.3 × 10−5 s−2 and the calculated value, K ≈ 1.5 × 10−5 s−2; for the synthesis of barium tungstate, K2a = 21 × 10−6 s−2 and the calculated value, K ≈ 9.0 × 10−6 s−2; for the mechanical destruction of silver oxalate, K3a = 3.5 × 10−8 s−2 and the calculated value, K ≈ 6 × 10−8 s−2. So one can see that the theory providing the calculation of

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High-energy ball milling

the kinetics of MA reactions is in quantitative agreement with the kinetic parameters determined experimentally, within an order of magnitude.

2.7

Conclusions

Using different ball mills as examples, it has been shown that, on the basis of the theory of glancing collision of rigid bodies, the theoretical calculation of t–P–T conditions and the kinetics of mechanochemical processes are possible for the reactors that are intended to perform different physicochemical processes during mechanical treatment of solids. According to the calculations, the ‘physicochemical’ effect of mechanochemical reactors is due to short-time impulses of pressure (ΔP = σ ∼ 1010–1011 dyn cm−2) with shift, and temperature ΔT(x, t). The highest temperature impulse ΔT ∼ 103 K are caused by the ‘dry friction’ phenomenon. Typical spatial and time parameters of the impact–friction interaction of the particles with a size R ∼ 10−4 cm are as follows: localization region, Δx ∼ 10−6 cm; time, Δt ∼ 10−8 s. On the basis of the obtained theoretical results, the effect of short-time contact fusion of particles treated in various comminuting devices can play a key role in the mechanism of activation and chemical reactions for wide range of mechanochemical processes. This role involves several aspects, that is, the very fact of contact fusion transforms the solid phase process onto another qualitative level, judging from the mass transfer coefficients. The spatial and time characteristics of the fused zone are such that quenching of non-equilibrium defects and intermediate products of chemical reactions occurs; solidification of the fused zone near the contact point results in the formation of a ‘nanocrystal or nanoamorphous state’. The calculation models considered above and the kinetic equations obtained using them allow quantitative ab initio estimates of rate constants to be performed for any specific processes of mechanical activation and chemical transformation of the substances in ball mills.

2.8

References

anderson ol (1965), ‘Determination and some applications of the isotropic elastic constant polycrystalline systems received from the data for monocrystals’, in Physical Acoustics, Principles and Methods, vol. III, part B: Lattice Dynamics, Mason WP (ed.), Academic Press, New York and London, 56–119. avvakumov eg and urakaev fkh (1982), ‘Kinetics of solid-phase mechanochemical reactions depending on conditions of ball milling’, in Kinetics and Mechanism of Solid-Phase Reactions, Kemerovo, KemGU, 3–9 (in Russian). avvakumov eg, urakaev fkh and tatarintseva mi (1983), ‘About two modes of course of mechanochemical reactions depending on conditions of ball milling’, Kinetics and Catalysis, 24, 227–9 (in Russian). barret p (1973), Cinétique Hétérogène, Authier-Villars, Paris.

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batuev gs, golubkov yuv, efremov ak and fedosov aa (1969), Engineering Methods of Impact Processes Investigation, Mashinostroeniye, Moscow (in Russian). belyaev eyu, lomovsky oi, ancharov ai and tolochko bp (1998), ‘Investigation of the reaction zone structure under mechanochemical synthesis of metal disilicides by a method of local diffractometry’, Nuclear Instr Methods Phys Res A, 405, 435–9. bhadeshia hkdh (2000), ‘Mechanically alloyed metals’, Proc Roy Microscopical Soc, 35, 95–102. bowden fp and persson pa (1961), ‘Deformation, heating and melting of solids in high-speed friction’, Proc Roy Soc Lond A, 260, 433–58. bowden fp and tabor d (1974), Friction: An Introduction to Tribology, Science Study Series No. 41, Heinemann Educational, Oxford, 178 pp. bowden fp and tabor d (1986), The Friction and Lubrication of Solids, International Series of Monographs on Physics, 2nd edition, Oxford University Press, USA, 372 pp. brach rm (1989), ‘Rigid body collisions’, ASME J Appl Mech, 56, 133–8. brennan jn (1957–58), Bibliography on Shock and Shock Excited Vibrations, in two volumes, Pennsylvania State University, College of Engineering and Architecture, Engineering Research Bulletin, nos. 68, 69. See: http://openlibrary.org/b/ OL21498394M/Bibliography_on_shock_and_shock_excited_vibrations._ Volume_II. brown me, dollimore d and galwey ak (1980), Reactions in the Solid State, Elsevier Science, Amsterdam. butyagin pyu (2003), ‘Diffusion and deformation models of mechanochemical synthesis’, Colloid J, 65, 648–51. butyagin pyu (2006), ‘From spontaneous dispersion to mechanical alloying’, Colloid J, 68, 397–403. butyagin pyu and streletskii an (2005), ‘The kinetics and energy balance of mechanochemical transformations’, Phys Solid State, 47, 856–62. chalmers b (1964), Principles of Solidification, Wiley, New York. chattopadhyay pp, manna i, talapatra s and pabi sk (2001), ‘A mathematical analysis of milling mechanics in a planetary ball mill’, Mater Chem Phys, 68, 85–94. clark spjr (1966), Handbook of Physical Constants, Geological Society of America, Connecticut. coriolis g (1835), Théorie Mathématique des Effets de Jeu de Billard, Paris. delassus e and peres j (1923–24), ‘Note sur le choc en tenant compte du frottement de glissement’, Nouv Ann Math, 2, 383–91. delmon b (1969), Introduction a la Cinétique Hétérogène, Éditions Technip, Paris. delogu f, monagheddu m, mulas g, schiffini l and cocco g (2000), ‘Impact characteristics and mechanical alloying processes by ball milling: Experimental evaluation and modeling outcomes’, Int J Non Equilibrium Processing, 11, 235–69. dinnik an (1952), ‘Impact and stress of solids’, in Selected Works (Vol. 1), AN USSR, Kiev (in Russian). el-eskandarany ms (2001), Mechanical Alloying for Fabrication of Advanced Engineering Materials, William Andrew Publishing, Noyes, 260 pp. flemings mc (1974), Solidification Processing, Mc Graw-Hill Book Co, New York. goldsmith w (1960), IMPACT. The Theory and Physical Behaviour of Colliding Solids, Edward Arnold, London.

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gorskii fk and mikulich as (1964), ‘Size of interphase surface energy of sodium chloride on a border crystal-melt’, Crystallization Mechamism and Kinetics, Syrota NN (ed.), Nauka i tekhnika, Minsk, 71–8 (in Russian). harris cm and piersol ag (eds) (1961), Harris’ Shock and Vibration Handbook, McGraw-Hill, Toronto, in three volumes. heertijes pm and bouter ja (1965), ‘Subcooling effects in a crystallisation system’, Chem Proc Eng, 46, 654–8. hertz h (1895), Gesammelte Werke, J.A. Barth (Arthur Meiner), Leipzig, Vol. 1. horák z (1948), ‘Impact of a rough ball spinning round its vertical axis onto a horizontal plane’, Trans Dept Tech Univ Prague (Sb. CVUT Vys. Sk. Stroj. Inz., Prague), 1–16. horák z and pacáková i (1961), ‘The theory of the spinning impact of imperfectly elastic bodies’, Czechoslovak J Phys, 11, 46–65. ismail ka and stronge wj (2008), ‘Impact of viscoplastic bodies: dissipation and restitution’, J Appl Mech, 75, 061011 (5 pages). johnson kl (1985), Contact Mechanics, Cambridge University Press, Cambridge, UK. kajdas ck (2005), ‘Importance of the triboemission process for tribochemical reaction’, Tribol Int, 38, 337–53. keller jb (1986), ‘Impact with friction’, ASME J Appl Mech, 53, 1–4. kilchevsky na (1976), Dynamic Contact Stress of Solids. Impact, Naukova dumka, Kiev (in Russian). kobayashi k (1995), ‘Formation of coating film on milling balls for mechanical alloying’, JIM (Materials Trans), 36, 134–7. koch cc (2003), ‘Top-down synthesis of nanostructured materials: Mechanical and thermal processing methods’, Rev Adv Mater Sci, 5, 91–9. kopylov av, avvakumov eg and urakaev fkh (1978), ‘Mechanochemical reactions of barium carbonate with oxides of metals of IV, V, and VI group’, Izv Sib otd AN SSSR: Ser khim nauk, 9/4, 8–14 (in Russian). korchagin ma and lyakhov nz (2008), ‘Self-propagating high-temperature synthesis in mechanoactivated compositions’, Russian J Phys Chem B, 2, 77–82. kubenko vd (1999), ‘Local wave theory of the collision of elastic bodies. Noncentral impact – two-dimensional problem in the ideal-fluid approximation’, Int Appl Mech, 35, 1155–66. landau hg (1950), ‘Heat conduction in a melting solid’, Quart Appl Math, 8, 81–94. love aeh (1927), A Treatise on the Mathematical Theory of Elasticity, 4th edition, University Press, Cambridge. lü l and lai mo (1998), Mechanical Alloying, Springer, Berlin, 276 pp. lykov av (1965), Theory of Energy and Mass Transfer, Prentice Hall, New York, 392 pp. lykov av (1967), Teoriya teploprovodnosti, Moskva, Vysshaya shkola, 600 pp (in Russian). lyubov bya and roytburd al (1962), ‘About influence of overcooling on border of the unit of phases for speed of moving of front of crystallization in conditions directed heat transfer’, in Crystallization and Phase Transitions, Syrota NN (ed.), Minsk, AN BSSR, 226–34 (in Russian). maw n, barber jr and fawcet jn (1976), ‘The oblique impact of elastic spheres’, Wear, 38, 101–14.

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miani f and maurigh f (2004), ‘Mechanosynthesis of nanophase powders’, in Encyclopedia of Nanoscience and Nanotechnology, Marcel Dekker. mio h, kano j, saito f and kaneko k (2002), ‘Effects of rotational direction and rotation-to-revolution speed ratio in planetary ball milling’, Mater Sci Eng A, 332, 75–80. painlevé p (1895), Lec¸ons sur le Frottement, Paris. panovko yag (1977), Introduction to the Theory of Mechanical Impact, Nauka, Moscow (in Russian). panovko yag (1985), Mechanics of Deformable Solids. Modern Concepts, Mistakes and Paradoxes, Nauka, Moscow (in Russian). peres j (1923–24), ‘Choc en tenant compte de frottement’, Nouv Ann de Math, 2, 98–107. peres j (1923–24), ‘Choc de deux solides avec frottement’, Nouv Ann de Math, 2, 216–31. persson bnj (2000), Sliding Friction: Physical Principles and Applications, New York, Springer Verlag, New York. poisson sd (1817), Mechanics, Longman, London, in two volumes. popov vl, psakhie sg, shilko ev, dmitriev ai, knothe k, bucher f and ertz ì (2002), ‘Friction coefficient in “rail–wheel” contacts as a function of material and loading parameters’, Physical Mesomechanics, 5(3), 17–24. pustov lyu, kaloshkin sd, cherdyntsev vv, tomilin ia, shelekhov ev and salimon ai (2001), ‘Experimental measurement and theoretical computation of milling intensity and temperature for the purpose of mechanical alloying kinetics description’, Mater Sci Forum, 360–2, 373–8. regel vr, slutsker ai and tomashevsky ee (1974), Kinetic Nature of Solids Strength, Nauka, Moscow (in Russian). revesz a and takacs l (2007), ‘Coating metals by surface mechanical attrition treatment’, J Alloys Compounds, 441, 111–14. routh ej (1897), Dynamics of a System of Rigid Bodies, Macmillan and Co., London, vol. 1. schlichting h (1951), ‘Einige exakte Lösungen für die Temperaturverteilung in einer laminaren Strömung’, Zeitschrift für angewandte Mathematik und Mechanik, 31, 78–83. schmitz t, action j, ziegert j and sawyer w (2003), ‘Dynamic friction coefficient measurements: device and uncertainty analysis’, in Proceedings of the 18th ASPE Annual Meeting, October 26–31, Portland, OR (on CD). schneider u (1968), ‘Makroskopische und mikroskopische Eigenschaftsanderungen von Feststoffpulvern infolge starker mechanischer Beanspruchung in Mühlen’, Aufbereit Techn, 9, 567–73. stoimenov lg (1992), ‘Solution of the problem of oblique impact of bodies: Impact models for rough bodies’, Int Appl Mech, 28, 477–83. strickland-constable rf (1968), Kinetics and mechanism of crystallization from the fluid phase and of the condensation and evaporation of liquids, Academic Press, New York. stronge wj (1990), ‘Rigid body collisions with friction’, Proc Roy Soc London Ser A, 431, 169–81. suryanarayana c (2001), ‘Mechanical alloying and milling’, Progr Mater Sci, 46, 1–184.

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takacs l (2002), ‘Self-sustaining reactions induced by ball milling’, Prog Mater Sci, 47, 355–414. takacs l and torosyan ar (2007), ‘Surface mechanical alloying of an aluminum plate’, J Alloys and Compounds, 434, 686–8. timoshenko s (1955), Vibration Problems in Engineering, D. Van Nostrand, Toronto, 3rd edition in collaboration with Young DH. ubellohde ar (1965), Melting and Crystal Structure, Clarendon Press, Oxford. urakaev fkh (2004), ‘Theoretical bases of mechanochemical processes in ball milling’, Eurasian Chemico-Tech J, 6, 239–54. urakaev fkh (2007), ‘Mechanodestruction of minerals at the crack tip (overview): 1. Experiment’, Phys Chem Minerals, 34, 351–61. urakaev fkh (2008), ‘Mechanodestruction of minerals at the crack tip (overview): 1. Theory’, Phys Chem Minerals, 35, 231–9. urakaev fkh (2009), ‘Mineral processing by the abrasive-reactive wear’, Int J Mineral Processing, 92, 58–66. urakaev fkh and avvakumov vv (1978), ‘Mechanism of mechanochemical reactions in grinding apparatus’, Izv Sib otd AN SSS: Ser khim nauk, 3/7, 10–16 (in Russian). urakaev fkh and boldyrev vv (1999a), ‘Calculation of the physicochemical parameters of mechanochemical reactors’, Inorg Mat, 35, 189–96. urakaev fkh and boldyrev vv (1999b), ‘Formation of X-ray amorphous material during mechanical activation (by the example of NaCl)’, Inorg Mater, 35, 302–5. urakaev fkh and boldyrev vv (1999c), ‘Kinetics of mechanochemical processes in grinding machines’, Inorg Mater, 35, 405–12. urakaev fkh and boldyrev vv (2000a), ‘Mechanism and kinetics of mechanochemical processes in comminuting devices 1. Theory’, Powder Technol, 107, 93–107. urakaev fkh and boldyrev vv (2000b), ‘Mechanism and kinetics of mechanochemical processes in comminuting devices 2. Applications of the theory. Experiment’, Powder Technol, 107, 197–206. urakaev fkh and shevchenko vs (2007), ‘Phenomenology, kinetics and application of abrasive-reactive wear during comminution (Overview)’, KONA (Powder and Particle), 25, 162–79. urakaev fkh, avvakumov vv and jost h (1982), ‘Mechanodestruction of silver oxalate’, Izv Sib otd AN SSSR: Ser khim nauk, 7/3, 9–14 (in Russian). vasil’ev ls and lomaeva sf (2003), ‘Mechanism of saturation of nanocrystalline powders with interstitial impurities upon mechanical dispersion’, Colloid J, 65, 639–47. vick b and furey mj (2001), ‘A basic theoretical study of the temperature rise in sliding contact with multiple contacts’, Tribol Int, 34, 823–9. vick b and furey mj (2003), ‘An investigation into the influence of frictionally generated surface temperatures on thermionic emission’, Wear, 254, 1155–61. vick b, furey mj and iskandar k (2000), ‘Theoretical surface temperatures generated from sliding contact of pure metallic elements’, Tribol Int, 33, 265–72. zegzhda sa (1997), Collision of Elastic Solids, Saint Petersburg State University, 316 pp (in Russian).

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3 Kinetic behaviour of mechanically induced structural and chemical transformations G. M U L A S, Università degli Studi di Sassari, Italy; and F. D E L O G U, Università degli Studi di Cagliari, Italy

Abstract: The kinetics of mechanically activated chemical reactions is an open field of investigation, with a number of mature problems that need to be dealt with in order to promote significant progress in mechanochemistry. Attention is here focused on the considerable efficiency exhibited by mechanochemical reactions when suitably referred to individual deformation events. Three different cases are covered in detail, which concern the degradation of toxic pollutants, the storage of hydrogen in hydride form and the enhancement of catalyst activity to produce hydrocarbons. All of these have been selected in view of their importance for future energetic scenarios and environmental policies, the emergence of which will modify the current assets of industrialized countries. Key words: mechanochemistry, kinetics, gas–solid reactions, pollutants degradation, catalysis.

3.1

Introduction

Concerns relating to the increasing demand for energy and to the impact of present technologies on environment and climate are forcing industrialized and developing countries to seek alternative solutions towards providing a sustainable future (Holdren, 2008). This is a tremendous challenge to science and engineering, which are called in to convert the unsustainable practices of industrialized countries to sustainable ones while improving the sustainability of living standards in developing countries (Holdren, 2008). The development of a fully sustainable future requires, on the one hand, more efficient utilization of resources and energy (Holdren, 2008). On the other, it implies a reduction in both wastes and in the environmental impact of human activities (Holdren, 2008). Both these issues have to cope with the needs of advancing fundamental and applied science (DOE, 2007). However, it must be noted that the demand for a sustainable future must necessarily pass across a number of intermediate scenarios in which the major role will be played by suitable rethinking of known methodologies 45 © Woodhead Publishing Limited, 2010

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(DOE, 2007). Within this framework, mechanochemistry has the opportunity to attract considerable interest. First, it is easy to apply and simple, even though its simplicity is only apparent. Second, under suitable conditions it allows very fast chemical processes. This chapter will focus precisely on these latter features, after a brief general introduction to mechanochemistry itself.

3.2

Mechanochemistry

Mechanochemistry is a form of chemistry activated by non-hydrostatic mechanical stresses (Heinicke, 1984; Levitas, 2004). In general terms, it results from the application of mechanical forces to solid phases and their consequent deformation (Heinicke, 1984; Butyagin, 1989; Gutman, 1998; Suryanarayana, 2001; Levitas, 2004). At the atomic level, the mechanical deformation of crystalline structural arrangements can be regarded as a distortion or modification of the coordination shells of individual atoms (Heinicke, 1984; Butyagin, 1989; Gutman, 1998; Suryanarayana, 2001; Levitas, 2004). In turn, the atomistic processes can be generically described as local structural excitations (Heinicke, 1984; Butyagin, 1989; Gutman, 1998; Suryanarayana, 2001; Levitas, 2004). These excitations pull the system away from thermodynamic equilibrium, which usually promotes a significant enhancement in the chemical reactivity (Gutman, 1998; Levitas and Zarechnyy, 2006; Delogu and Cocco, 2006). Local structural excitations are expected to have characteristic nature and lifetimes, which depend on a variety of intertwined factors including crystalline lattice geometry, intensity of mechanical stresses and temperature (Heinicke, 1984; Butyagin, 1989). Owing to the fast dynamics of relaxation processes in the solid phase (Delogu and Cocco, 2006), local excited states can be expected to have quite short lifetimes, roughly on the order of nanoseconds (Butyagin, 1989; Delogu and Cocco, 2006). In spite of this, promotional effects on the chemical reactivity can be expected whenever excited atoms involved in local structural excitations directly interact with other chemical species (Henicke, 1984; Butyagin, 1989). According to the brief description above, mechanochemical reactions rely substantially upon the response of any given solid phase to plastic mechanical deformation (Heinicke, 1984). In the most widespread methodology, mechanical deformation is operated by processing powders in socalled ball mills (Heinicke, 1984; Suryanarayana, 2001). Ball mills exhibit various geometries and mechanical actions (Suryanarayana, 2001). However, apart from the differences related to the design details, ball mills usually consist of stainless steel reactor chambers fixed either to platforms or to mechanical arms which are moved by electrical motors along eccentric trajectories (Suryanarayana, 2001). The powder charge is placed into the

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reactor chamber together with suitably chosen stainless steel balls. The reactors move along a selected trajectory such that the balls inside are forced to collide with each other and with the walls of the reactor chamber. When a powder charge is also present, a fraction of powder particles become trapped between the colliding surfaces of the milling tools. During the course of such events, relatively intense impulsive forces apply to the particles, inducing severe plastic deformation and a variety of correlated processes (Heinicke, 1984; Butyagin, 1989; Suryanarayana, 2001; Delogu et al., 2004). In turn, plastic deformation induces the accumulation of point, line and planar defects into the solid phase crystalline lattice (Heinicke, 1984; Honeycombe, 1984). Structural defects give rise to distorted local structures characterized by an intrinsic enhancement of chemical reactivity (Suryanarayana, 2001). The above-mentioned defective structures are static ones, that is, they refer to the content of lattice defects finally attained as a consequence of the relaxation processes following the mechanical deformation event. However, it can be reasonably expected that an enhancement in chemical reactivity could also be exhibited by atomic species involved in local deformation-induced rearrangements during the course of the deformation process and not only when the system has reached local equilibrium. The atoms undergoing rapid modifications of their local arrangement caused by the operating mechanical stresses necessarily exhibit different chemical properties from atoms at rest in a given structural configuration, whatever it is (Delogu, 2009). The local excited states connected with ongoing deformation processes are expected to trigger chemical transformations with possibly different kinetics and thermodynamics from relaxed defective systems, under conditions where the atoms participating in local excitations interact with chemical species in the surroundings (Delogu, 2009). The capability of mechanical activation to enhance chemical reactivity has already found various applications in areas of science and technology as different as powder metallurgy (Suryanarayana, 2001), mineral processing (Balaz, 2008) and organic synthesis (Rodriguez et al., 2007). Mechanical processing of powders has proved particularly valuable in the production of metastable phases such as amorphous alloys and nanostructured phases (Suryanarayana, 2001). However, it also attracts considerable interest within the framework of studies concerning hydrogen storage (DOE, 2003), pollutant abatement (Mulas et al., 1997, Loiselle et al., 1997, Monagheddu et al. 1999) and heterogeneous catalysis (DOE, 2007). Mechanical activation has only recently begun to exhibit all of its intrinsic potential despite being utilized in a variety of research activities for some time. There are a few main reasons for this (Heinicke, 1984; Butyagin, 1989; Gutman, 1998; Suryanarayana, 2001; Levitas, 2004). In the first place, the

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large amount of experimental data is extremely difficult to correlate owing to the variety of working conditions and experimental apparatuses. Second, solid phase structural transformations and interface processes exhibit a considerable degree of complexity, so that the characterization of the reactive behaviour is generally unsatisfactory. Finally, a suitable mathematical formalism to address properly thermodynamic and kinetic studies by taking account of the dynamics of powder processing in experimental devices is still lacking. Nevertheless, all of the above-mentioned points can be satisfactorily dealt with (Delogu et al., 2009; Delogu and Mulas, 2009). The emerging result provides striking experimental evidence common to all mechanically activated processes (Delogu et al., 2009; Delogu and Mulas, 2009). The rate of a given structural or chemical transformation induced by mechanical deformation is orders of magnitude higher than the one observed for the same transformation under thermal activation conditions (Delogu et al., 2009; Delogu and Mulas, 2009). This surprising evidence naturally comes from a detailed analysis of mechanically activated transformations as a function of local deformation events instead of time (Delogu et al., 2009; Delogu and Mulas, 2009). Focusing precisely on this latter issue, the present chapter aims to show the efficiency, and then the potential, of mechanical activation processes in the areas of hydrogen storage, toxic chemical degradation and catalytic transformations.

3.3

Outline of experimental methods

All of the mechanically activated transformations discussed below were carried out in a Mixer/Mill mod. 8000 from Spex CertiPrep Inc. (Metuchen, NJ, USA), which is one of the most popular commercial ball mills used in research. The milling device is equipped with a stainless steel cylindrical reactor, the internal chamber of which has radius and height of about 2 cm and 6 cm, respectively. Once the desired mass mp of powder and number of stainless steel balls are introduced, the chamber is sealed in the desired atmosphere and fixed on the clamping assembly of the mechanical arm. The arm is connected to the mill electrical motor through an eccentric fulcrum, which allows the reactor to move along a three-dimensional trajectory that can be described as a combination of rotations and translations. The frequency of such motion amounts to about 14.6 Hz (Delogu et al., 2004). As briefly mentioned earlier, the balls undergo a sequence of collisions with each other and with the reactor walls during the mill operation. At each collision, a small fraction of the powder charge is trapped between a pair of colliding surfaces and submitted to mechanical processing (Delogu et al., 2004).

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The fraction of powder charge trapped in individual collisions exhibits a dependence on the powder charge mp itself. In particular, the powder fraction trapped at collision increases on the average with the powder mass mp. A parallel increase in the inelasticity of collisions takes place. Almost perfectly inelastic collisions already occur when the powder charge occupies at least the 1% of the total volume of the reactor chamber (Manai et al., 2001; Delogu et al., 2004). For typical transition metal and transition metal oxide systems, such circumstances allow the processing of a mass m* of powders per collision roughly equal to 1 mg. This powder mass, relatively small when compared to typical powder charges mp of about 8 g, is processed at collision for a time interval that can be equalled to the time duration τ of the inelastic collision itself, which is roughly equal to 1 ms (Manai et al., 2001; Delogu et al., 2004). Owing to the particular dynamics of balls and powders inside the reactor chamber (Manai et al., 2001), both mp and τ are approximately insensitive to the exact amount mp of powder. Whereas the use of a relatively large amount of powders ensures an inelastic collision regime, the additional use of a single milling ball permits the establishment of regular milling dynamics (Manai et al., 2001). In fact, the ball can only undergo collisions with the internal reactor walls and with the base of the container in particular (Manai et al., 2001; Delogu et al., 2004). This allows the ball to follow periodic trajectories with two collisions per cycle of reactor displacement (Manai et al., 2001; Delogu et al., 2004). More specifically, collisions occur at a frequency N of about 29.2 Hz (Manai et al., 2001; Delogu et al., 2004). All the experimental trials were carried out with a powder charge mp of solid reactants or catalysts equal to 8 g. These powders were invariably manipulated under an inert argon atmosphere during the reactor chamber charge and discharge operations. With this aim in mind, a glove box with oxygen, nitrogen and water content below 2 ppm was used. When necessary, the atmosphere inside the reactor chamber was suitably modified by carrying out filling–emptying–refilling cycles with the desired pure gas or gaseous mixture. In these cases, a suitably manufactured cylindrical reactor was utilized, with the top and bottom caps equipped with valves for the inlet and outlet of gases. The reactor chamber can, in principle, be kept at any desired temperature, although within specific ranges depending on the leak-proof ‘O-ring’ systems. Suitable external jackets allow the circulation of heating and cooling fluids. In the absence of jackets, simple air jet cooling suffices to keep the average system temperature in the range between 300 and 330 K. The average temperature is measured by a thin lamella-shaped Ptresistance thermometer fixed on the external reactor wall or embedded in it. As for the local temperature experienced by powders, indirect clues

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suggest that the local rise caused by collisions amounts to about 100 K (Delogu and Cocco, 2008a). Analyses of solid phases were carried out by wide-angle X-ray diffraction (XRD), performed on powder sampled at selected times during mechanical treatment. XRD patterns were generally collected with a Rigaku D/Max diffractometer equipped with a Cu Kα radiation tube and a graphite monochromator in the diffracted beam. N2 adsorption was used to measure the specific surface area Sp of the different solid phases according to the BET method (Thomas and Thomas, 1997). Analytical tests employed a Fisons Sorptomatic 9600 apparatus. Liquid and gaseous phases were analysed by gas chromatography resorting to different detectors. In the former case, a Finnigam Tracker Mass Spectrometer GC/MS PE 8420 device was used. A Fisons 8000 apparatus equipped with a HWD detector was employed to monitor CO and H2 gases, whereas light hydrocarbons and other gaseous organic phases were analysed with a Perkin-Elmer 8600 apparatus equipped with a FID detector. High purity solid, liquid and gaseous reactants were invariably used.

3.4

Degradation of chlorinated aromatics

The behaviour of two different chlorinated aromatic compounds was selected for study in the presence of mechanically activated solid reactants, namely the chlorobenzene (C6H5Cl) and hexachlorobenzene (C6Cl6). At room temperature the former compound is in the liquid phase, whereas the latter is a solid. Degradation reactions were carried out over inorganic substrates such as MgO, CaO and CaH2. Reactants were mixed in order to have an atomic ratio of alkaline-earth metal to chlorine ratio equal to 15 : 1 (Mulas et al., 1997). The two toxic compounds exhibit markedly different reactivity. In particular, the mechanical treatment of C6Cl6 with CaH2 generally induces a very fast chemical conversion, mostly into benzene (C6H6) and CaCl2 with extremely large exothermic effects (Mulas et al., 1997). In contrast, the mechanical processing of C6H5Cl with CaH2 results in a gradual reaction, also yielding the same above-mentioned main products (Mulas et al., 1997, Loiselle et al., 1997). The chemical behaviour triggered by the mechanical activation of the reactant mixtures containing C6Cl6 is exemplified by the temperature profile shown in Fig. 3.1. Here, the temperature T of the mechanochemical reactor is reported as a function of the time t of mechanical processing. It can be seen that the reactor temperature T first undergoes a gradual increase connected with the thermal effects originated from the partial dissipation at impact of the kinetic energy of the single milling ball. The quantity of heat dissipated at collisions is a linear function of time t, so

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Temperature (K)

340 330 320 310 300 290 0

2000

4000 6000 Time, t (s)

8000

10000

3.1 The temperature (K) of the stainless steel reactor as a function of the time t of mechanical processing. Data refer to the reaction of CaH2 and hexachlorobenzene. The temperature spike marks the onset of the self-propagating high-temperature reaction.

that a rough estimate of this quantity can be obtained by the initial slope of the temperature curve. However, the process of heat generation at impact is soon counterbalanced by the heat dissipation process at the reactor external surfaces. The result is a curved trend that represents a compromise between heat generation and dissipation processes. After a certain milling time has elapsed, it is possible to observe a sharp thermal spike superposed on the gradual temperature rise and in a sense independent of it. The sudden temperature rise marks the onset of a very fast chemical reaction, the characteristics of which are only partially known. However, it is possible to state a few important features. First, the chemical process is self-propagating, that is, it does not need any further energy supply once triggered. Second, the chemical reaction occurs on very short timescales, roughly on the order of 1 s. Third, the exothermal effects perfectly correlate with the reaction enthalpy, so that no significant amount of intermediate species is present. Fourth, the reaction probably involves the formation of radical species, as suggested by the detection of negligible traces of chemical compounds, the formation of which relies upon radical condensation reactions. The evidence that mechanical processing succeeds in activating selfpropagating behaviour immediately points out the difference between mechanical activation and conventional thermally activated processes. In fact, no self-sustaining reaction takes place when the reactants are submitted to a progressive temperature increase, unless very high heating rates are imposed. On the contrary, the chemical reaction possibly proceeds at relatively low conversion rates. This does not mean that the self-propagating process cannot be triggered under purely thermal conditions. Nevertheless, such conditions must be quite severe in terms of heating rate and total

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energy transferred to reactants. More specifically, conditions must correspond to the ones generally met by self-sustaining high-temperature synthesis (Munir and Anselmi-Tamburini, 1989). It follows that mechanical activation is able to impose equally severe processing conditions on reactant powders (Takacs, 2002). In contrast to the previous case, the mechanical treatment of reactant mixtures containing C6H5Cl only results in a gradual conversion of reactants in the final main products. A typical kinetic curve is shown in Fig. 3.2, where the molar fraction of C6H5Cl reacted is reported as a function of the time t of mechanical processing. It can be seen that the conversion is described by a sigmoidal curve, which indicates that the reaction reaches completion on timescales on the order of 12 h. The gradual conversion curve can be used to gain insight into the reactive behaviour at individual collisions. Starting from evidence that the chemical reaction stops immediately when the mechanical processing is interrupted, the observed conversion rate can be connected with the frequency N of collisions. With this aim in mind, a necessary assumption is that the chemical events activated by individual deformation events at each collision are independent of each other. Correspondingly, the local chemical reaction initiated by a given collision is not influenced by the following one. In general terms, this assumption can appear quite arbitrary. However, it is based on evidence that local structural excitations induced by deformation events exhibit relaxation times on the order of 1 ns (Delogu, 2009) and involve small percentages of the powder charge, typically on the order of 0.1% (Delogu et al., 2004). Taking into account that successive collisions are separated by time intervals of about 30 ms (Delogu et al., 2004), it must be expected that any chemical reactivity triggered by mechanical deformation at a given collision is definitely quenched when the immediately successive collision occurs. Under these circumstances, the overall chemical reaction corresponds to the resultant of the whole sequence of discrete reactive events at individual collisions. Consequent to such assumption, the instantaneous conversion rate r at any selected point on the kinetic curve in Fig. 3.2 can be referred to individual collision events by dividing it by the frequency N of collisions. This provides a measure of the amount of reactants reacted at each collision. In the present case, the maximum conversion rate, referred to C6H5Cl, is attained at about 4 h and amounts to about 4.8 × 10−7 mol s−1. The frequency N of collisions being equal to about 29.2 Hz, the amount of C6H5Cl transformed at each collision is roughly equal to 1.6 × 10−8 mol. This value can be utilized to estimate the specific conversion rate of reactants at individual collision events. With this aim, the above-mentioned value must be referred to the mass m* ≈ 1 mg of powder trapped during the collision and to the time duration τ ≈ 1 ms of collisions, which represents a reasonable estimate

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Conversion α (%)

100 80 60 40 20 0 0

2

4

6

8

10

12

14

Time, t (h)

3.2 Fraction α of reacted chlorobenzene as a function of the time t of mechanical processing. Data refer to the reaction with CaH2.

of the time interval during which mechanical forces operate on powders. The resulting specific reaction rate amounts then to about 1.64 × 10−2 mol g−1 s−1. A way to interpret the above-mentioned result is by comparing the observed overall reaction rate and the possible overall reaction rate in the case in which negligible time intervals separate successive collisions. In the former case, the gradual degradation of C6H5Cl reaches 80% completion after about 8.5 × 105 collisions, that is, after 8 h. In the second case, provided that the same number of collisions is required, 80% conversion would be reached after only a few seconds of treatment. Therefore, the mechanochemical process in principle exhibits considerable efficiency, with very high specific rates of chemical reaction. In spite of the general observations above, it is worth noting here that this aspect can be adequately pointed out only through a suitable comparison of mechanochemical processes and thermal ones. This approach has not been possible in the present case, since the thermal activation of reactant powders does not produce any gradual conversion in the case of C6Cl6, whereas it induces the evaporation of C6H5Cl in the other. However, comparison of mechanically and thermally activated reactions will be the focus of the following sections.

3.5

Hydrogen absorption on Mg2Ni/Ni composite

The H2 absorption processes were carried out on powder charges mp of 8 g of a Mg2Ni/Ni composite. A stainless steel ball of 8 g was used. The H2 pressure was kept constant at 0.4 MPa (Delogu and Mulas, 2009). The kinetics of the H2 absorption process is described by the curve shown in Fig. 3.3, where the number n of H2 moles absorbed is reported as a function of the time t of mechanical processing. The curve exhibits

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0.12 ns+d 0.08

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ns 0.04

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3.3 Number n of H2 moles absorbed as a function of the time t of mechanical processing. The vertical dotted line separates the H2 absorption processes performed under static and dynamic conditions, whereas the horizontal dotted lines indicate their respective plateau values ns and ns+d.

two distinct portions relating to the H2 absorption processes carried out with the mechanochemical reactor, respectively, at rest and in motion. The two cases will be hereafter referred to as static and dynamic processing conditions. It can be seen that the number n of H2 moles absorbed under static conditions tends to reach a plateau value ns. The mechanical treatment was started at time t0, roughly 60 min before the number of hydrogen moles absorbed could reach the plateau value under static processing conditions, that is, when powders have already absorbed n0 ≈ 4.8 × 10−2 H2 moles. The mechanical processing induces a definite increase in the number n of H2 moles absorbed, which however tends to reach a plateau value ns+d larger than ns. The observed behaviour suggests that the H2 absorption reaction over Mg2Ni/Ni under static conditions is limited by some sort of thermodynamic or kinetic limitations. However, these are by-passed under dynamic conditions, which promote a significant absorption beyond the apparent asymptotic threshold ns. According to XRD analyses, the above-mentioned behaviour can be ascribed to a change in the H2 absorption kinetics. In fact, static processing conditions substantially induce the formation of the Mg2NiH0.3 hydride, which is relatively poor in H content. When the same powders have also been mechanically processed, the predominant phase is instead the Mg2NiH4 hydride, which is the hydride richest in H formed by Mg2Ni systems. Static and dynamic processing conditions produce, therefore, very different results. Under static conditions, H2 is absorbed according to an apparent secondorder rate equation: dn 2 = Ks ( ns − n) [3.1] dt the solution of which is © Woodhead Publishing Limited, 2010

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10

n/(ns–n)

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3.4 n/(ns − n) as a function of the time t of mechanical processing. The line best-fit is also shown.

n = Ks ns t ns − n

[3.2]

Accordingly, the first portion of the curve in Fig. 3.3 can be linearized by reporting the quantity n/(ns − n) as a function of time t. The resulting plot is shown in Fig. 3.4. According to Equation [3.2], the slope of these linear trends is equal to Ksns, where Ks is the apparent rate constant of the H2 absorption process under static conditions. The best-fit of data in Fig. 3.4 indicates that ns and Ks amount to about 5.4 × 10−2 mol and 2.6 mol min−1, respectively. It immediately appears that ns is far from the value of 0.148 mol necessary for complete conversion into Mg2NiH4, but is about 5 times larger than the 1.1 × 10−2 mol necessary for the complete conversion into Mg2NiH0.3. Under dynamic conditions, the H2 absorption kinetics exhibits an exponential character. The scaled number n* = n − n0 of H2 moles absorbed under dynamic conditions and the scaled time t* = t − t0 were defined suitably to analyse the second portion of the curve in Fig. 3.3. As shown in Fig. 3.5, the quantity ln(n*s+d − n*), being n*s+d = ns+d − n0, changes linearly with t*. It follows that: ln(n*s+d − n*) = K*s+dt* + lnn*s+d

[3.3]

which is a solution of the rate equation: dn * = K*s+d (n*s+d − n*) dt *

[3.4]

The quantity K*s+d represents the apparent rate constant for the H2 absorption process performed under dynamic conditions on powders that have already absorbed n0 H2 moles after a time interval t0 under static conditions. The best-fit of data in Fig. 3.5 yields for n*s+d and K*s+d estimates of about 6.0 × 10−2 mol and 4.5 × 10−3 min−1, respectively. It is worth noting here that ns © Woodhead Publishing Limited, 2010

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ln(n*s+d –n*)

–3 –4 –5 –6 –7 –8 0

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Time, t* (min)

3.5 ln(n*s+d − n*) as a function of the scaled time t*. The line of best-fit is also shown.

+ n*s+d ≈ 0.114 mol represents the total number of H2 moles absorbed. The above-mentioned experimental findings allow a quantitative comparison of experimental data obtained under static and dynamic conditions. Owing to the exponential character of the kinetic curve, the H2 absorption under dynamic conditions can be referred to independent reactive events taking place at individual collisions (Delogu and Cocco, 2007; Delogu and Cocco, 2008b). Under these conditions, the rate of mechanically activated processes only depends on the mass m* ≈ 1 mg of powders involved in individual collisions. In addition, H2 absorption at collision should take place for time periods roughly equal to the impact duration τ ≈ 1 ms. When the mechanical treatment starts, the scaled number n* of H2 moles absorbed under dynamic conditions is equal to zero. After the first collision has occurred, the number of H2 moles absorbed is approximately equal to K*s+dn*s+d, which then represents the initial rate rs+d of H2 absorption. This quantity amounts to about 4.5 × 10−6 mol s−1, which means that about 1.5 × 10−7 mol of H2 are absorbed at the first collision. This H2 amount is absorbed by roughly 1 mg of powders within time intervals of about 1 ms. Therefore, the specific H2 absorption rate at the first collision is equal to about 0.15 mol g−1 s−1. This value must be compared with the specific rate of H2 absorption under static conditions approximately after n0 H2 moles have been absorbed in a time period t0. It follows from Equation [3.1] that the rate rs of H2 absorption under static conditions is equal to Ks(ns − n0)2, which roughly amounts to 1.6 × 10−6 mol s−1. The specific H2 absorption rate under static conditions can now be obtained by dividing this rs value by the total powder mass mp = 8 g. The specific rate amounts to about 2.0 × 10−7 mol g−1 s−1. According to the above-mentioned specific rate values, H2 absorption under dynamic conditions is about six orders of magnitude faster than under static conditions. This difference cannot be explained in terms of

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specific surface areas. In fact, the Mg2Ni/Ni powders interacting with H2 at individual collisions should exhibit a specific surface area S*p on the order of 1 × 105 m2 g−1, values characteristic of catalytic carbon or silica aerogels. In contrast, the BET specific surface area Sp is equal to only about 15 m2 g−1 before and after any mechanical treatment. In the light of calculations above, the enhanced H2 absorption capability of mechanically processed powders must find a different rationalization. This can be looked for in so-called mechanochemical effects, that is, in effects ascribable to the highly excited atomic structures that appear at the surface of powder particles as a consequence of deformation processes. More specifically, the extremely high specific rates of H2 absorption at individual collision events can be related to the increase of the surface density of reactive sites, to the enhancement of their chemisorption efficiency or to both these factors. Rough calculations carried out to pursue research further along these lines indicate that the activation energy Ea of the H2 chemisorption should drop from about 64 to 37 kJ mol−1. Alternatively, local temperatures should rise by about 400 K. Of course, no discrimination is possible between surface, temperature and energy effects. In contrast, it is highly probable that the enhanced H2 absorption capability of mechanically processed powders could result from a combination of all such effects.

3.6

Catalytic hydrogenation of carbon monoxide

The hydrogenation of carbon monoxide (CO) to produce methane (CH4) and other hydrocarbons was carried out over 8 g of a powder catalyst with a composition of (Co50Fe50)0.2(TiO2)99.8. A gaseous mixture of H2 and CO was introduced into the reactor chamber with a 1 : 3 CO:H2 ratio to maximize the reaction yield in CH4. The total pressure inside the reactor chamber was set at 0.3 MPa (Delogu et al., 2009). The conversion data for the CO hydrogenation process are given in Fig. 3.6, where the moles n of hydrocarbon produced are shown as a function of the time t of mechanical processing. Substantially, CO is selectively converted to CH4 with an initial rate of about 7.8 × 10−7 mol h−1. This conversion must be ascribed to the metallic Co-Fe phase, which, in 8 g of catalyst, is equal to 0.023 g. The amount r of CO converted on the average in individual impact events can be estimated by dividing the overall conversion rate by the frequency N of collisions, which is equal to about 29.2 Hz. Roughly 7.4 × 10−12 mol of CO are converted at each collision. Remembering that the time duration τ of collisions is equal to about 1 ms and that the mass m* of powder involved in each collision amounts roughly to 1 mg, the specific conversion rate r* = r τ−1m*−1 at individual collisions amounts to approximately 7.4 × 10−6 mol g−1 s−1. The specific

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3.6 n moles of hydrocarbon produced over the (Co50Fe50)0.2(TiO2)99.8 catalyst as a function of the time t of mechanical processing.

catalytic activity must be referred to the area s* of the catalytically active surface involved in individual collisions. This corresponds to a turnover frequency fmech, that is, the number of reactant moles reacted per time unit per active site of the catalytically active surface under mechanochemical −1 conditions. The quantity fmech is equal to rτ−1s*C−1 o Feρs , where ρs is the average surface atomic density of the Co50Fe50 solid solution. As s*CoFe = m*CoFeSp,CoFe, −1 the expression above is also equal to r*S*C−1 o Feρs , where Sp,CoFe is the specific surface area of only the Co50Fe50 solid solution, which is the catalytically active phase. The average surface atomic density ρs can be estimated by the arithmetic mean of the atomic densities of individual Co and Fe surfaces reported in the literature (Anderson, 1975). A ρs value of about 1.6 × 1019 at m−2 is obtained. The quantity s*CoFe can be estimated under the hypothesis that the Co50Fe50 domains supported on TiO2 correspond to individual crystallites. This hypothesis permits the estimation of the maximum specific surface area Sp,CoFe available to reaction by considering spherical Co50Fe50 domains with a diameter equal to the average crystallite size L. In this case, the maximum specific surface area Sp,CoFe is equal to 6ρ−1L−1, where ρ is the average density of the Co50Fe50 solid solution. Evaluating this quantity as the arithmetic average of Co and Fe densities, ρ amounts to about 8.5 g cm−3. As L is roughly equal to 20 nm, the maximum specific surface area Sp,CoFe amounts to about 35 m2 g−1. The total surface area s*CoFe of the Co50Fe50 domains available for reaction in individual collisions is therefore roughly equal to 9.8 × 10−5 m2, and the mass m*CoFe of Co50Fe50 phase involved in collisions equal to about 2.8 × 10−6 g. The mechanochemical turnover frequency fmech of the CO hydrogenation process is then equal to about 2.84 molecules atom−1 s−1. The fmech value obtained for the (Co50Fe50)0.2(TiO2)99.8 catalyst can be now compared with the value of turnover frequency fth obtained over a similar

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catalyst under thermal activation conditions (Duvenhage and Coville, 1997). This catalyst has a (Co50Fe50)5(TiO2)95 chemical composition and was tested under flow conditions at a temperature of 493 K and a total pressure of a 1 : 2 CO : H2 mixture equal to 1 MPa (Duvenhage and Coville, 1997). The working conditions are therefore slightly different from the ones employed during the course of mechanical activation. The thermal turnover frequency fth for this catalyst amounts to about 1.4 × 10−2 molecules atom−1 s−1 (Duvenhage and Coville, 1997). Therefore, at least two orders of magnitude separate the above mentioned fth value from the mechanochemical turnover frequency fmech of about 2.84 molecules atom−1 s−1 obtained for the (Co50Fe50)0.2(TiO2)99.8 catalyst. The difference of two orders of magnitude represents only the minimum possible difference between the activity of the two catalysts. A few simple considerations suggest, in fact, that the difference in specific activity, or in turnover frequencies, must be larger. First, both temperature and pressure are higher in the case of thermal activation, which assures higher frequencies of collisions between gaseous species and catalyst surface. Second, the actual specific surface area Sp,CoFe value of the supported Co50Fe50 phase is expected to be smaller than the rough estimate based on the average crystallite size, since individual crystallites are expected to form nanostructured aggregates. Third, the time interval τ over which the chemical transformations are assumed to occur has been set equal to the time duration of an inelastic collision. However, such an estimate only represents an upper bound for the collision duration, since partially inelastic collisions often last only 0.1 ms (Manai et al., 2001). It follows that the mechanochemical turnover frequency fmech could be two to three orders of magnitude larger than the thermal fth one. Correspondingly, the mechanically activated catalytic process could be two or three orders of magnitude more efficient than the thermally activated one. This provides an indirect estimate of the importance of mechanochemical effects on the chemical behaviour of solid phases.

3.7

Future trends

The kinetics of mechanically activated processes is an area of research with a number of open questions. The brief overview presented here is specifically aimed at pointing out a few outstanding topics in mechanochemical kinetics, the investigation of which could promote significant advances in the field. According to the data and approaches discussed, there is a lack of fundamental knowledge about the mechanisms with which mechanically activated processes take place at individual collisions. In fact, the experimental evidence only refers to overall transformations observed on the macroscopic scale. Referring to individual collisions represents an important,

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though still unsatisfactory, step toward the microscopic scale. There is a stringent need to perform careful investigations from both the experimental and theoretical points of view on the chemical response of materials embedded in a reactive environment while being submitted to ongoing mechanical deformations. The challenging objectives can be summarized in two main points of a middle-term programme: • •

set-up of experimental methodologies to investigate the chemical reactivity of solid, liquid and gaseous reactants at individual collisions; development of theoretical and numerical models able to address the fundamental issues intrinsically connected with a far-from-equilibrium chemistry.

Both these points imply considerable effort in tentative exploration of chemical reactivity under extreme conditions at the limit of application of present technologies. Which method can be applied to investigate the way a given amount of kinetic energy is distributed at individual collisions? How can reliable information be obtained at the very short timescales characteristic of collisions? How do portions of matter really behave when submitted to ongoing deformation during interaction with liquid and gaseous environments? These and a number of other questions regarding the intimate nature of mechanochemical processes clearly emphasize the unsatisfactory degree of knowledge in the field at present. There is a strong need for cooperation between research groups with different expertise, as well as to identify viable applications of current techniques to practical problems in mechanochemistry. New ideas are also needed that could permit a significant breakthrough in this field. And all of this must happen in the shortest possible time, if mechanochemistry has to play a role in the intermediate energetic scenarios that are becoming every day closer.

3.8

References

anderson j r (1975), Structure of Metallic Catalysts, Academic Press, London. balaz p (2008), Mechanochemistry in Nanoscience and Minerals Engineering, Springer, Berlin. butyagin p yu (1989), ‘Active states in mechanochemical reactions’, Sov Sci Rev Chem B, 14, 1–134. delogu f (2009), ‘Molecular dynamics of perturbed rearrangements at perturbed interfaces’, Phys Rev B, 80, 014115. delogu f and cocco g (2006), ‘Numerical simulations of atomic-scale disordering processes at impact between two rough crystalline surfaces’, Phys Rev B, 74, 035406. delogu f and cocco g (2007), ‘The size refinement of Cu crystallites under mechanical processing conditions: a phenomenological modeling approach’, J Mater Sci, 42, 4356–63.

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delogu f and cocco g (2008a), ‘Weakness of the “hot spots” approach to the kinetics of mechanically induced phase transformations’, J Alloys Compd, 465, 540–6. delogu f and cocco g (2008b), ‘Kinetics of amorphization processes by mechanical alloying: a modeling approach’, J Alloys Compd, 436, 233–40. delogu f and mulas g (2009), ‘Hydrogen absorption processes in Mg2Ni-based systems: thermal and mechanochemical kinetics’, Int J Hydrogen Energy, 34, 3026–31. delogu f, mulas g, schiffini l and cocco g (2004), ‘Mechanical work and conversion degree in mechanically induced processes’, Mater Sci Eng A, 382, 280–7. delogu f, mulas g, and garroni s (2009), ‘Hydrogenation of carbon monoxide under mechanical activation conditions’, Appl Catal A: Gen, 366, 201–5. doe report (2003), Basic Research Needs for the Hydrogen Economy, Report from the US Department of Energy (DOE) Basic Energy Sciences Workshop on Hydrogen Production, Storage, and Use, Argonne National Laboratory, Illinois (USA) May 13–15 (available at http://www.sc.doe.gov/bes/hydrogen.pdf). doe report (2007), Basic Energy Needs: Catalysis for Energy, Report from the US Department of Energy (DOE) Basic Energy Sciences Workshop, Bethesda, Maryland (USA), August 6–8 (available at http://www.sc.doe.gov/bes/reports/list. html). duvenhage d j and coville n j (1997), ‘Fe:Co/TiO2 bimetallic catalysts for the Fischer–Tropsch reaction. I. Characterization and reactor studies’, Appl Catal A: Gen, 153, 43–67. gutman e m (1998) Mechanochemistry of Materials, Cambridge International Science Publishing, Cambridge. heinicke g (1984), Tribochemistry, Akademie-Verlag, Berlin. holdren j p (2008), ‘Energy and sustainability’, Science, 315, 737. honeycombe r w k (1984), The Plastic Deformation of Metals, Edward Arnold, Baltimore. levitas v i (2004), ‘Continuum mechanical fundamentals of mechanochemistry’, in High Pressure Surface Science and Engineering, Gogotsi Y and Domnich V (eds), Institute of Physics, Bristol, 159–292. levitas v i and zarechnyy o m (2006), ‘Kinetics of strain-induced structural changes under high pressure’, J Phys Chem B, 110, 16035–46. loiselle s, branca m, mulas g, cocco g (1997), ‘Selective mechanochemical dehalogenation of chlorobenzenes over calcium hydride’, Environ Sci Technol, 31, 261–5. manai g, delogu f and rustici m (2001) ‘Onset of a chaotic dynamics in a ball mill: attractors merging and crisis induced intermittency’, Chaos, 12, 601–9. monagheddu m, mulas g, doppiu s, cocco g, raccanelli s (1999), ‘Reduction of polychlorinated dibenzodioxins and dibenzofurans in contaminated muds by mechanically induced combustion reactions’, Environ Sci Technol, 33, 2485–8. mulas g, loiselle, schiffini l, and cocco g (1997), ‘The mechanochemical selfpropagating reaction between hexachlorobenzene and calcium hydride’, J Sol State Chem, 129, 263–70. munir z a and anselmi-tamburini u (1989), ‘Self-propagating exothermic reactions: The synthesis of high-temperature materials by combustion’, Mater Sci Rep, 3, 277.

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rodriguez b, bruckmann a, rantanen t, and bolm c (2007), ‘Solvent-free carboncarbon bond formations in ball mills’, Adv Synth Catal, 349, 2213–33. suryanarayana c (2001), ‘Mechanical alloying and milling’, Prog Mater Sci, 46, 1–184. takacs l (2002), ‘Self-sustaining reactions induced by ball milling’, Prog Mater Sci, 47, 355–414. thomas j m and thomas w j (1997), Principles and Practice of Heterogeneous Catalysis, VCH, Weinheim.

© Woodhead Publishing Limited, 2010

4 Materials design through mechanochemical processing M A M O R U S E N N A, Keio University, Japan

Abstract: Controlled shear stressing on a dry solid mixture at ambient temperature enables charge transfer and short range interparticulate atomic transfer. They result in formation of hetero-bridging bonds, leading to nucleation of various composites or complexes, for example, perovskites or spinels. Similar phenomena could be applied to rationalizing organic syntheses or compounding of different materials in various genres. After surveying general and theoretical bases of materials design via solid state routes, case studies on electroceramics and novel organic synthesis by using moderate mechanical stressing are given using the concept of soft-mechanochemical processes, where the use of chemical spontaneity like self-organization or molecular recognition is emphasized. Key words: soft-mechanochemical processes, charge transfer, hetero-bridging bonds, perovskite ceramics, coordination compounds, Diels–Alder reaction.

4.1

Introduction

It is always a challenge to tailor materials pinpointed for particular applications. Since the characteristics of the materials are so diverse, the parameters that need to be controlled in materials design are necessarily numerous. One of the smartest methods in materials design is first to assemble atoms or molecules in a well-defined way in a given space. This is widely carried out by computer simulation.1–3 Although the concept is elegant and perfect, outcome of these studies cannot be applied to actual production processes. Various attempts have been made, to design materials and fabricate them, where atomic or molecular manipulation is regarded as an ultimate goal.4 Since this is extremely tedious, if possible at all, people have tried to let atoms self-organize themselves. This is realized to some extent in various thin film technologies.5 They can be stacked to give a three-dimensional (3D) structure, often called artificial lattices.6,7 On an industrial basis, however, the well-defined methodologies given above are mostly unaffordable. This is not just due to the very minute amount of thin film processing. Since most materials are handled in the 63 © Woodhead Publishing Limited, 2010

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form of fine particulates, we need to start by designing particulates. This is far more complicated than for thin films, since fine particulate states cannot be prepared without introducing a significant amount of lattice defects and near-surface anomalies of solids, although the finely controlled technology of two-dimensional (2D) film formation can be smartly converted into socalled 3D films, which are actually an ensemble of fine particulates. When we go down to the level of several tens of nanometers, we need to regard them as an ensemble of surface anomalies without any stable parts in the bulk solid. In addition, we have to take huge number of extra parameters into account, that is, geometrical, physical and chemical ones, since we almost always have to deal with a large ensemble of particulate matters or grains. Thus, we have to abandon sophisticated beam technology and look into various methods of preparing fine particles with reasonable scalability. This chapter will discuss how the general concepts of materials design mentioned above could be associated with mechanochemistry. The structure of the chapter is as outlined below. •

• •

• •

The benefits of a mechanochemical route for materials design: after introducing the brief histories of mechanochemical processing for materials synthesis, the importance of microscopic phenomena at solid interfaces is emphasized. The theoretical aspects of modern mechanochemistry: anomalies in the near-surface region and nano-sized particles are briefly reviewed. Composite oxides for electromagnetic application: this is the main part of the chapter, explaining the concept of soft-mechanochemical processes with examples of several case studies for perovskite ceramic particles carried out in the author's laboratory. Organic synthesis and utilization of chemical spontaneity: the role of charge transfer complex on the solid state Diels–Alder reaction under mechanical stressing is introduced. Summarizing remarks and the future outlook.

4.2

Benefits of a mechanochemical route for materials design

First, the chemical implications of mechanical stressing on the solid state will be briefly reviewed. Traditional solid state processes are represented by crushing mountain rocks to acquire various natural resources, notably metals. Powder technology in mineral processing in its early stage includes size reduction, sorting or classification, among others, by floatation. Subsequent metallurgical processes are subdivided into wet and dry processes, called hydro- and pyro-metallurgy. Mechanical activation plays a particularly significant role in hydrometallurgy, which begins with a dissolution

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process. The presence of finer particles with a larger specific surface area, together with higher reactivity at the solid/liquid interface prove to be beneficial. This is one of the origins of practical application of mechanochemistry.8,9 Thus, an enhanced rate of dissolution and apparent solubility is the main attraction in industry of using mechanical stressing on solids in a one phase system. Introduction of various lattice imperfections is key.10–12 This often leads to phase transformation13–15 and amorphization as well.16–18 While this kind of mechanical activation is of practical importance, it is not always directly associated with materials design, except in some special applications like pharmaceutics. A more interesting mechanochemical issue that is closely related to materials design is putting stress on a mixture of solids. One typical example is so-called mechanical alloying, where two or more metallic species in the form of powder are mixed in a mill, become quasi atomically homogeneous alloys,19–25 and could even exceed the solubility limits predicted by phase diagrams. Here, the process might seem quite simple, that is where different metallic species penetrate the boundary via a short-range transfer, called mechanical diffusion, as schematically illustrated in Fig. 4.1. Things are, however, not so straightforward when mixing coloured marble. Whether or not mixture really builds up is believed to be governed by the heat of mixing,26–29 although, the mechanical alloying state is far from thermodynamic equilibrium so that phase diagrams cannot be the sole measure of the phase availability of alloys. This illustrates another peculiarity of mechanical alloying, where the process goes in most cases through a multilayered structure and there is a large contact area for mechanical diffusion owing to the ductility of the ingredient metals.30,31 The concept of mechanical alloying was extended to the polymeric materials to give ‘polymer alloys’ under a more or less similar procedure,32–34 although the chemical procedure involved is quite different since the unit of species is a complicated molecule with high molecular weight. Polymer alloys, often called ‘polymer blends’ are formed by repeated breakage and (a)

(b)

(c)

4.1 Schematic illustration of forced diffusion at the contact point of dissimilar particles. This simple process explains the principle of mechanical diffusion, which is a basic process of mechanochemistry, along with charge transfer, explained below.

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rebuilding of molecular chains with and without side chains via widely varying radical species. Finally, the concept of mechanical alloying has been further extended to more general compounds involving oxides or chalcogenides. These processes are then renamed as ‘mechanosynthesis’, where products are obtained straight from the milling pot, normally without any heating process.35–38 There are a number of outstanding books in this area.39–41 At this stage, it is worth emphasizing that the technological benefits of mechanical stressing relating with materials design are mainly associated with a cross-solid interaction, triggered by many different factors, including short-range diffusion. A typical attempt exerts mechanical stress on the solid mixture to obtain alloys or composite oxides outright from the milling devices. In the case of conventional alloying, we usually need vacuum furnace and temperature higher than the melting points of ingredient metallic species, or at least higher than the eutectic temperature. Melting of oxide mixtures is in most cases out of question owing to their extremely high melting points. Many of them tend to decompose prior to melting. Therefore, the apparent merits of mechanosynthesis are quite obvious. However, there are a number of disadvantages or pitfalls. First of all, mechanosyntheses are in most cases carried out in energy intensive milling devices after operating for a long time, often hundreds of hours. A direct consequence is high power consumption, as well as high cost of devices and their maintenance. At the same time, the level of contamination from abraded materials in the milling media or vessels is high, up to several percent, and hence not tolerable for many industrial products including electroceramics. The limited extent of scale up in energy intensive milling devices counts as another demerit. Meanwhile, as energy used in industrial processes became more important, a headwind developed against mechanosynthesis, since generally milling devices are notorious for their low energy efficiency. As the costly mechanochemical processes were applied to more high-value-added materials, another serious problem arose, contamination, which becomes more serious with increasing energy intensity. It is therefore important at this stage to re-examine the merit of mechanical stressing more critically to find out what kind of solid state reactions are really enhanced by exerting external mechanical energy on the reactants and whether materials processing under mechanical stress is the best method compared to other non-conventional methods like photochemical, sonochemical, magnetochemical, hydrothermal or high-pressure processes. Here, we must concentrate more scientific effort on considering how processes related to mechanochemistry can survive in the area of materials design and in an era where there is increasing emphasis on the energy cost and ecological burden. Homogenization of solid mixtures at the atomic or molecular level was believed to be possible only by heating to elevated temperatures where the

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mobility of the transporting species are high enough. As we have used metal oxides in making ceramic products, a vast number of new technological processes have been developed with respect to materials design and syntheses. Mechanochemistry and mechanochemical processes are representative of these. Conventional ceramic processes, often called ‘shake-and-bake’, are regarded as old-fashioned, belonging to the 20th century. When we look at the solid–solid grain boundary, we recognize that charge transfer and electronic redistribution under mechanical stressing is unique. The change in the electronic states associated with cross-boundary charge transfer plays an important role in a number of functional materials. Intimate mixing under controlled shear stressing in a dry solid process at ambient temperature enables controlled charge transfer and short range atomic transfer across boundaries of solids, which no other methods, like photo-, magneto- or plasma chemical processes, could achieve. By the same token, seemingly more advanced processes of solution–chemical, colloid– chemical or sol–gel processes are not always capable of controlling the charge transfer or associated electronic band structure, even if they are superior with regard to homogeneity of mixing. It is also important to mention that these processes are developing parallel to solid state chemistry, or to be more exact, studies on the reactivity of solids. A conventional tumbling ball mill can certainly achieve intimate mixing of particles with simultaneous size reduction of individual particles, although only to a limited extent. The use of more intensive mills, such as vibromills or planetary mills, therefore became prevalent in the hope that much finer and more homogeneous starting mixtures for subsequent calcination processes could be obtained. Up until the end of the last century, energy intensive milling machines were developed for these purposes, including vibration and planetary mills. When it turned out that the technological development of milling devices with higher energy density was not always moving in the right direction, the problem was approached in reverse, that is, how can we use mechanical activation or mechanochemical processes with minimum external stressing, by combining with other processes, among them traditional chemical or thermal processes. This is the basis of soft-mechanochemistry,42–47 where charge transfer across the grain boundary of dissimilar particles enables bridging bond formation under mild conditions. This will be explained below in more detail after we have discussed the theoretical aspects of mechanochemistry.

4.3

Theoretical aspects of modern mechanochemistry

Another aspect of mechanochemical interest is the mechanism of solid-state processes under mechanical stressing. Physical chemists tend to

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attribute a mechanochemical process to excitation of the electronic energy state to reduce the energy gap between the highest occupied molecular orbital (HOMO) and the lowest unoccupied molecular orbital (LUMO). Indeed, molecular distortion brings about energy excitation, understood as the inverse Jahn–Teller effect.48–50 This is just the reverse of the photochemical process where photo-induced excitation brings about molecular distortion. This is different from the simple route via a radical formation, which mechanochemists frequently observe.51–55 Mechanochemical processes are inevitably associated with grain downsizing, since milling processes are almost always involved. Classically, it is well known that brittle materials turn into plastic, under the concept of microplasticity.56,57 Primarily, downsizing of the solid particles becomes more difficult since the probability of pre-existing microcracks decreases with decreasing particle size. More important is the decrease in the energy efficiency for comminution with decreasing particle size. This was smartly demonstrated by Kapur et al.58 Similar phenomena must be observed more carefully when the downsizing of the solid particle proceeds into the nanometric regime. Gutkin and co-workers59,60 and Ovid’ko and Sheinerman61 suggested that the migration of the grain boundary changes significantly when a solid particle size is less than 10 nm. Inspired by stress-induced grain boundary migration, they extended their ideas to dislocation nucleation and nano-crack formation, which are attributable to the coexistence of dislocations and disclinations and their combined effects. A number of attempts have been made to explain anomalies in the nearsurface region of solids, where there are not only unusually highly concentrated lattice defects but also a loosening of chemical bonds, which are very significant in mechanochemical processes, particularly when we try to combine these concepts with nanostructured material. One of the most important things to be inferred is the charge transfer across the particulate boundary. The importance of cross boundary charge transfer cannot be overemphasized in rational solid state chemical processes, since this can be done at room temperature. Mechanical stressing is very appropriate for this purpose, as will be seen later on.

4.4

Composite oxides in electromagnetic applications

4.4.1 Principle of soft mechanochemical effects Some of the soft mechanochemical processes mentioned above were conducted in the author’s laboratory, as introduced in several review articles.46,47,62–69 Some case studies are briefly introduced below.

© Woodhead Publishing Limited, 2010

Materials design through mechanochemical processing (a)

(b) MS-0

TG

M-48

Weight loss 5%

Exothermic

Weight loss

M-1/2

M-24

TG

MS-1/12

M-0

M-8

DTA

MS-1/6 MS-1/2

Exothermic

DTA

5%

69

MS-1 MS-2 MS-3

300

400

500 800 900 Temperature (°C)

300

400

500 800 900 Temperature (°C)

4.2 Change in the thermal gravimetry (TG)–DTA profiles of Mg(OH)2 (a) and an Mg(OH)2–SiO2 mixture (b) during milling. The milling time is indicated at the end of the sample names in hours.

One of the earliest works in this field was the base–acid system, Mg(OH)2– SiO2, where dehydration and amorphization of Mg(OH)2 during milling was found to take place much faster when they coexist with SiO2 fine particles.70 The differential thermal analysis (DTA) endothermic peak caused by dehydration of Mg(OH)2 persisted even after vibro-milling Mg(OH)2 alone for 48 h, as shown in Fig. 4.2(a). The same peak entirely disappeared after milling for 3 h with SiO2 coexisting in a stoichiometric amount with magnesium monosilicate, as shown in Fig. 4.2(b). What is actually taking place at the contact points between Mg(OH)2 and SiO2 is rather straightforward. Since Mg(OH)2 is highly basic, in contrast to the highly acidic SiO2, neutralization reaction takes place, as schematically shown in Fig. 4.3. An important consequence of this surface neutralization is the formation of a bridging bond between two metallic species, M-I and M-II, Si and Mg in this particular case, respectively, abridged by an oxygen atom, as shown in Fig. 4.3(b). Note that there are almost always OH groups on the surface of metal oxide, a silanol group, Si-OH, in the case of silica. Evidence of the formation of this kind of a hetero-bridging bond (HBB), or a hetero-metalloxane bond (HMB), is given in Fig. 4.4, where the intensity of the Si–O–Si symmetric

© Woodhead Publishing Limited, 2010

High-energy ball milling (a)

(b) Acidic (hydr)oxide M-I O M-II

M-I M-II

M-1 O H

Mechanical stressing

Basic (hydr)oxide

M-1 O H2O M-2

H O M-2

4.3 Schematic illustration of HBB formation during milling of (hydr)oxide mixture with charge transfer and consequent neutralization at the contact point of dissimilar particles.

Mg(OH)2–SiO2 Si–O–Si

MS-1/6

O–H

MS-1

Transmittance

70

MS-2

3500

2500

1000 600 1600 Wavenumber (cm–1)

4.4 Change in the FT-IR spectra of the Mg(OH)2–SiO2 mixture due to milling. The milling time is indicated at the end of the sample names in hours.

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stretching vibration bond decreases with the milling time of the Mg(OH)2– SiO2 mixture. The same principle works for a number of mixed oxide systems. Comparative studies on a series of basic hydroxides combined with acidic oxides were carried out to study the effect of basicity on the trend of bridging bond formation.71,72 More detailed studies of the soft mechanochemical process were carried out for the system comprising Ca(OH)2 and SiO2.73–75 The mechanisms of HBB formation were examined by various methods, among others, in situ infrared (IR), X-ray photoelectron spectroscopy (XPS) and solid state proton nuclear magnetic resonance (NMR).73 The main outcome of these experimental studies shows the increase in the polarization of surface OH groups caused by mechanical stressing. This enables a proton formed from the surface silanol group to serve as a Brønsted acid. Thus, a liberated proton can jump into the electron rich oxygen of the OH group in Ca(OH)2. This was verified by a molecular orbital study,76 as shown in Fig. 4.5. The probability of this kind of proton jump, a typical cross-boundary charge transfer, inevitably increases under mechanochemical conditions because of the compressive stress exerted on the powder mass. It is also important to note that the consequent HBB, Si–O–Ca, is stabilized by the introduction of an oxygen vacancy around Ca, as shown in Fig. 4.6. This is again an explicit manifestation of a mechanochemical process, since introduction of an oxygen vacancy is more probable under mechanical stressing due to the microplasticity mentioned above.57 The probability of HBB formation could be generalized on the basis of the well-established concept of electronegativity equalization.77–80 In addition to the acid–base mechanisms, an alternative process via radical formation and recombination should not be overlooked.74

(a)

(b) SiO4

Si

0.2 nm

O

Z2

Ca Oxygen vacancy

H

Z1

CaOx (x = 4)

Si

Z

X

Si O

Z2Z1

O

Ca

O

Ca

Y Z2–Z1 = (i) 0.2

(ii) 0.0

(iii) 0.0

4.5 Molecular orbital calculation demonstrating the interfacial neutralization due to mechanical activation of the Ca(OH)2–SiO2 mixture. (a) Model cluster system employed; (b) detachment of H2O as a consequence of neutralization.

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High-energy ball milling 0.95

0.14

Overlap population (–)

0.12

0.90

Ca–O

0.10

0.85

0.08

0.80

0.06

0.75

0.04

0.70 O–Si

0.02 0.00

5

2 4 3 Coordination number X (–)

0.65

1

0.60

4.6 Bond overlap population between Ca and O, and O and Si in Ca(OH)2 and SiO2, respectively.

4.4.2 Application to perovskite ceramics (1): BaTiO3 The next step in materials design is combined with the subsequent heating process, called ‘calcination’ in the community of ceramics. In the author’s laboratory, a series of studies have been carried out for Mg–O–Ti or Ba–O– Ti systems, leading to MgTiO381–83 or BaTiO3,84–89 respectively. On the basis of the schematic explanation given in Fig. 4.3, it might be almost instinctively understood that a soft mechanochemical method brings about a more ready reaction just because of the preformed nuclei. As a matter of fact, the temperature of reaction completion decreased significantly.86,88 A more explicit description is given below of the effect of the mechanical activation of the BaTiO3 starting mixture. Starting powders, BaCO3 and TiO2 were mixed in the molar ratio 1 : 1 and predispersed to obtain an aqueous slurry. Samples S and W, with and without agitation milling, respectively, were obtained by freeze drying the respective slurry. Sample S was further dry vibro-milled for 5 h to obtain sample Sb.84 The corresponding increase in the homogeneity is shown in Table 4.1 by taking the variation coefficient of local composition determined by electron probe microanalysis. As seen from thermogravimetry and the corresponding differential profiles given in Fig. 4.7, the initiation and termination temperature of CO2 evolution, which tallies fairly well with the process of BaTiO3 formation,88 shift significantly toward lower temperatures with better dispersion (W to

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Table 4.1 Homogeneity of the starting mixtures for the preparation of BaTiO3

Sample

Coefficient of variation of X-ray intensity of Ti (σ/average)

W S Sb

0.324 0.074 0.063

100

TG (%)

80 60 W

40

Sb

20

S

0 W

DTG

S

Sb 200

400

600 Temperature (°C)

800

1000

4.7 TG (upper) and DTG (lower) profiles of the BaCO3 and TiO2 mixtures.

S) and subsequent vibro-milling (S to Sb). We further observed the mixture under high resolution transmission electron microscopy (HR-TEM) in an attempt to identify these particles. As shown in Fig. 4.8(a) and (b) for the sample S and Sb, individual particles above 100 nm with a high contrast are identified as BaCO3, while particles of 40–80 nm with smoother surface and lower contrast are identified as TiO2. When sample Sb is heated up to 650°C, the micrograph exhibits two categories of the particles, as shown in Fig. 4.8(c). The selected area diffraction pattern shown in Fig. 4.8(d)

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High-energy ball milling S

BaCO3

Sb

TiO2

50 nm

50 nm (a)

(b)

Sb-650°C

S-750°C

80 nm (c) (e)

(d)

40 nm

9 nm (f)

4.8 TEM micrographs of the BaCO3 and TiO2 mixtures ((a) and (b)) with those of quenched samples ((c) and (d), and (e) and (f)).

identifies the larger particle as BaCO3, as we assigned the spot pattern to belong to the space group Pnma from the fitting under the lattice parameters, a = 0.6433 nm, b = 0.5315, and c = 0.8904 nm. On the other hand, assignment of the small particle in sample S heated up to 750°C, shown in Fig. 4.8(e), was less straightforward. While the size is similar to that of TiO2, the morphology is quite different. The crystallinity of the small particle is relatively low, as shown in Fig 4.8(f). Although we failed to obtain either a clear fringe pattern or diffractograms, the small particle observed in sample S, heated up to 750°C, is highly likely to be BaTiO3. This is also supported by the X-ray diffractogram shown in Fig. 4.9, where we observed a broad diffraction peak from BT, other than those of unreacted BaCO3 and TiO2. Note that we also observed BaCO3 in sample S heated up to 750°C, or conversely, similar small BT particles in the sample Sb heated up to 650°C. Scanning electron micrographs are shown in Fig. 4.10 for the samples S and Sb after heating at 1000°C for 2 h. From the cumulative volume distribution curves obtained by the image analyses, the median diameter, D50, was significantly smaller for sample Sb with a sharper size distribution, as shown in Table 4.2. The particles of the fired product after calcining at

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Intensity (a.u.)

BT BaCO3 TiO2 anatase TiO2 rutile Sb-900°C Sb-750°C Sb-650°C S-900°C S-750°C 20

25

30 35 40 2q (CuK a) (deg)

45

S-650°C 50

4.9 XRD profiles of quenched samples. (a)

(b)

S

S

0.48 mm Sb

0.8 mm Sb

0.48 mm

0.8 mm

4.10 SEM micrographs of BaTiO3 obtained by calcining at 1000°C.

1000°C from sample S are the mixture of spherical and angled particles. In contrast, those from sample Sb are spherical with sharper size distribution. From these experiments, it is concluded that wet agitation milling deagglomerates the starting mixture and results in the clear separation of the first and second steps of BaTiO3 formation. Additional dry vibro-milling, on the other hand, disintegrates the primary particles of both ingredients significantly elevating the reactivity of the reactant mixture with the participation of the mechanochemical effects, favouring the smaller particle size with the narrower size distribution. We attribute these positive

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High-energy ball milling Table 4.2 Granulometric properties of BaTiO3 obtained from starting mixtures S and Sb

Average particle size (μm) D50% (μm) Slope: M value

S

Sb

0.12 0.19 3.8

0.20 0.25 5.0

M is defined as an average slope of the cumulative distribution curve, i.e. M = 1/(log D80 − log D20), where D80 and D20 denote the cumulative particle diameter at 80% and 20%, respectively. Larger M means sharper particle size distribution. (a)

(b)

10 μm (c)

1 μm (d)

1 μm

10 μm

4.11 Scanning electron micrographs of the reaction mixture for PMN-xPT. Left and right columns: before and after mechanical homogenization, respectively. (a) and (b): x = 0. (c) and (d): x = 0.1.

effects to efficient mechanical activation without losing the state of high dispersion.

4.4.3 Application to perovskite ceramics (2): PMN and PZN In our case study on the complex perovskite, PbMg1/3Nb2/3O3 (PMN) with and without coexistence of PbTiO3 (PMN-xPT with x = 0, 0.1 or 0.2),90 we started from a stoichiometric mixture comprising respective oxides with the exception of Mg(OH)2, which is a stable hydroxide, instead of MgO. As shown in Fig. 4.11, the degree of agglomeration of the starting mixture substantially decreases after milling with a multi-ring mill91 for 60 min. While total deagglomeration by powder processing is generally not possi-

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ble, the size distribution of the agglomerates always narrows toward homogenization in a 100 nm regime. Homogeneity in the nanometer regime was also quantified by using the standard deviation of the local atomic ratio from the overall nominal one determined by energy dispersive X-ray analysis (EDX) installed in the transmission electron microscope.90 It is, however, more important to notice that the mixture is chemically homogenized as well. As shown in Fig. 4.12, the oxygen 1s XPS spectra changes from being multimodal, reflecting the individual starting materials to that of the solid solution, that is, the end product. The increase in the relative intensity of Pb–O–multi component B-site ions by mechanical treatment is attributed to the formation of new chemical bonds as a result of chemical interactions between dissimilar components. When we calcined the mechanically homogenized mixtures mentioned in the previous section at 850°C for 4 h, we obtained always phase pure perovskite, as shown in Fig. 4.13. The morphology of the calcined powders is also much more uniform when we start from the homogenized mixture. The lattice constant decreases linearly with x, as shown in Fig. 4.14, indicating the formation of a uniform PMN-PT solid solution. As expected, the dielectric properties exhibit large differences, with typical characteristics of the relaxer in the case starting from the homogenized mixture, as shown in Fig. 4.15. The obvious superiority of the mechanically pre-

Oxygen 1s

(a)

x=0

(b)

(a)

x = 0.2

(b)

540

536 532 528 524 Binding energy (eV)

520

4.12 X-ray photoelectron spectra of the reaction mixture for PMN-xPT with x = 0 and 0.2; (a) before and (b) after milling for 60 min.

© Woodhead Publishing Limited, 2010

(111)

X=0

(110)

(100)

(311)

(400)

(311) (222) (110)

High-energy ball milling

(100)

78

X = 0.1

X = 0.2

20

30 2q (CuK a) (deg) (a)

40 20

30 2q (CuK a) (deg) (b)

40

4.13 X-ray diffraction (XRD) profiles of PMN-xPT after firing at 850°C for 4 h; (a) and (b) correspond to the starting materials without and with mechanical homogenization, respectively.

8

0.4055

6

0.4050 4 0.4045 2

0.4040 0.4035

0

0.1 X

0.2

Average grain size, dm (mm)

Lattice parameter (nm)

0.4060

0

4.14 Lattice parameter, a, and average grain size of PMN-xPT after firing at 850°C for 4 h.

treated products is primarily attributed to the elimination of the second phase, pyrochlore. We were further challenged to apply a similar technique to PZN, where Mg in PMN is substituted by Zn. Difficulty is chiefly attributed to the thermodynamic instability of PZN, so that we expect kinetic stabilization to

© Woodhead Publishing Limited, 2010

Materials design through mechanochemical processing 25000

Dielectric constant

20000

1 kHz 10 100 500 1000

79

Milled

15000

10000

5000

0 –50

Intact

0

50 Temperature (°C)

100

150

4.15 Frequency dependence of the dielectric constant of PMN-0.1PT after sintering at 1200°C for 2 h.

reach the crystalline state as rapidly as possible and to quench before decomposition or further phase transformation takes place. Since the hurdle for pure PZN was too high, we tried to prepare PZN-PMN solid solution with a minimum proportion of PMN. We first mixed PbO, Mg(OH)2, Nb2O5 and 2ZnCO3⋅3Zn(OH)2⋅H2O in the desired stoichiometry.92 The compound 2ZnCO3⋅3Zn(OH)2⋅H2O possesses higher basicity than any other commercially available Zn salts in powdery form, favouring the principle of HBB formation. A set of new XRD peaks corresponding to perovskite appear at the early stages of milling. The relative intensity of perovskite peaks increases with milling time. On subsequent heating up to 400°C and air quenching, perovskite became the sole crystalline phase. As shown in Fig. 14.16, the percentage of perovskite decreases and again increases after showing a minimum at temperatures between 700°C and 800°C for the mixtures up to x = 0.8. We attributed the increase in the perovskite to the incorporation of PMN into PZN due to restabilization of the perovskite phase. No recovery of perovskite is observed for components with x = 0.9 and x = 1.0 (pure PZN). It has been theoretically calculated that substitution of Pb(II) at the A-site of the perovskite by Ba up to 10 mol% or Zn at the B-site by Mg up to 7.4 mol% will stabilize PZN. In actual fact, however, an additive concentration as high as 40 mol% Mg to PZN (PZxM1−xN for x = 0.6) was necessary for phase pure perovskite. The instability of perovskite PZxM1−xN is thus much higher than theoretically expected. Rapid formation of a

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80

High-energy ball milling 100

% Perovskite

80

60

40

20

0

400

500

600 700 800 Temperature (°C)

900

4.16 Plots of percentage perovskite against firing temperature for mixtures milled for 180 min. Different curves corresponding to different x values are shown from top to bottom: 0 (PMN), 0.2, 0.4, 0.6, 0.7, 0.8, 0.9 and 1.0 (PZN).

PZxM1-xN solid solution at relatively low temperatures, where the lattice strain is still accommodated in the perovskite structure, is the essence of a successful synthesis. A continuous increase in the X-ray diffraction peak breadth of perovskite was observed with the increasing fraction of PZN, x, as shown in Fig. 4.17.

4.5

Organic synthesis and utilization of spontaneous chemical reactions

4.5.1 Diels–Alder addition reaction The Diels–Alder (DA) reaction counts as one of the key steps in the total organic syntheses of complicated natural compounds and hence is examined in depth, together with the mechanisms of the syntheses.93–95 The reaction involves diene and dienophile, both having C=C bonds which form adducts with a six-membered ring to form two new C–C bonds. It is generally accepted that a smaller HOMO–LUMO gap (HLG) between the HOMO of the diene and the LUMO of the dienophile favours the reaction. Indeed, introduction of electron-accepting groups to the dieno-

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Increase in FWHM (deg)

0.12 (110) (200) (211)

0.1 0.08 0.06 0.04 0.02 0

0

0.2

0.4

0.6

0.8

1

x

4.17 Increase in the full-width at half-maximum (FWHM) of XRD peaks with fraction of PZN, x, for mixtures milled for 180 min and fired at 1100°C for 1 h.

Table 4.3 Yields of Diels–Alder reaction under various conditions Entry

R1

R2

Condition

Time (h)

Yield (%)c

1 2 3 4 5 6 7

Br CH2Cl H Me Me Me Me

Br CH2Cl H H Me Me Me

Agate 2 mm ampa Agate 2 mm ampa Agate 2 mm ampa Agate 2 mm ampa Agate 2 mm ampa Mixingb Chloroform solution

5 5 5 5 5 48 5

0 0 0 8 26 Trace 40

a

Reactions were carried out under milling with Pulverisette 0. Mixing two powdered compounds in a rotary evaporator. c Yield determined by 1H-NMR (chloroform-d solution). b

phile and/or electron-donating groups to the diene enhances the DA reaction. The DA reaction can also take place in the solid state.96–98 In the latter case, we may enhance the reactivity via a mechanochemical route, that is by introducing molecular strain on the molecular crystals.99,100 This kind of solvent-free synthesis is increasingly explored in pursuit of green chemical concepts. Table 4.3 exhibits the yields of the mechanochemical DA reaction between anthracene (AN) derivatives and benzoquinone (BQ), determined by 1H-NMR. For the reaction we used a small vibro-mill.99 It is noteworthy that the yield was increased by introducing an electron-donating group to AN, in accordance with the principle of decreasing HLG mentioned above, even for the present mechanochemical processes, which in most cases are

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High-energy ball milling O

Yield (%)

Me

100 90 80 70 60 50 40 30 20 10 0

+ O

Me

O Mechanical activation

Me

O R

With Binol, mechanochemical With Binol in chloroform at RT Chloroform solution at RT without additives Without additives

0

1

2 3 Reaction time (h)

4

5

4.18 Effect of alcohols on the rate of Diels–Alder reaction (in stoichiometric amounts). RT, room temperature.

non-Arrhenian.10 No appreciable reaction was observed in the physical mixture without milling. Poor mixing and hence a smaller number of contact points could also count for this, together with the HLG issue mentioned above. Upon addition of 1,1-bi-2-naphthol (Binol, BN) to the reaction system, dimethylanthracene (DMA) and BQ, the rate of the DA reaction increased remarkably toward completion, as shown in Fig. 4.18. It is known that a ladder-like molecular complex is formed between racemic BN and BQ. A more complicated complex is also reported to form, via a solid-state route, between BN, BQ and AN with the aid of mechanical stressing.101 Formation of this kind of charge transfer complex (CTC) serves as strong incentive for rational mechanochemical processes. To exploit the utilization of these charge transfer complexes further, we noted autogenous eutectic formation with automatic fusion during CTC formation between BQ and thymol (TM). As we conducted a mechanochemical DA reaction between DMA and BQ in the presence of the BQ-TM CTC, further acceleration of the mechanochemical DA reaction was confirmed.100 Immediately after adding TM to BQ, autogenous fusion was observed. When a small bit of crystalline DMA was attached to this mixture, it automatically fused and formed a CTC. The color of the entire mixture turned into that of the DA adducts and solidified again after 30 min, where the yield was 95%. Reaction kinetic curves for these CTC-aided mechanochemical DA reactions are given in Fig. 4.19. It should also to be mentioned that the intensity of mechanical stress needed for this kind of paste-like reaction mixture is much lower than those for dry powdered systems.

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Materials design through mechanochemical processing

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100 90 80 Yield (%)

70 60

Additives 200 mol% TM 10 mol% TM 5 mol% Binol None

50 40 30 20 10 0

0

0.5

1 Reaction time (h)

1.5

2

4.19 Diels–Alder reaction with a eutectic complex. TM, thymol.

4.5.2 Correlation between molecular distortion and change in the frontier orbitals It is well known that the symmetry of crystals and molecules is attributed basically to the symmetry of the electronic wave function, or more straightforwardly, to the spacial distribution of the electronic crowds. These symmetries are very specific to particular substances and those in gas or liquid phase cannot be altered in the usual manner. In the molecular solid state, however, the distorted state of molecules is quenchable through the plasticity of these crystals. In the case of transition metal complexes, a change in the electron orbitals upon mechanical stressing is observed by different analytical methods, for example, Mössbauer or XPS spectra.102–104 In addition, we also observed that by starting from the distorted crystals, the rate of complex formation increases. Furthermore, a mechanochemical process can produce a mono-coordinated Fe(II) complex with 1,10-phenanthroline (phen), when we mill Fe(II) oxalate with phen.105 The usual thermal process yields only bi-coordinates. Excitation of the electron energy states by mechanical activation can be explained by an inverse Jahn–Teller effect,106 also referred at the beginning of the Section 4.3. An explosive with a highly symmetrical molecule, hexahydro-1,3,5-trinitro 1,3,5-triadine, RDX, is analysed by initiating its fast reaction, with the loss of its molecular symmetry, by a shock wave,107 to explain the mechanisms of its mechanical detonation rationally. A similar discussion can take place for the solid state DA reactions mentioned previously. When two reacting molecules, AN and BQ, approach each other, the coplanar AN molecule folds along the symmetry axis C10– C11 so that the dihedral angle decreases, as shown in Fig. 4.20.108 The total energy of the system, as simultaneously plotted in Fig. 4.20 exhibits a

© Woodhead Publishing Limited, 2010

1000

190

900

180

800

170

700

160

600 500

150 r

140

400 300

130

θ

200

120

100

110

0

Dihedral angle, θ (deg)

High-energy ball milling

Energy of system (kcal mol–1)

84

100 Intermolecular distance, r (Å)

4.20 Change in the energy of the system and dihedral angle of anthracene with decreasing intermolecular distance calculated using the AM1 method. The dihedral angle of anthracene at the state of activated complex is 150°. Note that the results given here refer to a model comprising two free molecules instead of those in an actual solid state reaction, where molecules are bound to the solid surface.

maximum, corresponding to the formation of the activated complex, at a dihedral angle of approximately 150°. When intermolecular HLG values for four derivatives of AN and BQ are plotted against the dihedral angle of the former, we recognize the decrease in HLG with increasing extent of folding of AN derivatives, as shown in Fig. 4.21.108 Although the computation is semi-empirical, using SPALTAN ’02, with relatively loose rigorousness, the results shown in Fig. 4.21 could be correct as a general trend. The reliability of the computation is evident for the point where the HLG takes the smallest value for DMA (curve c) compared at the same dihedral angle. It is thus verified that the mechanically derived molecular strain increases the reactivity of solids. Furthermore, charge transfer, often occurring autogenously, brings about a molecular complex or CTC. By utilizing these phenomena, mechanochemical processes can be rationalized and their application can also be extended deeper into the area of chemical syntheses of organic adducts and complexes with or without central metallic species. These processes are compatible with the world trend toward using green chemical processes.

4.6

Conclusions and future trends

An overview is given above of some possible applications of mechanochemical principles in materials design with technical affordability.

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Materials design through mechanochemical processing

HOMO–LUMO gap (eV)

6.6

85

d

6.5

e

6.4 a

6.3 6.2

b

6.1

c

6 145

150

155 160 170 165 Dihedral angle, θ (deg)

175

180

4.21 The intermolecular HOMO–LUMO gap of (a) anthracene, (b) 9-methyl anthracene, (c) 9,10-dimethyl anthracene, (d) 9,10-bis(chloromethyl)anthracene, and (e) 9,10-bis(bromomethyl) anthracene.

Mechanical stressing on the large scale is quite costly and contamination by the equipment is unavoidable. The process is often stochastic and localized. Care should, therefore, be taken not try to apply a mechanochemical process unless there is an absolute merit and necessity. Charge transfer across a boundary of dissimilar solids with simultaneous polarization, size reduction and homogenization without external heating are the main merits of applying mechanochemistry to various functional materials including metal coordination compounds or molecular complexes via smart or even autogenous charge transfer. This could lead to innovation in organic syntheses starting from the solid state, involving solvent-free and hence green chemical processes. Since mechanical stress often brings about unusually excited states, it could enable some noble metals that have been unavoidable in catalysis to be replaced with less expensive metals. Mechanochemical processes or mechanically aided charge transfer is expected to be applied further to biocompatible complexes109,110 and pharmaceutics.111,112

4.7

References

1. ping w., ‘Digital materials design: computational methodologies as a discovery tool’, Mater Res Bull, 2006, 31, 995–8. 2. sun w., lin f. and hu x., ‘Computer-aided design and modelling of composite unit cells’, Composites Sci Technol, 2001, 61, 289–99. 3. tomonori i., ‘Recent progress in computer-aided materials design for compound semiconductors’, J Appl Phys, 1995, 77, 4845. 4. hagrman p. j., finn r. c. and zubieta j., ‘Molecular manipulation of solid state structure: influences of organic components on vanadium oxide architectures’, Solid State Sci, 2001, 3, 745–74.

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46. senna m., ‘A straight way toward phase pure complex oxides’, J Eur Ceram Soc, 2005, 25, 1977–84. 47. senna m., kinoshita t., abe y., kishi h., ando c., doshida y. and stojanovic b., ‘Smart soft-mechanochemical syntheses of well-crystallized pure phase fine particulates of mixed oxides for electroceramics’, J Eur Ceram Soc, 2007, 7, 4301–6. 48. luty t., ordon p. and eckhardt c. j., ‘A model for mechanochemical transformations: Applications to molecular hardness, instabilities, and shock initiation of reaction’, J Chem Phys, 2002, 117, 1775–85. 49. watanabe h., matsui e., ishiyama y. and senna m., ‘Solvent free mechanochemical oxygenation of fullerene under oxygen atmosphere’, Tetrahedron Lett, 2007, 48, 8132–7. 50. fishman r. s., okamoto s. and reboredo f. a., ‘Inverse Jahn-Teller transition in bimetallic oxalates’, Phys Rev Lett, 2008, 101, 116402/1–4. 51. potter w. d. and scott g., ‘Mechanochemical reactions of polymers. Formation of free radicals in stressed rubber’, Eur Polym J, 1971, 7, 489–97. 52. chemerisov s. d., grinberg o. ya., tipikin d. s., lebedev ya. s., kurreck h. and moebius k., ‘Mechanically induced radical pair formation in porphyrin-quinone and related donor-acceptor mixtures. Unusual stability and zero-field splittings’, Chem Phys Lett, 1994, 218(4), 353–61. 53. chemerisov s. d., grinberg o. y., tipikin d. s., lebedev y. s., kurreck h. and moebius k., ‘Mechanically induced radical pair formation in porphyrinquinone and related donor-acceptor mixtures. Unusual stability and zerofield splittings’, Chem Phys Lett, 1994, 218, 353–61. 54. tanaka y., zhang q., saito f., ikoma t. and tero-kubota s., ‘Dependence of mechanochemically induced decomposition of mono-chlorobiphenyl on the occurrence of radicals’, Chemosphere, 2005, 60(7), 939–43. 55. billik p., plesch g., brezova v., kuchta l., valko m. and mazur m., ‘Anatase TiO2 nanocrystals prepared by mechanochemical synthesis and their photochemical activity studied by EPR spectroscopy’, J Phys Chem Solids, 2007, 68, 1112–16. 56. beke b., The Process of Fine Grinding, Dr. W. Junk Publishing, Budapest, 1981. 57. smekal a. g., ‘The microplasticity of hard substances’, Plansee Proc, 1956, 1955, 28–40. 58. kapur p. c., schönert k. and fuerstenau d. w., ‘Energy-size relationship for breakage of single particles in a rigidly mounted roll mill’, Int J Mineral Processing, 1990, 29, 221–33. 59. gutkin m. yu. and ovid’ko i. a., ‘Homogeneous nucleation of dislocation loops in nanocrystalline metals and ceramics’, Acta Materialia, 2008, 56, 1642–9. 60. gutkin m. yu., mikaelyan k. n. and ovid’ko i. a., ‘Athermal grain growth through cooperative migration of grain boundaries in deformed nanomaterials’, Scripta Materialia, 2008, 58, 850–3. 61. ovid’ko i. a. and sheinerman a. g., ‘Nanocrack generation at dislocationdisclination configurations in nanocrystalline metals and ceramics’, Phys Rev B, 2008, 77, 054109-1–7. 62. senna m., ‘Incipient chemical interaction between fine particles under mechanical stress – a feasibility of producing advanced materials via mechanochemical routes’, Solid State Ionics, 1993, 63–5, 3–9.

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63. senna m., ‘More chemistry for finer particle technology’, Trans Inst Chem Eng (Part A), 1998, 76, 767–74. 64. senna m., ‘Chemical powder technology- a new insight into atomic processes on the surface of fine particles’, Adv Powder Technol, 2002, 13(2), 115– 38. 65. senna m., ‘Consequences of molecular strain on the solid state addition reaction’, J Mater Sci, 2004, 39, 4995–5001. 66. senna m., ‘Preparation of functional materials via non-conventional routes’, Ann Chem Sci Mater, 2002, 27(6), 3–14. 67. senna m., ‘Boundary-free nanocomposites – Design, preparation and charactrization’, Mater Sci Eng A, 2005, A412, 37–42. 68. senna m., ‘Smart solid state syntheses of well-crystallized phase pure mixed oxides for electroceramics’, J Korean Ceram Soc, 2006, 43(11), 680–7. 69. senna m., ‘Rational mechanochemical processes with less intensive stressing for their affordable application’, J Thermal Anal, 2007, 90(1), 107–13. 70. liao j. f. and senna m., ‘Enhanced dehydration and amorphization of Mg(OH)2 in the presence of ultrafine SiO2 under mechanochemical conditions’, Thermochim Acta, 1992, 210, 89–102. 71. liao j. f. and senna m., ‘Mechanochemical dehydration and amorphization of hydroxides of Ca, Mg and Al on grinding with and without SiO2’, Solid State Ionics, 1993, 66, 313–19. 72. watanabe t., liao j. f. and senna m., ‘Changes in the basicity and species on the surface of Me(OH)2-SiO2 (Me = Ca, Mg, Sr) mixtures due to mechanical activation’, J Solid State Chem, 1995, 115(2), 390–4. 73. watanabe t., isobe t. and senna m., ‘Mechanisms of incipient chemical reaction between Ca(OH)2 and SiO2 under moderate mechanical stressing, I. A solid state acid–base reaction and charge transfer due to complex formation’, J Solid State Chem, 1996, 122, 74–80. 74. watanabe t., isobe t. and senna m., ‘Mechanisms of incipient chemical reaction between Ca(OH)2 and SiO2 under moderate mechanical stressing, II. Examination of a radical mechanism by an EPR study’, J Solid State Chem, 1996, 122, 291–6. 75. watanabe t., isobe t. and senna m., ‘Mechanisms of incipient chemical reaction between Ca(OH)2 and SiO2 under moderate mechanical stressing, III. Changes in the short range ordering through the mechanical and thermal processes’, J Solid State Chem, 1997, 130, 284–9. 76. fujiwara y., isobe t., senna m. and tanaka j., ‘Effects of reduced coordination number for Ca on the electron redistribution during Ca–O–Si bridge bonding from CaO or Ca(OH)2 and SiO2’, J Phys Chem, 1999, 103, 9842–6. 77. sanderson r. t., ‘An interpretation of bond lengths and a classification of bonds’, Science, 1951, 114, 670–2. 78. sanderson r. t., ‘Electronegativities in inorganic chemistry’, J Chem Educ, 1952, 29, 539–44. 79. sanderson r. t., ‘Electronegativity and bond energy’, J Am Chem Soc, 1983, 105(8), 2259–61. 80. watanabe t., isobe t. and senna m., ‘Electronegativity equalization during mechanochemical addition reaction’, Chem Sustainable Dev, 1998, 6(2–3), 165–70.

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81. hamada k. and senna m., ‘Mechanochemical effects on the properties of starting mixtures for PbTiO3 ceramics by using a novel grinding equipment’, J Mater Sci, 1996, 31, 1725–28. 82. baek j. g., isobe t. and senna m., ‘Mechanochemical effects on the precursor formation and microwave dialectic characteristics of MgTiO3’, Solid State Ionics, 1996, 90, 269–79. 83. hamada k., yamamoto s., nagao m. and senna m., ‘Measurement of compressive and shear forces in multi-ring media mill’, J Chem Eng Jpn, 1997, 30(4), 756–9. 84. ando c., yanagawa r., chazono h., kishi h. and senna m., ‘Nuclei-growth optimization for fine-grained BaTiO3 by precision-controlled mechanical pretreatment of starting powder mixture’, J Mater Res, 2004, 19(12), 3592–9. 85. oguchi k., ando c., chazono h., kish h. and senna m., ‘Effects of glycine on the solid-state synthesis of barium titanate micro-particles with high tetragonality’, J Phys IV France, 2005, 128, 33–9. 86. yanagawa r., ando c., chazono h., kishi h. and senna m., ‘Well-crystallized tetragonal BaTiO3 micro particles via a solid-state reaction preceded by agglomeration-free mechanical activation’, J Am Ceram Soc, 2007, 90, 809–14. 87. ando c., kshi h., oguchi h. and senna m., ‘Effects of bovine serum albumin on the low temperature synthesis of barium titanate microparticles via a solid state route’, J Am Ceram Soc, 2006, 89(5), 1709–12. 88. ando c., suzuki t., mizuno y., kishi h., nakayama s. and senna m., ‘Evaluation of additive effects and homogeneity of the starting mixture on the nucleigrowth processes of barium titanate via a solid state route’, J Mater Sci, 2008, 43, 6182–92. 89. ando c., tsuzuku k., kobayashi t., kishi h., kuroda s. and senna m., ‘Function of glycine during solid-state reaction toward well-crystallized fine particulate barium titanate’, J Mater Sci: Materials in Electronics, 2009, 20, 844–50. 90. beak j. g., isobe t. and senna m., ‘A new fabrication technique for perovskite 0.9Pb(Mg1/3 Nb2/3)O3. 0.1TiO2 ceramics via a soft-mechanochemical route’, J Am Ceram Soc, 1997, 80, 973–81. 91. hamada k., yamamoto s., nagano m. and senna m., ‘Measurement of compressive and shear forces in multi-ring media mill’, J Chem Eng Jpn, 1997, 30, 756–9. 92. shinohara s., isobe t. and senna m., ‘Synthesis of phase pure Pb(ZnxMg1−x)Nb2/3 O3 up to x = 0.7 from a single mixture via a soft-mechanochemical route’, J Am Ceram Soc, 2000, 83, 3208–10. 93. ess d. h., jones g. o. and houk k. n., ‘Conceptual, qualitative, and quantitative theories of 1,3-dipolar and Diels–Alder cycloadditions used in synthesis’, Adv Synth Catalysis, 2006, 348(16 + 17), 2337–61. 94. atherton j. c. c. and jones s., ‘Diels–Alder reactions of anthracene, 9-substituted anthracenes and 9,10-disubstituted anthracenes’, Tetrahedron, 2003, 59(46), 9039–57. 95. dennis n., ‘Addition reactions: cycloaddition’, Organic Reaction Mechanisms, 2003, 1998, 453–85. 96. kiselev v. d., iskhakova g. g., kashaeva e. a., potapova l. n. and konovalov a. i., ‘Diels–Alder reaction volumes in the solid state and solution’, Russ Chem Bull, 2004, 53(11), 2490–5. 97. radha kishan k. v. and desiraju g. r., ‘Crystal engineering: a solid state Diels– Alder reaction’, J Org Chem, 1987, 52(20), 4640–1.

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98. kim j. h., hubig s. m., lindeman s. v. and kochi j. k, ‘Diels–Alder topochemistry via charge-transfer crystals: novel (thermal) single-crystal-to-single-crystal transformations’, J Am Chem Soc, 2001, 123(1), 87–95. 99. watanabe h. and senna m., ‘Acceleration of solid state Diels–Alder reactions by incorporating the reactants into crystalline charge transfer complexes’, Tetrahedron Lett, 2005, 46, 6815–18. 100. watanabe h. and senna m., ‘A Diels–Alder reaction catalyzed by eutectic complexes autogenously formed from solid state phenols and quinones’, Tetrahedron Lett, 2006, 47, 4481–4. 101. cheung e. y., kitchin s. j., harris k. d. m., imai y., tajima n. and kuroda r., ‘Direct structure determination of a multicomponent molecular crystal prepared by a solid-state grinding procedure’, J Am Chem Soc, 2003, 125, 14658–9. 102. tsuchiya n., tsukamoto a., ohshita t., isobe t., senna m., yoshioka n. and inoue h., ‘Anomalous spin crossover of mechanically strained iron(II) complexes with 1,10-phenantroline with their counterions, NCS- and PF6-’, J Solid State Chem, 2000, 153, 82–91. 103. tsuchiya n., tsukamoto a., ohshita t., isobe t., senna m., yoshioka n. and inoue h., ‘Stress-induced ligand field distribution and consequent multi-mode spin crossover in FeII(phen)2(NCS)2 and FeII[HB(pz)3]2’, Solid State Sci, 2001, 3, 705–14. 104. ohshita t., tsukamoto a. and senna m., ‘Change in the magnetic properties of [FeII(phen)3](PF6)2 in the solid state by combining grinding and annealing’, Physica Status Solidi (a), 2004, 201, 762–71. 105. ohshita t., hisore k., komai m. and senna m., ‘Synthesis of mono-coordinate iron (II)–phen complex via a solid-state ligand exchange process from iron (ii) oxalate dihydrate at room temperature under mechanical stressing’, Synth Reactivity Inorg Metal-Organic Nano-Metal Chem, 2005, 25, 355–8. 106. kiefer j. k., tranter r. s., wang h. and wagner a. f., ‘Thermodynamic functions for the cyclopentadienyl radical: The effect of Jahn–Teller distortion’, Int J Chem Kinetics, 2001, 33, 834–45. 107. luty t., ordon p. and eckhardt c. j., ‘A model for mechanochemical transformations: Applications to molecular hardness, instabilities, and shock initiation of reaction’, J Chem Phys, 2002, 117, 1775–85. 108. pradipta m. f., watanabe h. and senna m., ‘Semiempirical computation of the solid phase Diels–Alder reaction between anthracene derivatives and p-benzoquinone via molecular distortion’, Solid State Ionics, 2004, 172, 169–72. 109. wang l., nemoto r. and senna m., ‘Microstructure and chemical states of hydroxyapatite/silk fibroin nanocomposites synthesized via a wet-mechanochemical route’, J Nanoparticle Res, 2002, 4, 535–40. 110. nemoto r., wang l., aoshima m., senna m., ikoma t. and tanaka j., ‘Increasing the crystallinity of hydroxyapatite nanoparticles in composites containing bioaffinitive organic polymers by mechanical stressing’, J Am Ceram Soc, 2004, 87(6), 1014–17. 111. senna m. and nakayama s., ‘Preparation and properties of nano-amorphous organic and inorganic particles via chemical and mechanochemical routes’, J Alloys Compds 2009, 483, 265–70. 112. nakayama s., watanabe t. and senna m., ‘Rapid amorphization of molecular crystals by sorption of solvent molecules in the presence of hydrophilic matrice’, J Alloys Compd, 2009, 483, 217–21.

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5 Kinetic processes and mechanisms of mechanical alloying F. D E L O G U, Università degli Studi di Cagliari, Italy; and G. M U L A S, Università degli Studi di Sassari, Italy

Abstract: This chapter focuses on the study of the kinetics of mechanically activated alloying processes. This is a difficult area of study that has great importance for any substantial progress in the field from both the fundamental and applied points of view. Various questions arise when a rationalization of experimental evidence is attempted. These concern aspects as different as the dynamics of the apparatuses used for mechanical activation on the macroscopic level and the intimate features of individual reactive events induced by mechanical stresses on the microscopic level. A brief overview of such intertwined questions is given together with a possible line of approach to the characterization of mechanical alloying mechanisms operating on the atomic scale (atomistic). Key words: mechanical alloying, phase transformation, kinetics, atomistic mechanisms.

5.1

Introduction

Mechanical alloying (MA) is a powder metallurgy methodology widely employed to synthesize metallic alloys starting from pure elements (Heinicke, 1984; Butyagin, 1989; Suryanarayana, 2001). MA is usually performed in suitably designed ball mills, inside which milling tools undergo a series of collisions with each other (Heinicke, 1984; Butyagin, 1989; Suryanarayana, 2001). MA is generally thought to proceed through the accumulation of structural defects in crystalline lattices and the intimate mixing of elemental species at interfaces consequent on the mechanical deformation of powder particles trapped between the surfaces of colliding milling tools (Suryanarayana, 2001). However, any agreement is lost as soon as attention is focused on the mechanisms that may underlie the MA process on the atomic scale. Hypotheses regarding the mechanisms range from the oldest scenarios of hot spot theory (Benjamin, 1992; Koch, 1992) and defect-enhanced diffusion (Schwarz and Johnson, 1983; Johnson, 1986; Hellstern and Schultz, 1986; Cahn and Greer, 1996) to the most recent ones of shear-induced 92

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mixing, interface roughening and contact melting (Martin, 1984; Bellon and Averback, 1995; Hammerberg et al., 1998; Urakaev and Boldyrev, 2000; Vick et al., 2000; Fu et al., 2001; Lund and Schuh, 2003; Delogu and Cocco, 2005a,b; Odunuga et al., 2005). While passing from the former approach to the latter ones, the emphasis gradually shifts from inherently thermal processes to coupled thermally and mechanically induced phenomena. Along these lines, MA reactions can be tentatively connected to local rearrangements at sliding interfaces that originate from the action of non-hydrostatic mechanical stresses operating on the atomic scale (Levitas, 2004; Martin, 1984; Bellon and Averback, 1995; Hammerberg et al., 1998; Urakaev and Boldyrev, 2000; Vick et al., 2000; Fu et al., 2001; Lund and Schuh, 2003; Delogu and Cocco, 2005a,b; Odunuga et al., 2005). In fact, relative sliding events are accompanied by the local generation of frictional heat, which can easily result in far-from-equilibrium conditions promoting a transient enhancement of chemical reactivity (Martin, 1984; Bellon and Averback, 1995; Hammerberg et al., 1998; Urakaev and Boldyrev, 2000; Vick et al., 2000; Fu et al., 2001; Lund and Schuh, 2003; Delogu and Cocco, 2005a,b; Odunuga et al., 2005). Owing to the lack of direct experimental evidence on the microscopic MA mechanisms, the above-mentioned scenarios are substantially based on conjecture or numerical findings. Of course, this is not the best way to pursue the research farther. Yet, valuable insight into the mechanisms of MA processes in ball mills can be still gained by investigating the apparent kinetics of phase transformations and relating the macroscopic evidence obtained to individual local deformation events. This chapter describes an attempt in this direction.

5.2

Fundamentals of mechanical alloying processes in ball mills

A ball mill is a relatively simple apparatus in which the motion of the reactor, or of a part of it, induces a series of collisions of balls with each other and with the reactor walls (Suryanarayana, 2001). At each collision, a fraction of the powder inside the reactor is trapped between the colliding surfaces of the milling tools and submitted to a mechanical load at relatively high strain rates (Suryanarayana, 2001). This load generates a local nonhydrostatic mechanical stress at every point of contact between any pair of powder particles. The specific features of the deformation processes induced by these stresses depend on the intensity of the mechanical stresses themselves, on the details of the powder particle arrangement, that is on the topology of the contact network, and on the physical and chemical properties of powders (Martin et al., 2003; Delogu, 2008a). At the end of any given collision event, the powder that has been trapped is remixed with the

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powder that has not undergone this process. Correspondingly, at any instant in the mechanical processing, the whole powder charge includes fractions of powder that have undergone a different number of collisions. The individual reactive processes at the perturbed interface between metallic elements are expected to occur on timescales that are, at most, comparable with the collision duration (Hammerberg et al., 1998; Urakaev and Boldyrev, 2000; Lund and Schuh, 2003; Delogu and Cocco, 2005a,b). Therefore, unless the ball mill is characterized by unusually high rates of powder mixing and frequency of collisions, reactive events initiated by local deformation processes at a given collision are not affected by a successive collision. Indeed, the time interval between successive collisions is significantly longer than the time period required by local structural perturbations for full relaxation (Hammerberg et al., 1998; Urakaev and Boldyrev, 2000; Lund and Schuh, 2003; Delogu and Cocco, 2005a,b). These few considerations suffice to point out the two fundamental features of powder processing by ball milling, which in turn govern the MA processes in ball mills. First, mechanical processing by ball milling is a discrete processing method. Second, it has statistical character. All of this has important consequences for the study of the kinetics of MA processes. The fact that local deformation events are connected to individual collisions suggests that absolute time is not an appropriate reference quantity to describe mechanically induced phase transformations. Such a description should rather be made as a function of the number of collisions (Delogu et al., 2004). A satisfactory description of the MA kinetics must also account for the intrinsic statistical character of powder processing by ball milling. The amount of powder trapped in any given collision, at the end of collision is indeed substantially remixed with the other powder in the reactor. It follows that the same amount, or a fraction of it, could at least in principle be trapped again in the successive collision. This is undoubtedly a difficult aspect to take into account in a mathematical description of MA kinetics. There are at least two extreme cases to consider. On the one hand, it could be assumed that the powder trapped in a given collision cannot be trapped in the successive one. On the other, it could be assumed that powder mixing is ideal and that the amount of powder trapped at a given collision has the same probability of being processed in the successive collision. Both these cases allow the development of a mathematical model able to describe the relationship between apparent kinetics and individual collision events. However, the latter assumption seems to be more reliable than the former one, at least for commercial mills characterized by relatively complex displacement in the reactor (Manai et al., 2001, 2004). A further obvious condition for the successful development of a mathematical description of MA processes is the one related to the uniformity

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of collision regimes. More specifically, it is highly desirable that the powders trapped at impact always experience the same conditions. This requires the control of the ball dynamics inside the reactor, which can be approximately obtained by using a single milling ball and an amount of powder large enough to assure inelastic impact conditions (Manai et al., 2001, 2004; Delogu et al., 2004). In fact, the use of a single milling ball avoids impacts between balls, which have a remarkable disordering effect on the ball dynamics, whereas inelastic impact conditions permit the establishment of regular and periodic ball dynamics (Manai et al., 2001, 2004; Delogu et al., 2004). All of the above assumptions and observations represent the basis and guidelines for the development of the mathematical model briefly outlined in the following. It has been successfully applied to the case of a Spex Mixer/ Mill mod. 8000, but the same approach can, in principle, be used for other ball mills.

5.3

A phenomenological model of mechanical alloying kinetics

During MA reactions by ball milling, powder particles undergo fracturing and coalescence processes so that their number changes (Heinicke, 1984; Butyagin, 1989; Suryanarayana, 2001). At the same time, complex morphological, microstructural and chemical transformations that affect each other take place (Heinicke, 1984; Butyagin, 1989; Suryanarayana, 2001). In contrast, the volume of powder trapped at each collision keeps approximately constant (Manai et al., 2004; Delogu et al., 2004). Indeed, the amount of powder trapped between the surfaces of colliding milling tools depends on the total amount of powder within the vial, on the ball size and on the relative impact velocity (Courtney, 1995; Maurice and Courtney, 1996). Therefore, the volume of powder trapped at individual collisions can be regarded as a key reference quantity in describing MA kinetics. It should be noted regarding this point that the mechanical load at collision is not uniformly distributed within this volume. On the contrary, as briefly mentioned before, it is distributed in the network formed by the points of contact between different particles (Martin et al., 2003; Delogu, 2008a). The extent to which chemical modifications occur depend on the intensity of the non-hydrostatic mechanical stresses generated at these points. Therefore, the rate of MA processes is related to the volume fraction of powder trapped at collision during which a mixing of elemental species actually takes place. Of course, whether or not such mixing could take place depends on whether the mechanical stresses in the volume of trapped powders reach the necessary intensity.

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Numerous questions arise in connection with the aforementioned observations. For example, how large must be the fraction of powder trapped in a given individual collision, if that collision is involved in deformation events severe enough to start the interfacial mixing of elemental species? How many individual collisions are necessary for this process to proceed? As the number of loading events experienced by a given quantity of powder is different from the total number of collisions, what is the relationship between the number of loading events undergone by a certain quantity of powder and the total number of collisions? As will be shown below, an approximate response to such questions can be given, provided that a few critical assumptions are made. Let us assume that the volume vimp of powders trapped at each individual collision remains constant during the course of the MA process. For the sake of simplicity, the average density ρ of the powders is also considered constant. It follows that the mass mimp of powder trapped on the average at each collision is equal to vimpρ. Therefore, the volume fraction ximp of powder trapped at individual collisions is equal to vimpv−1 p , where vp is the total powder volume mpρ−1 and mp is the total powder mass inside the reactor. The rate at which a given MA reaction proceeds will be necessarily proportional to the fraction ximp. Let us also assume that mixing chemical species at perturbed interfaces will occur only in the fraction of powder that has experienced critical loading conditions (CLCs) a number I of times. CLCs can be defined as the conditions under which a certain amount of powder experiences a mechanical stress σ higher than a critical threshold value σ0. The critical threshold value σ0 which allows a given MA process to occur is in turn defined as the lowest mechanical stress necessary to induce the mixing of elemental species at perturbed interfaces. A further necessary assumption is that at each individual collision, a fraction k of the powder charge mp in which σ is higher than σ0 can be found. Of course, k must be smaller than ximp. Finally, let us assume that the powders are perfectly remixed after each collision event, so that the powder mass inside the reactor is kept homogeneous and that each volume fraction k in which the powder mass mp can be ideally subdivided has the same probability of being involved in a collision event as any other. Let us define now χ0 as the volume fraction of powder charge that never experienced CLCs. Of course, χ0 is equal to 1 before the first collision has occurred. After it has taken place, a volume fraction χ1 of the powder trapped at collision has been submitted for the first time to CLCs. According to the above mentioned assumptions, χ1 corresponds to k. It follows that the fraction of powder charge never submitted to CLCs after the first collision has occurred is equal to χ0(1) = 1 − k.

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After the second collision has taken place, the fraction of powder never submitted to CLCs becomes equal to χ0(2) = 1 − k − k(1 − k) = (1 − k)2. The fraction of powder submitted to CLCs once is equal to χ1(2) = k + k(1 − k) − k2 = 2k(1 − k) and the fraction of powder submitted to CLCs for two times is equal to χ2(2) = k2. The process briefly sketched above continues as the number n of collisions increases. The fraction of powder χ0(n) never submitted to CLCs after n impacts is equal to: χ0(n + 1) = χ0(n) − kχ0(n)

[5.1]

The fraction of powder submitted to i CLCs after n impacts is instead equal to: χi(n + 1) = χi(n) − kχi(n) + kχi−1(n)

[5.2]

In fact, at the nth collision, a fraction kχi(n) of powder experiences CLCs for the (i + 1)th time, thus becoming part of the fraction χi+1(n). Conversely, a fraction kχi−1(n) experiences CLCs for the ith time and becomes part of the fraction χi(n). It is now worth noticing that usually the fraction k of powder that is submitted to CLCs at each collision is small with respect to the total amount of powder (Benjamin, 1992; Delogu et al., 2004; Delogu and Cocco, 2006, 2007). Under these conditions, Equations [5.1] and [5.2] can be written in the continuous form shown below: dχ0(n) = −kχ0(n)dn

[5.3]

dχ0(n) = −kχi(n)dn + kχi−1(n)dn

[5.4]

Equation [5.3] is solved by: χ0(n) = e−kn

[5.5]

whereas the solution to Equation [5.4] is (Szabò, 1969): χ i ( n) =

(kn)i i!

e- kn

[5.6]

As the total amount of powder in the reactor is constant, Equations [5.5] ∞

and [5.6] satisfy the constraint

∑ χ ( n) = 1 . i

i=0

Equations [5.5] and [5.6] represent the starting set of mathematical functions that must be combined in order to obtain a rough description of MA kinetics. With this aim, let us define j as the minimum number of CLCs that powders must experience to undergo MA. The total volume fraction α(n) of powders involved in MA processes at the nth collision is then equal to:

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α(n) = ∑ χi ( n) ≅ 1 − kne− kn … − i= j

( kn ) j − 1 ( j − 1)!

e- kn

[5.7]

provided that the number n of collisions that have taken place is so large that it tends to infinity. Of course, the kinetic curves can also be obtained by considering the mass fraction αm(n) of powders involved in MA processes at the nth collision event. This quantity can be expressed as: n

α m (n) = ∑ χi ( n) ≅ 1 − km ne− km n … − i= j

(km n) j − 1 ( j − 1)!

e- km n

[5.8]

where km represents the apparent rate constant of the MA process referred to the mass fraction of powder involved in individual collision events. Accordingly, if a single CLC is necessary to induce the MA process, and then j is equal to 1, the mathematical function describing the evolution of the volume fraction of alloyed phase as a function of the total number n of collisions is α(n) = 1 − e−kn. Should two CLCs be required, j would be equal to 2 and the volume fraction of alloyed phase would be α(n) = 1 − (1 + kn)e−kn. The volume fraction k, which measures the fraction of trapped powder that experiences CLCs at collision and then undergoes the MA process, can also be regarded as an apparent rate constant for the MA reaction (Szabò, 1969). Effectively, the volume fraction of powders experiencing CLCs at collision events determines the rate of the MA process. Provided that all the above-mentioned assumptions can be made, the mathematical model discussed here has the remarkable merit of providing the possibility of plotting best-fitting experimental kinetic curves and thus estimating the number j of CLCs necessary to induce the MA process and the apparent rate constant k. It is precisely the shape of kinetic curves that permits discrimination between different cases. However, this possibility is to a large extent dependent on the accuracy of the experimental data. Only refined experimental measurements of the relative amount of alloyed phases can indeed be used profitably to further this aim. A few practical examples of the usefulness of Equations [5.7] and [5.8] are given in the following section, where experimental data concerning different MA reactions under different conditions are considered.

5.4

Collecting and analysing experimental data on the kinetics of mechanical alloying reactions

The necessary conditions for obtaining a satisfactory quantification of the kinetics of MA processes are given by accurate characterization of the solid

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phases formed during the mechanical treatment and by the evaluation of their relative quantities. Any experimental technique able to give information on such aspects can, in principle, be used to quantify the kinetics, but X-ray diffraction (XRD) methods are not the most widespread ones (Klug and Alexander, 1974). Their fundamentals will not be dealt with here. It suffices to remember that obtaining quantitative information on the phase transformation modelling of solid phases is not easy. Long data acquisition times and optimized procedures and equipment are required, the best choice being represented by intense synchrotron radiation facilities (Enzo, 1998). Once accurate data have been acquired, a suitable analysis of XRD patterns must be carried out (Enzo, 1998). This is generally done through a numerical reconstruction of experimental XRD profiles based on a set of mathematical functions with adjustable parameters (Lutterotti et al., 1998). The method, usually referred to as the Rietveld method (Lutterotti et al., 1998), essentially consists of a least-square best-fitting. The different adjustable parameters have a physical meaning such that their value can be used to gain information on the relative amount of a given phase, on its composition and on its microstructural features (Lutterotti et al., 1998). The data obtained are affected by error bars, the size of which depends on the accuracy of data collection and XRD pattern analysis. Under the condition where the error bars are quite small, the data concerning the mass fraction of the different phases can be used to describe the kinetics of the mechanically induced phase transformations. A few examples of the way Equations [5.7] and [5.8] can be employed to gain information on microscopic events governing MA processes are given below. Although mechanically induced phase transformations of pure substances cannot be considered MA processes, it is nevertheless instructive to see what happens in this case. The first example discussed concerns the transition of Co powders from the hexagonal close-packed (hcp) lattice to the face-centred cubic (fcc) one, which has been studied in detail in previous work (Delogu, 2008b). The stable crystallographic phase for Co at room temperature is the hcp one. When it is submitted to mechanical treatment, a transition to the fcc crystallographic phase takes place. The experimental data describing the kinetic modelling of pure Co powders are shown in Fig. 5.1, where the mass fractions αm(n) of hcp and fcc Co phases are reported as a function of the number n of collisions. It can be seen that the hcp phase undergoes an exponential decrease, whereas the fcc one undergoes a complementary exponential increase. The set of data corresponding to the mass fraction of fcc phase is satisfactorily best-fit by Equation [5.8] provided that j is set equal to 1. In other words, a single collision event suffices for Co powders to trigger the transition from the hcp to the fcc phase. The best-fit curve exhibits a km value approximately equal to 1.2 × 10−6 when collisions involve a stainless steel ball of 12 g and take place with an impact energy of about

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1.0 0.8 0.6 0.4 0.2 0.0

0

3

6

9

12

15

18

5

Number of collisions, n (10 )

Mass fraction, a m (n)

5.1 The mass fractions αm(n) of hcp (䊐) and fcc (䊊) Co phases as a function of the number n of collisions. The set of data representing the mass fraction of the fcc phase is best-fitted by Equation [5.8].

1.0 0.8 0.6 0.4 0.2 0.0

0

1

2

3

4

5

6

7

Number of collisions, n (106)

5.2 The mass fractions αm(n) of crystalline (䊐) and amorphous (䊊) NiTi2 phases as a function of the number n of collisions. The set of data representing the mass fraction of the amorphous phase is bestfitted by Equation [5.8].

0.1 J (Delogu, 2008b). This means that, as the powder charge mp is equal to 8 g, only 9 μg of powder experiences CLCs at individual collisions. Analogous behaviour is exhibited by NiTi2 intermetallic powders (Delogu and Cocco, 2003). In this case, the mechanical treatment of NiTi2 powders induces a transition from the ordered crystalline fcc phase to an amorphous one in which long-range crystalline order is completely lost. As shown in Fig. 5.2, where the mass fractions αm(n) of crystalline and amorphous NiTi2 phases are reported as a function of the number n of collisions, the kinetics has a clear exponential character. The modelling is strictly similar to that of the Co powders cited above and, correspondingly, the increase of the mass fraction of the amorphous phase is satisfactorily reproduced by Equation [5.8] with j equal to 1. Although impact energy and ball mass are the same as in the previous case, the apparent kinetic constant km is nevertheless smaller, taking a value of about 0.3 × 10−6 (Delogu and Cocco, 2003). Correspondingly, the mass of powder involved in amorphization processes

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1.0 0.8 0.6 0.4 0.2 0.0 0.0

0.5

1.0

1.5

2.0

2.5

6

Number of collisions, n (10 )

5.3 Mass fraction αm(n) of the amorphous Ni40Ti60 binary alloy as a function of the number n of collisions. The set of data is best-fitted by Equation [5.8].

at individual collisions amounts roughly to 2 μg of a total of 8 g inside the reactor. The case of binary mixtures, that is of true MA processes, exhibits remarkably different features. For example, let us consider the case of Ni40Ti60 powder mixtures, which have been studied in detail in previous work (Cocco et al., 2000; Delogu et al., 2004). Data obtained for an impact energy of about 0.1 J and with a stainless steel ball 12 g in mass are shown in Fig. 5.3, where the mass fraction αm(n) of the amorphous phase is reported as a function of the number n of collisions. It can be seen that the kinetics no longer exhibits an exponential character. On the contrary, data are arranged according to a sigmoidal trend, which points to the existence of some intermediate stage between reactant powders and final amorphous alloy (Szabò, 1969). The experimental points in Fig. 5.3 can be satisfactorily interpolated by Equation [5.8] provided that the parameter j is set equal to 2. It follows that powders must undergo CLCs at least twice before an amorphous phase could be generated. The best-fitting procedure indicates that the apparent kinetic constant km of the MA process amounts roughly to 1.5 × 10−6 (Cocco et al., 2000). Therefore, once the second individual collision has occurred, the powder charge contains approximately 18 pg of amorphous phase. The experimentally obtained km values are very small, so that the amount of powder transformed at individual collisions is correspondingly very small. This suggests that the mechanical stresses at impact are not uniformly distributed within the volume of powders trapped at collisions. On the contrary, a highly inhomogeneous distribution of the mechanical loads can be expected, so that only limited regions around the points of contact between powder particles are actually involved in phase transformation processes, and not all of them are under the same conditions. A very similar modelling is observed for all the binary mixtures composed by Ni and the elements of the Ti group in the periodic table. In all

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cases, two CLCs are needed to induce the formation of a binary amorphous alloy (Cocco et al., 2000). The same is observed for binary mixtures containing Ti and a transition metal belonging to the fourth period of the periodic table (Delogu, 1999). Of course, the nature of the final product changes with the chemical nature of the binary mixture. In particular, whereas Fe–Ti, Co–Ti, Ni–Ti and Cu–Ti mixtures give rise to amorphous alloys, V–Ti and Cr–Ti ones exhibit the formation of metastable crystalline solid solutions (Delogu, 1999). Again, only two CLCs are required for the final phase to appear. The conceptual framework behind the experimental kinetic evidence and the theoretical interpretation discussed above is, however, much more complex than the approximate description based on a few key parameters such as j and km. First, it should be noted that km strongly depends on the impact energy, that is on the amount of kinetic energy transferred to powders at individual collisions (Delogu, 1999; Cocco et al., 2000; Delogu et al., 2004; Delogu and Cocco, 2007; Delogu, 2008a,b). This dependence is typically linear and points to the existence of a threshold energy below which no structural transformation can take place (Delogu, 1999; Cocco et al., 2000; Delogu et al., 2004; Delogu and Cocco, 2007; Delogu, 2008a,b). Second, and much more interesting, j also depends on the energy of individual collisions (Delogu and Cocco, 2007). This means that the number of CLCs to which the powder must be submitted to undergo a MA process depends on the energy transferred to the powder at impact. In turn, this again suggests that the extent to which elemental species alloy at collision depends on the intensity of mechanical stresses at the points of contact between surfaces. Unfortunately, this is not amenable to direct experimental investigation. It is, however, the area in which numerical simulation methods can be most fruitfully employed. These will be dealt with in some detail in the next section. Before doing this, it is worth mentioning a particular kinetic regime observed in a few systems consisting of a cyclic mixing–demixing process that takes place under mechanical treatment conditions. As broadly discussed in previous work (Johnson et al., 2006; Lee, 2008), the mechanical treatment occasionally gives rise to cyclic modelling in microstructural and phase transformations. This is the case, for example, in the cyclic amorphization of Co-based systems (El-Eskandarany et al., 1997, 2002; ElEskandarany and Inoue, 2007). Previous investigations have shown that the occurrence of cyclic MA processes critically depends on the interplay of various mechanical processing parameters, including impact energy, temperature and powder mixture composition (Johnson et al., 2006; Lee, 2008). All of this clearly suggests that a simple and comprehensive description of MA kinetics is unlikely to be obtained on the basis of the current knowledge. On the contrary, considerable efforts should be devoted to clarifying

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various fundamental aspects of MA processes. Among these, their atomistic features are probably the most important.

5.5

Mechanisms of chemical mixing on the atomic scale

As briefly mentioned in the Introduction, relatively few studies have been carried out that shed light on the intimate details of mechanically induced processes taking place on the atomic scale. All of them consist substantially of numerical investigations (Martin, 1984; Bellon and Averback, 1995; Hammerberg et al., 1998; Urakaev and Boldyrev, 2000; Vick et al., 2000; Fu et al., 2001; Lund and Schuh, 2003; Delogu and Cocco, 2005a,b; Odunuga et al., 2005; Vo et al., 2009). Phenomenological approaches have laid emphasis in the past on the capability of shear stresses of inducing stochastic atomic displacements to be able to withstand, at least under suitable conditions, opposite thermodynamic forces (Martin, 1984; Martin and Gaffet, 1990). However, the mechanisms according to which such displacements take place remained for a long time completely obscure. A very elegant experimental study based on electron microscopy and atom probe tomography techniques recently allowed insight to be gained into the atomistic processes governing MA in immiscible binary systems (Zgahl et al., 2002a,b; Wu et al., 2006). It was shown that the particular nanostructure of immiscible alloys arises as a consequence of a relatively complex interplay between the tendency to phase separation, which is a consequence of the thermodynamic driving force operating in the system, and the disordering phenomena due to forced mixing. The mechanistic details of such forced mixing were further investigated in successive works based on numerical simulation methods (Odunuga et al., 2005; Delogu, 2008c; Vo et al., 2009). These have shown that intense plastic deformation regimes can generate unusual atomic flows, mediated by dislocation motion, reminiscent of the ones observed in turbulent fluids. The occurrence of disordered atomic displacements and rearrangements has been also pointed out in two- and three-dimensional molecular dynamics simulations aimed at studying the atomistic processes at sliding interfaces (Hammerberg et al., 1998; Fu et al., 2001; Delogu and Cocco, 2005a,b). In particular, numerical findings suggested the existence of a hierarchy of processes involving the creation of defective atoms, that is atoms with defective coordination. These assure the attainment of a relatively high degree of local mobility and allow the occurrence of mixing phenomena, which represent the fundamental mechanistic step for obtaining a so-called driven alloy. Finally, once the mechanical load has been removed, or has locally ceased its action, the clusters of defective atoms progressively relax, without

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Square of layer thickness, Δ2 (nm2)

necessarily completely disappearing. All of these processes must be regarded as a relatively complex sequence of intertwined phenomena dependent on a number of parameters, including the intensity of interatomic forces, the intensity of the mechanical stress and the duration of the deformation event (Hammerberg et al., 1998; Fu et al., 2001; Delogu and Cocco, 2005a,b). Interestingly, simulations pointed out that the time dependence of the volume of crystalline phase involved in chemical mixing and disordering processes is analogous to the one exhibited by isothermal vacancy-mediated diffusion mechanisms (Fu et al., 2001; Delogu and Cocco, 2005a). An example relevant to the case of amorphization reactions is shown in Fig. 5.4, where the square of the thickness Δ of the layer of crystalline phase that underwent disordering upon mechanical loading is reported as a function of the time t for the case of a Ni–Zr sliding interface (Delogu and Cocco, 2005a,b). It can be seen that the data are arranged according to a linear plot despite the conditions for the growth of the amorphous interfacial layer being far from similar to the ones of isothermal diffusion. Rather, the displacement of atoms is mediated by the clusters of defective atoms and the formation of particular rotating structures, such as the one shown in Fig. 5.5. These groups of atoms, submitted to local torques, are able to displace individual atoms at relatively high rates, enabling them to overcome the ideal interface and alloy into the other elements. Of course, the continuous nucleation and propagation of dislocations and partial dislocations permit these atoms to migrate successively far from the interface region (Delogu and Cocco, 2005a,b). Mixing and disordering processes during deformation are accompanied by continuous, though inhomogeneous, transfer of mechanical energy to the system. This can produce significant local rises of temperature. As shown in Fig. 5.6, temperatures as high as 1500 K can be reached. However, local heat is rapidly dissipated owing to the very steep temperature gradients operating in relatively small volumes. For example, data in Fig. 5.7 indicate that 30 20 10 0 0.0

0.5

1.0

1.5 2.0 Time, t (ns)

2.5

3.0

5.4 Square of the thickness Δ of the amorphous layer, formed upon mechanical loading at a Ni–Zr sliding interface, as a function of the time t. The line of best-fit is also shown.

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Local temperature, T (K)

5.5 Cluster of atoms undergoes rotation, as indicated by the large white arrows. Small arrows indicate the direction of displacement of the individual atoms. White and grey indicate, respectively, Ni and Zr atoms. The dashed horizontal line ideally marks the interface between Ni and Zr crystals. The small cross indicates the instantaneous axis of rotation of the cluster.

2000 1500 1000 500 0.2

0.3

0.4 Time, t (ns)

0.5

0.6

5.6 Local temperature T as a function of time t. Data refer to a square domain with sides about 1.7 nm long at the interface between Ni and Zr crystals.

temperatures as different as about 700 and 1000 K are present in a small square region with sides about 1.7 nm long (Delogu and Cocco, 2005a,b). Further numerical studies have also shown that the topology of colliding surfaces and the geometry of their collision can significantly affect the resulting chemical mixing and disordering processes, as well as the local distribution of temperature (Delogu and Cocco, 2006). In particular, it was

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5.7 Map of instantaneous temperature (in K) for a square domain with sides about 1.7 nm long at the interface between Ni and Zr crystals.

shown that the collision between rough surfaces, more effectively than smooth surfaces, results in a local disordering and temperature increase. The displacement of atomic species is mediated by the formation of force chains along which repulsive forces preferentially operate, thus determining the occurrence of atomic rearrangements through the possible path at lowest potential energy. Despite the considerable efforts devoted to the field and the small advances achieved, the intimate mechanisms governing mechanically induced processes remain a matter of debate. The role of localized mechanical stresses as well as of local structural perturbations need suitable clarification. Owing to the complicated nature of mechanochemical transformations, which are governed by intertwined hierarchies of elementary events on the atomic scale, numerical simulations are expected to play an increasingly important role in the definition of detailed mechanistic scenarios.

5.6

Future trends

In the light of the considerable work carried out recently by different research groups around the world, it is astonishing that the view given in a paper published about 14 years ago (Khina and Froes, 1996) can still stand. The large gap between macroscopic kinetic evidence and microscopic transformation mechanisms must be filled. Similarly, the thermodynamic description of MA processes must be improved in order to include the nonequilibrium features typical of atomic rearrangement processes taking place at mechanical loading events. Despite apparently well defined objectives, a great deal of work remains to be done from both the experimental and theoretical points of view.

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With regard specifically to the kinetic and mechanistic aspects of MA processes, at least four main topics can be clearly pointed out that need particularly urgent attention in order to attain a significant advance in the field. 1. Refinement of global models linking the operation of milling devices to the dynamics of powders inside the reactor. 2. Refinement of local models describing the relationship between the mechanical action exerted by milling tools at impact events and the corresponding behaviour of powders. 3. Integration of suitably improved kinetic models based on a statistical approach to the description of phase evolution with phenomenological kinetic laws accounting for the phase transformation behaviour on a microscopic level. 4. Rationalization of phenomenological kinetic evidence on the basis of atomistic models of mechanically induced atomic rearrangement processes. Each and every one of the above-mentioned areas of investigation requires the dedication of intense research effort. Effectively, all of them define a complex and variegated conceptual framework extending from the dynamics of milling devices to the displacement of atomic species. Between these extreme boundaries, there is the evident need to relate the motion of milling tools to the regime of impacts inside the reactor, to shed light on the behaviour of powder particles submitted to high strain rate compaction at impact, and to connect the topology of the net formed by the points of contact between the powder particles to the local mechanical stresses operating on them. In turn, this would allow realistic studies to be carried out on the local deformation processes and the corresponding response of interfaces between both neighbouring particles and inside the same particles to be carried out, in order to rationalize the cold-welding and fragmentation behaviour of powder particles and to gain insight into the stress and temperature conditions under which MA reactions take place at the atomistic levels. Finally, it should be possible to link the mechanical response of the materials to the local alloying behaviour, eventually identifying the finest aspects of atomistic mechanisms. The above-mentioned considerations clearly define a series of complementary studies aimed at creating deeper and better connections between the different scales involved in MA processes. The briefly sketched plan of research activity outlined here may permit a deepening of the current knowledge on MA. It also invites more intense and frequent collaboration between the different researchers involved in the field. It is with the sincere wish that such collaborations can be realized in the near future that this chapter ends.

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5.7

References

bellon p and averback r s (1995), ‘Nonequilibrium roughening of interfaces in crystals under shear: application to ball milling’, Phys. Rev. Lett., 74, 1819–22. benjamin j s (1992), ‘Fundamentals of mechanical alloying’, Mater. Sci. Forum, 88–90, 1–17. butyagin p yu (1989), ‘Active states in mechanochemical reactions’, Sov. Sci. Rev. Chem. B, 14, 1–134. cahn r w and greer a l (1996), ‘Metastable states of alloys’, in Physical Metallurgy: fourth, revised and enhanced edition, Cahn R W and Haasen P (eds), Elsevier Science BV, Amsterdam, 1723–1830. cocco g, delogu f and schiffini l (2000), ‘Toward a quantitative understanding of the mechanical alloying process’, J. Mater. Synth. Proc., 8, 167–80. courtney t h (1995), ‘Process modeling of mechanical alloying’, Mater. Trans., JIM, 36, 110–22. delogu f (1999), Solid Phase Reactivity under Mechanical Processing Conditions. Structural Evolution and Phase Transformation Kinetics, PhD Thesis, University of Sassari, Sassari. delogu f (2008a), ‘A combined experimental and numerical approach to the kinetics of mechanically induced phase transformations’, Acta Mater., 56, 905–12. delogu f (2008b), ‘Kinetics of allotropic phase transformation in cobalt powders undergoing mechanical processing’, Scripta Mater., 58, 126–9. delogu f (2008c), ‘Forced chemical mixing in model immiscible systems under plastic deformation’, J. Appl. Phys., 104, 073533. delogu f and cocco g (2003), ‘Impact-induced disordering of intermetallic phases during mechanical processing’, Mater. Sci. Eng. A, 343, 314–17. delogu f and cocco g (2005a), ‘Molecular dynamics investigation on the role of sliding interfaces and friction in the formation of amorphous phases’, Phys. Rev. B, 71, 144108. delogu f and cocco g (2005b), ‘Numerical simulations of structural modifications at a Ni-Zr sliding interface’, Phys. Rev. B, 72, 014124. delogu f and cocco g (2006), ‘Numerical simulations of atomic-scale disordering processes at impact between two rough crystalline surfaces’, Phys. Rev. B, 74, 035406. delogu f and cocco g (2007), ‘Kinetics of amorphization processes by mechanical alloying: a modeling approach’, J. Alloys Compd., 436, 233–40. delogu f, mulas g, schiffini l and cocco g (2004), ‘Mechanical work and conversion degree in mechanically induced processes’, Mater. Sci. Eng. A, 382, 280–7. el-eskandarany m s and inoue a (2007), ‘Mechanically induced cyclic metastable phase transformations of Zr2Ni alloys’, Phys. Rev. B, 75, 224109. el-eskandarany m s, aoki a, sumiyama k and suzuki k (1997), ‘Cyclic crystalline– amorphous transformations of mechanically alloyed Co75Ti25’, Appl. Phys. Lett., 70, 1679–82. el-eskandarany m s, aoki a, sumiyama k and suzuki k (2002), ‘Cyclic phase transformations of mechanically alloyed Co75Ti25 powders’, Acta Mater., 50, 1113–23. enzo s (1998) ‘Diffraction of amorphous and nanocrystalline alloys prepared by solid state reactions’, Mater. Sci. Forum, 269–72, 363–72. fu x y, falk m l and rigney d a (2001), ‘Sliding behavior of metallic glass. Part II. Computer simulations’, Wear, 250, 420–30.

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hammerberg j e, holian b l, roder j, bishop a r and zhou s j (1998), ‘Nonlinear dynamics and the problem of slip at material interfaces’, Physica D, 123, 330–40. heinicke g (1984), Tribochemistry, Akademie-Verlag, Berlin. hellstern e and schultz l (1986), ‘Amorphization of transition metal Zr alloys by mechanical alloying’, Appl. Phys. Lett., 48, 124–7. johnson w l (1986), ‘Thermodynamic and kinetic aspects of the crystal to glass transformation in metallic materials’, Prog. Mater. Sci., 30, 81–134. johnson w l, lee j k and shiflet g j (2006), ‘Thermodynamic treatment of cyclic amorphization during ball milling’, Acta Mater., 54, 5123–33. khina b b and froes f h (1996), ‘Modelling mechanical alloying – Advances and challenges’, J. Metals, 48, 36–8. klug h p and alexander l e (1974) X-ray Diffraction Procedures, John Wiley & Sons, New York. koch c c (1992), ‘The synthesis of non-equilibrium structures by ball-milling’, Mater. Sci. Forum, 88–90, 243–62. lee j k (2008), ‘On cyclical phase transformations in driven alloy systems’, Metall. Mater. Trans. A, 39, 964–75. levitas v i (2004) ‘Continuum mechanical fundamentals of mechanochemistry’, in High Pressure Surface Science and Engineering, Gogotsi Y and Domnich V (eds), Institute of Physics, Bristol, 159–292. lund c and schuh c a (2003), ‘Atomistic simulation of strain-induced amorphization’, Appl. Phys. Lett., 82, 2017–19. lutterotti l, ceccato r, dal maschio r and pagani e (1998) ‘Quantitative analysis of silicate glass in ceramic materials by the Rietveld method’, Mater. Sci. Forum, 278–281, 87–92. manai g, delogu f and rustici m (2001) ‘Onset of a chaotic dynamics in a ball mill: attractors merging and crisis induced intermittency’, Chaos, 12, 601–9. manai g, delogu f, schiffini l and cocco g (2004), ‘Mechanically-induced selfpropagating combustions: experimental findings and numerical simulation results’, J. Mater. Sci., 39, 5319–25. martin g (1984), ‘Phase stability under irradiation: ballistic effects’, Phys. Rev. B, 30, 1424–36. martin g and gaffet e (1990) ‘Mechanical alloying: far from equilibrium transitions?’, J. Phys. (Paris), C-4, 15, 71–7. martin c l, bouvard d and shima s (2003), ‘Study of particle rearrangement during powder compaction by the discrete element method’, J. Mech. Phys. Solid, 51, 667–93. maurice d r and courtney t h (1996), ‘Milling dynamics. Part III. Integration of local and global modeling of mechanical alloying devices’, Metall. Mater. Trans. A, 27, 1981–6. odunuga s, li y, krasnochtchekov p, bellon p and averback r s (2005), ‘Forced chemical mixing in alloys driven by plastic deformation’, Phys. Rev. Lett., 95, 045901. schwarz r b and johnson w l (1983), ‘Formation of an amorphous alloy by solid-state reaction of the pure polycrystalline metals’, Phys. Rev. Lett., 51, 415–18. suryanarayana c (2001), ‘Mechanical alloying and milling’, Prog. Mater. Sci., 46, 1–184.

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szabò z g (1969), ‘Kinetic characterization of complex reaction systems’, in Comprehensive Chemical Kinetics. The theory of kinetics, Bamford C H and Tipper C F H (eds), Elsevier Science Publishers, Oxford, vol. 2, 1–80. urakaev f kh and boldyrev v v (2000), ‘Mechanism and kinetics of mechanochemical processes in comminuting devices: 2. Applications of the theory. Experiment’, Powder Tech., 107, 197–206. vick b, furey m j and iskandar k (2000), ‘Theoretical surface temperatures generated from sliding contact of pure metallic elements’, Tribol. Int., 33, 265–71. vo n q, odunuga s, bellon p and averback r s (2009), ‘Forced chemical mixing in immiscible alloys during severe plastic deformation at elevated temperatures’, Acta Mater., 57, 3012–19. wu f, isheim d, bellon p and seidman d n (2006), ‘Nanocomposites stabilized by elevated-temperature ball milling of Ag50Cu50 powders: an atom probe tomographic study’, Acta Mater., 54, 2605–13. zgahl s, hytch m j, chevalier j- p, twesten r, wu f and bellon p (2002a), ‘Electron microscopy nanoscale characterization of ball-milled Cu–Ag powders. Part I: Solid solution synthesized by cryo-milling’, Acta Mater., 50, 4695–709. zgahl s, twesten r, wu f and bellon p (2002b), ‘Electron microscopy nanoscale characterization of ball milled Cu–Ag powders. Part II: Nanocomposites synthesized by elevated temperature milling or annealing’, Acta Mater., 50, 4711–26.

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6 Mechanochemical synthesis of complex ceramic oxides T. R O JAC and M. K O S E C, Jozˇef Stefan Institute, Slovenia

Abstract: Mechanochemical synthesis has recently been recognized as a promising method for synthesizing a variety of technologically important ceramic oxides with complex compositions. This chapter begins by presenting some recent results related to the mechanisms and kinetics of mechanochemical reactions in oxide systems. In the second part of the chapter, attention is turned to the literature data on direct mechanochemical synthesis and the mechanochemical activation-based synthesis of a variety of complex oxides having different properties. This includes ferroelectric oxides and related materials, magnetic oxides and oxides with semiconducting and catalytic properties. Key words: mechanochemical synthesis, mechanochemical activation, high-energy milling, complex ceramic oxides, reaction mechanisms.

6.1

Introduction

Mechanochemical synthesis has recently received a lot of interest for processing ceramic powders, opening up new ways of producing technologically important oxides with complex compositions. One of the most important research areas is the study of the mechanisms and kinetics of mechanochemical reactions, which form the basis of the further development of mechanochemical synthesis. The first part of the chapter is devoted to this subject and discusses various topics, including the relation between milling conditions and phase formation, the course of mechanochemical reactions including amorphization and nucleation phenomena triggered by high-energy milling, the formation of intermediate phases during a reaction, the evolution of the crystallite size and the particularities of the nanocrystalline products obtained. In general, complex ceramic oxides exhibit a variety of properties, making them attractive for a range of different applications. The second part of the chapter highlights the main results related to both direct mechanochemical synthesis (referring to the formation of the final product directly during milling) and the mechanochemical activation-based synthesis (referring to high-energy milling, used to increase the reactivity of powders, 113 © Woodhead Publishing Limited, 2010

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which is followed by an annealing step to induce the formation of the final product) of complex oxides with a variety of properties. This section includes ferroelectric, antiferroelectric, piezoelectric, relaxor and multiferroic oxides as well as oxides with magnetic, semiconducting and catalytic properties. From the extensive literature data available some key points are highlighted, in particular the influence of mechanochemical activation on both the phase formation during subsequent annealing and the preparation of the final ceramics, the influence of milling conditions, such as humidity, and the effect of the hydration state of the reagents on the course of the mechanochemical reaction as well as the issue of contamination during milling.

6.2

Mechanisms and kinetics of mechanochemical reactions in complex oxide systems

One of the main tasks when exploring the basics of the mechanochemical synthesis is to correlate the parameters that describe the milling operation with the formation of the final product during milling. Rojac et al. (2006b) were the first to apply the concept of the milling map, which was originally developed for metal systems, to the synthesis of ceramic oxide powders. They chose the mechanochemical reaction between Na2CO3 and Nb2O5, which yields NaNbO3, as a model system. The reaction was carried out under different milling conditions. From the milling parameters characteristic of a planetary mill (Fig. 6.1a) the authors calculated the ball-impact energy (ΔEb) (amount of kinetic energy of a milling ball released to the powder during a collision, in J) and the ball-impact frequency (νt) (number of ball impacts per unit time, in s−1) applied in each milling experiment using the Equations [6.1] and [6.2], respectively: ⎛ ⎛ W ⎞ D − db ⎞ 2 ⎛ W ⎞ 1 ⎛ π d3 ⎞ ΔEb = ⋅ ⎜ ρb b ⎟ ⋅ Wp2 ⎜ ⎜ v ⎟ ⎛ v ⋅⎜ 1 − 2 v ⎟ − ⎝ ⎠ ⎝ ⎠ ⎝ Wp ⎠ 2 6 2 ⎝ ⎝ Wp ⎠ 2

⎛ W ⎞ D − db ⎞ ⎛ Wv ⎞ ⎛ Dv − db ⎞ 2 ⎞ − 2 Rp ⎜ v ⎟ ⋅ ⎛ v ⎟ ⎝ Wp ⎠ ⎝ 2 ⎠ ⎜⎝ Wp ⎟⎠ ⎝ 2 ⎠ ⎠ 2

νt = NbK(Wp − Wv)

[6.1]

[6.2]

where K is a constant depending on the geometry and on the ball diameter (Rojac et al., 2006b). The mechanochemical formation of NaNbO3 under specific energy and frequency conditions was followed using X-ray diffraction (XRD) analysis. This allowed the authors to construct the milling map shown in Fig. 6.1b. Their results suggested that a minimum or critical cumulative kinetic energy

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(a) Diameter db

Diameter Dv

Number Nb

Height Hv Milling balls

Milling vial

Density rb

Powder weight mp

Milling parameters

Operation Disc (plate) rotational speed Wp

Mill Distance between rotational axes Rp

Vial rotational speed Wv

(b) 200 10 min DEb∗ (mJ per hit)

150 15 min 100

30 min

50

60 min

0 1

10

100

1000

–1

Ecum (kJ g )

6.1 (a) Milling parameters for the case of a planetary mill and a cylindrical milling vial and (b) the milling map of NaNbO3. Notation: ΔE*b ball-impact energy, Ecum cumulative kinetic energy, 䊊 no NaNbO3 is detected in the mixture, 䊉 NaNbO3 is detected in the mixture according to XRD analysis; the milling times of the first appearance of NaNbO3 according to XRD analysis are given on the figure (from Rojac et al., 2006b).

(the weight-normalized product of the ball-impact energy, the ball-impact frequency and the milling time) of between 7 and 12 kJ g−1, which does not depend on the ball-impact energy applied, is needed to trigger the formation of the NaNbO3. Milling for a duration equivalent to a cumulative kinetic energy lower than this critical value, however, will not result in the

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formation of NaNbO3. In addition, by comparing several oxide systems on the basis of the experimental data available in the literature the authors also stressed the importance of the critical ball-impact energy, which determines, in a similar manner, whether a certain product will form or not under specific milling conditions. The mathematical model used to construct the milling maps described above assumes that collisions are the prevailing mechanism by which the energy is transferred from the milling balls to the powders, while friction is neglected. However, in some circumstances the friction component might play a crucial role in the reaction kinetics, as was explained by Jean and Nachbaur (2008) for their mechanochemical synthesis of ZnFe2O4. They used a Fritsch Pulverisette 4 planetary mill, which allows independent adjustment of the rotational frequency of the supporting disc and that of the grinding vials. On the basis of the extent of the mechanochemical reaction between ZnO and α-Fe2O3 they estimated the milling efficiency for various ratios of the two frequencies, denoted as k in Equation [6.3]. The results are shown in Fig. 6.2:

600 550

Disc rotation speed (r.p.m.)

500 450 400

+ A No efficient milling

350

+ B High-efficiency milling

+ C High-efficiency milling

Pure ZnFe2O4 after 18 h

Pure ZnFe2O4 after 12 h

+ E Low-efficiency milling

+ F High-efficiency milling

Trace of ZnFe2O4 after 18 h

Pure ZnFe2O4 after 36 h

300 250 200 150

+ D No efficient milling

100 50 0 200

300

400

500

600

700 –1

Vial rotation speed, x

800

900

1000

(r.p.m.)

6.2 Graphical representation of the milling efficiency estimated for various ratios of disc and vial rotational frequencies (mechanochemical synthesis of ZnFe2O4). The end products of the milling experiments performed at different frequency ratios are given on the graph (from Jean and Nachbaur, 2008).

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Wv Wp

117 [6.3]

In Equation [6.3], Wp and Wv are the rotational frequencies of the supporting disc and of a grinding vial of the planetary mill, respectively. The negative sign indicates that the rotations of the disc and the vial are in opposite directions. Abdellaoui and Gaffet (1995), in their work on metal systems, reported that the end product obtained by milling depends on the injected mechanical power, defined as the product of ball-impact energy and ball-impact frequency. Because the injected mechanical power increases as a function of the rotational frequencies of the disc and the vial, Jean and Nachbaur (2008) were able to explain the increasing milling efficiency in experiments A to C or D to F (Fig. 6.2), with increasing injected power. However, this did not explain, for example, the low milling efficiency in the case of experiment A, where no reaction occurred, during which more injected power was used than in experiment F, where the spinel phase was obtained. This observation meant that it was also necessary to include a friction component in the discussion. Based on literature data, which show that the tangential component of the impact force during a collision event is nearly zero for |k| < 1, while its absolute value increases rapidly for |k| > 1 (Chattopadhyay et al., 2001), the authors related the high milling efficiency of experiment F (|k| = 3.6) to the large frictional force during milling, while the low milling efficiency in experiment A (|k| = 0.5) is due to a low friction component. The results clearly showed that the injected power and the friction phenomenon have to be considered simultaneously, especially if the k-ratio is varied over a broad range. The importance of the milling conditions during the synthesis of a target oxide is best demonstrated by the example of BaTiO3. It was shown that single-phase BaTiO3 can be synthesized mechanochemically from a mixture of BaO and TiO2 by milling in a vacuum (Welham, 1998) or a nitrogen atmosphere (Xue et al., 2000a). However, several authors have reported their inability to synthesize BaTiO3 from a mixture of BaCO3 and TiO2 (Stojanovic´ et al., 1999; Gomez-Yanez et al., 2000; Brzozowski and Castro, 2000, 2003; Berbenni et al., 2001a; Pavlovic´ et al., 2002; Kong et al., 2002a; Goes et al., 2008). Xue et al. (2000a) claimed that it was impossible to prepare BaTiO3 by milling BaO and TiO2 in a shaker mill in air owing to the formation of BaCO3 from the BaO and atmospheric CO2. This point was discussed in the work of Welham (1998), who explained, using thermodynamic calculations, the high stability of BaCO3 in the reaction with TiO2 in comparison with BaO or Ba(OH)2. A positive free energy was calculated for the reaction between BaCO3 and TiO2 (ΔG298 = 60.3 kJ mol−1), in contrast to the negative

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free energies of the reactions between BaO and TiO2 (ΔG298 = −158 kJ mol−1) and Ba(OH)2 and TiO2 (ΔG298 = −60.7 kJ mol−1). However, as was mentioned in the paper, the influence of the milling conditions on the reaction should not be overlooked. As a matter of fact, in a later investigation, Stojanovic´ et al. (2005) showed that BaTiO3 can be synthesized from a BaO– TiO2 mixture in air using a planetary mill, even if a substantial amount of BaCO3 had already been observed in the initial stage of milling. In addition, a recent publication of Ohara et al. (2008a) demonstrates that single-phase BaTiO3 can be synthesized in just 12 min from a BaCO3–TiO2 mixture using a highly efficient attrition-type mill operating at a rotational frequency of 4000 min−1 (Fig. 6.3). In the past few years great efforts have been made to understand the basic mechanisms that govern the mechanochemical reaction. The mechanistic aspect of reactions of mixtures of oxides, which are, according to the literature, the most frequently used reagents in the mechanochemical synthesis of complex ceramic oxides, has been studied by Wang, Xue and coworkers (Wang et al., 2000a). Their work involves the synthesis of several materials, including ferroelectrics, relaxors and relaxor-ferroelectrics, such as PbTiO3 (PT) (Xue et al., 1998), Pb(Zr0.52Ti0.48)O3 (PZT) (Wang et al., 2000a), Pb(Mg1/3Nb2/3)O3 (PMN) (Wang et al., 1999, 2000a; Xue et al., 2002), Pb(Zn1/3Nb2/3)O3 (PZN) (Wang et al., 2000b), 0.36Pb(Mg1/3Nb2/3) (a)

(b) BaTiO3 BaCO3 TiO2

Powders

Intensity (a.u.)

After

Chamber Rotor Before

20

30

40

50

60

70

2θ (deg)

6.3 (a) XRD patterns of BaCO3–TiO2 powder mixture before and after 12 min of high-energy milling (from Ohara et al., 2008a) and (b) schematic view of the attrition-type mill employed (from Sato et al., 2006).

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O3–0.54Pb(Zn1/3Nb2/3)O3–0.1PbTiO3 (0.36PMN–0.54PZN–0.1PT) (Wan et al., 1999; Wang et al., 2000a, 2000b; Xue et al., 2002), 0.9Pb(Zn1/3Nb2/3) O3–0.1BaTiO3 (0.9PZN–0.1BT) (Xue et al., 1999a) and SrBi4Ti4O15 (SBIT) (Ng et al., 2002a). Wang et al. (2000a) explained the mechanochemical synthesis of PMN from a PbO–MgO–Nb2O5 mixture by dividing the process into different stages. In the initial part of the mechanochemical treatment they first observed a crystallite-size refinement of the initial oxides coupled with a partial amorphization of the mixture. At this stage the PMN phase formed from the activated mixture and, with further milling, the amount of PMN increased. According to XRD analysis, a single nanocrystalline PMN phase was obtained after 20 hours of milling. They also observed a sharpening of the XRD peaks of the PMN during its formation and estimated, using the Scherrer equation, an increase in the size of the PMN crystallites from 9 nm to 15 nm when extending the milling time from 5 hours to 20 hours. This suggested that the perovskite nanocrystallites underwent a steady growth during the mechanochemical treatment. The authors claimed that pyrochlore phases, such as Pb3Nb4O13 and Pb2Nb2O7, which are typically encountered as intermediate reaction products during the temperaturedriven solid-state reaction, were not involved in the process. They believe that the formation of the PMN occurs via nucleation from an activated amorphous matrix involving all three oxides, rather than being a result of an interfacial ‘reaction and diffusion’ mechanism, which is characteristic for thermally activated reactions and results in the formation of transitional pyrochlore phases. On the basis of similar observations in other systems, the authors explained the mechanochemical synthesis as a ‘nucleation and growth’ process in which the product nanocrystallites nucleate from a highly activated amorphous matrix and grow with continued milling. In Fig. 6.4 we indicate schematically the stages involved in the reaction mechanism. The mechanism proposed by Wang et al. (2000a) was partially put into question after the work of Kuscer et al. (2006, 2007). As a matter of fact, they observed the formation of pyrochlore phases during the mechanochemical synthesis of Pb(Mg1/3Nb2/3)O3 (PMN) and 0.65Pb(Mg1/3Nb2/3) O3–0.35PbTiO3 (PMN–PT) solid solutions. In the initial stage of the mechanochemical synthesis of PMN from a PbO–MgO–Nb2O5 powder mixture, a lead-niobate pyrochlore was formed first (Kuscer et al., 2006). Based on the results of a Rietveld refinement analysis coupled with an internal standard method, which was used to determine the amount of amorphous phase from the XRD background, a change in the pyrochlore’s lattice parameter, as well as the consumption of MgO from the crystalline and amorphous phases, were observed with further milling. This suggested that a lead–magnesium–niobate pyrochlore, that is Pb1.86Mg0.24Nb1.76O6.5,

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Mixture of oxides

Mechanochemical treatment

Reduction of crystallite size and amorphization

Mechanochemical treatment

Growth of product nanocrystallites

Mechanochemical treatment

Nucleation of the product from the amorphous phase

6.4 Schematic view of the ‘nucleation and growth’ mechanism proposed by Wang, Xue and co-workers (Wang et al., 2000a).

was also formed subsequently owing to the incorporation of the MgO into the pyrochlore structure. At the same time, the PMN perovskite phase appeared in the mixture. The amount of pyrochlore phase reached a maximum during milling, while the amount of perovskite phase increased almost linearly. Since the quantity and the composition of the amorphous phase changed in the course of milling, it was suggested that it takes part in reactions related to the formation of both perovskite and pyrochlore phases. As in the case of the PMN synthesis, Kuscer et al. (2007) showed that the formation of the PMN–PT solid solution in the initial stage of milling is characterized by nucleation from the amorphous phase. This is consistent with the previously discussed mechanism of Wang et al. (2000a) (Fig. 6.4). However, after a certain milling time, when the PbO, TiO2 and Nb2O5 were no longer identified in a crystalline form and MgO had disappeared from the amorphous phase, a significant decrease in the formation rate of PMN– PT was observed (Fig. 6.5). As there was no MgO in the amorphous phase, the authors assumed that the perovskite could not simply nucleate from the amorphous phase. The amount of pyrochlore decreased even though the amount and the composition of the amorphous phase did not change significantly with increasing milling time from 32 to 64 hours (see Fig. 6.5b, c). Based on this, the authors supposed that the formation mechanism of the perovskite phase changed to a pyrochlore-to-perovskite transformation, which is predominant after prolonged milling and occurs in a way similar to conventional solid-state synthesis. When comparing the mechanisms proposed by Kuscer et al. (2006, 2007) and Wang et al. (2000a) it should be mentioned that the latter put together their explanation on the basis of not observing any pyrochlore phases being formed during the mechanochemical synthesis of PMN and

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Mechanochemical synthesis of complex ceramic oxides (a) 70 18 16 14 12 10 8 6 4 2 0

121

Mass (%)

PbO TiO2 Nb2O5 MgO 0

10

20

30 40 50 Milling time (h)

(b) 70

60

70

Per

60 Mass (%)

50 AF

40 30 20 10

PY

0 0

10

20

30 40 50 Milling time (h)

60

70

Amorphous phase (mol)

(c) 0.35 0.3 PbO

0.25 0.2 0.15

NbO2.5

0.1

TiO2

MgO

0.05 0 0

10

20 30 40 50 Milling time (h)

60

70

6.5 Quantitative phase composition of high-energy milled PbO–MgO– Nb2O5–TiO2 powder mixture as a function of milling time. (a) Fraction of PbO, MgO, Nb2O5 and TiO2 in crystalline form; (b) fraction of perovskite PMN–PT (Per), pyrochlore (PY) and amorphous (AF) phases; and (c) calculated composition of the amorphous phase (from Kuscer et al., 2007).

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0.36PMN–0.54PZN–0.1PT. This observation was made primarily on the basis of an XRD analysis, where a peak overlapping with PbO makes identification of the pyrochlore less than straightforward. Secondly, the authors used a shaker mill for the synthesis, whereas a planetary mill was used in the work of Kuscer et al. (2006, 2007); this could result in some differences in the formation of pyrochlore phases. In order to elucidate this point, further studies should be more focused on the relation between the milling conditions and the formation of transitional pyrochlore phases during the mechanochemical synthesis of PMN and related solid solutions. Zdujic´ et al. (2006) studied the mechanochemical synthesis of Bi4Ti3O12, determining the amount of amorphous phase formed during milling and following the crystallite size as a function of the milling time. In addition to the stoichiometric Bi2O3–TiO2 mixture, a preformed Bi4Ti3O12 powder was also subjected to mechanochemical treatment. First, they found that the amorphization of the Bi4Ti3O12 compound was realized only when its average crystallite size decreased below 40 nm. The amorphization mechanism was therefore explained as a collapse of the crystal structure after a certain critical crystallite size was reached, with a subsequent dissolution of the nanocrystallites into the amorphous matrix. On the other hand, milling the Bi2O3–TiO2 mixture resulted first in the formation of Bi2(CO3)O2 as an intermediate product (due to the reaction of the Bi2O3 with atmospheric CO2), which reacted subsequently with the TiO2, finally forming the Bi4Ti3O12. In both cases, however, they found that the form of the Bi4Ti3O12 obtained by milling, that is amorphous or nanocrystalline, depends strongly upon the milling conditions applied. While a higher injected mechanical power (the product of the ball-impact energy and the ball-impact frequency) favours a nanocrystalline phase, a lower power results in an amorphous one. For example, by increasing the power during a milling experiment, the processed powder responded to the change in the milling conditions in such a way that the amorphous phase crystallized, to some extent, in the form of Bi4Ti3O12, while decreasing the power during the same milling experiment resulted in complete amorphization. The milling power was modified by changing both the ball-impact energy and the frequency. Therefore, the direction of the nanocrystalline-to-amorphous transition is determined most probably by the ball-impact energy itself since it is reasonable to expect that the change in the frequency can influence at most the kinetics of the transformation but not its course. A detailed study and identification of the intermediate phases formed during a mechanochemical synthesis was carried out by Rojac et al. (2006a), who focused on the reaction between Na2CO3 and Nb2O5. After milling the mixture, they observed a complete loss of the long-range periodicity of the Na2CO3 structure in the first 40 hours, before NaNbO3 was formed. However, this did not occur if the Na2CO3 was milled separately, without the presence of Nb2O5. The results suggested that the two reagents interacted during © Woodhead Publishing Limited, 2010

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milling, resulting in the amorphization of the Na2CO3. The coordination state of the carbonate ions (CO32−) was further examined by infrared spectroscopy. Simultaneously with the amorphization of Na2CO3, a lowering of the point-group symmetry of the CO32− ions was observed, which was attributed to the coordination of these ions into a carbonato complex. The amorphization of the Na2CO3 in the mixture with Nb2O5 was therefore explained in terms of complex formation, during which the CO32− ions must be rearranged from their original structure. The carbonato complex, which was found to be amorphous or eventually nanocrystalline, was considered to be an intermediate phase in the mechanochemical synthesis of NaNbO3. Further studies of the reaction mechanism in the Na2CO3–Nb2O5 system were focused on the second part of the reaction, the formation of NaNbO3 (Rojac et al., 2007, 2008a). In one investigation the formation of the niobate was studied by applying different ball-impact energies (Rojac et al., 2008a). Using a Rietveld refinement analysis complemented by an internal standard method and X-ray line-broadening analysis, the quantitative phase composition (the amounts of Na2CO3, Nb2O5, NaNbO3 and XRD amorphous phase) as well as the crystallite size and the quantity of microstrains in the Nb2O5 and NaNbO3 were determined as a function of the milling time. In addition, the amount of carbonate was followed using thermogravimetric analysis. The results showed that the NaNbO3 forms through an intermediate XRD amorphous phase, which is in agreement with the mechanism proposed by Wang et al. (2000a) and Kuscer et al. (2006, 2007) for systems involving oxides as reagents. The investigation of the mechanochemical synthesis of NaNbO3 by Rojac et al. (2007, 2008a) showed that after prolonged milling a constant ratio between the mass fractions of the NaNbO3 and the XRD amorphous phase was established, which did not change with increasing milling time. At this stage the mixture contained about 50% (in terms of mass fraction) of XRD amorphous phase. This suggested that a quasi-equilibrium condition is reached during the milling, like that observed by Zdujic´ et al. (2006) during the mechanochemical synthesis of Bi4Ti3O12. In a similar manner, the NaNbO3 crystallite size and the quantity of microstrains reached constant values, 17–18 nm and 1.0%, respectively, and did not differ significantly even if the milling experiments were performed under very different milling conditions. In contrast to the work of Wang et al. (2000a), who reported on the growth of PMN nanocrystallites during milling, there was no such tendency observed during the mechanochemical synthesis of NaNbO3. Namely, at low ball-impact energy (15 mJ per hit) a decrease of the perovskite crystallite size was observed with increasing milling time, while it remained constant throughout the experiment at higher ball-impact energy (370 mJ per hit). In addition, the final conversion of Na2CO3 in the mechanochemical reaction with Nb2O5 in terms of its decomposition was found to be lower © Woodhead Publishing Limited, 2010

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(around 70%) than usually obtained in a thermally activated reaction, where a nearly complete decomposition of the Na2CO3 is typically achieved. This suggests that, in terms of thermodynamics, there exists a difference between a mechanochemically induced and a thermally induced reaction. Kosova et al. (1999, 2000, 2001) compared the mechanochemical synthesis of LiMn2O4 starting from mixtures of LiOH–MnO2 or Li2CO3–MnO2. In one of their investigations, the purpose was to identify the intermediate state formed at the very beginning of the mechanochemical interaction between the reagents (Kosova et al., 1999). They observed a considerable difference in the reaction course depending on whether the lithium hydroxide or the carbonate was used in the initial mixture. In the case of the LiOH–MnO2 mixture it was shown that the valence state of the surface manganese ions, determined by X-ray photoelectron spectroscopy (XPS), changed from Mn4+ to a mixed state of Mn3+ and Mn4+ after 10 min activation. At the same time, the Li/Mn concentration ratio on the surface decreased. However, after 10 min of milling this ratio was found to be still 1.6 times higher than in the spinel LiMn2O4, indicating that the surface layer of the particles was enriched with amorphous LiOH. Based on these data, the authors reconstructed the changes occurring on the surface of the LiOH–MnO2 mixture in the following way: during the initial stage of milling LiOH, having a layered structure and exhibiting good plasticity, covers the MnO2 particles, forming an amorphous layer. A similar mechanism, which takes into account the difference in the mechanical properties of the reagents, was also suggested by Streletskii et al. (1995), in their case for the reaction between PbO and Nb2O5. The subsequent chemical interaction is associated with an electron transfer from the OH− ions of LiOH to the Mn4+ ions, partially reducing them to Mn3+ (Equations [6.4], [6.5] and [6.6]). At the same time, the lithium ions diffuse into the interior of the particle, resulting in a decrease of the lithium concentration on the surface. 4LiOH + 8MnO2 → 4Li(Mn3+,Mn4+)O4 + 2H2O + O2 −

[6.4]

−

4OH − 4e → 2H2O + O2

[6.5]

4Mn4+ + 4e− → 4Mn3+

[6.6]

A different picture was observed in the case of the Li2CO3–MnO2 mixture. No reduction of Mn4+ occurred and the Li concentration on the surface of the particles increased during activation from 1 min to 10 min. It was suggested that Li2CO3, being more brittle than LiOH, does not behave in the same way as the hydroxide. The authors proposed that a defect spinel-like phase Li2O⋅yMnO2 (2.5 ≤ y ≤ 4), with a Li content higher than in LiMn2O4, forms on the surface in this case, where the redox stage is absent and only brittle fracture occurs, Equation [6.7]: Li2CO3 + yMnO2 → Li2O⋅yMnO2 + CO2

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The reaction between LiOH and MnO2 was identified as a new, more complex type of reaction. The particularity of the so-called ‘soft mechanochemical reactions’ consists of the high reactivity of surface functional groups, notably, OH groups (Avvakumov et al., 2001). Details of these kinds of reactions are given in Chapter 14. It is well known that mechanochemical reactions can produce nanocrystalline products, for which specific properties are expected. Sˇepelák et al. (2007a) studied the reaction between NiO and α-Fe2O3 giving special attention to the structural particularities of the synthesized NiFe2O4. It was found that the ferrite nanoparticles exhibit a core-shell structure consisting of an ordered inner core surrounded by a disordered surface-shell region with a thickness of about 1 nm (Fig. 6.6). The size of the nanoparticles was estimated to be in the range 6–13 nm. Using Mössbauer spectroscopy, Sˇepelák and co-workers (2007a) showed that the bulk NiFe2O4 exhibits a fully inverse spinel structure with a collinear spin alignment or (Fe↓)[NiFe↑]O4 (for illustration see core in Fig. 6.6). In contrast, the mechanochemically synthesized nanoparticles Shell

Core

t = 1.0 nm

Core: (Fe ) [NiFe ] O4; λc = 1.00; Ψ[A]s = 0°; Ψ[B]s = 0° Shell: (Ni0.33 Fe0.67

) [Ni0.67 Fe1.33

] O4; λs = 0.67; Ψ[A]s = 28°; Ψ[B]s = 40°

6.6 Non-uniform core/shell structure of mechanochemically synthesized nanocrystalline NiFe2O4. The chemical formulae denote the atomic site occupancy and the spin alignment (Ψ[A]s, and Ψ[B]s, where A and B subscripts refer to A and B sites of the spinel structure, respectively) within the particle’s core and shell. λc denotes the degree of inversion defined as the fraction of A sites occupied by Fe3+ ions in 3+ ˇ the spinel structure (M12+−λFeλ3+)(M2+ λ Fe2−λ)O4, where M is Ni (from Sepelák et al., 2007a).

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showed different structural and magnetic features. While the core was considered to possess the same structure and magnetic spin alignment as the bulk ferrite, the shell was shown to be structurally disordered with a nearly random distribution of cations and a canted spin arrangement (see Fig. 6.6). The authors showed that this configuration influences the macroscopic magnetic properties of the ferrite, that is the saturation magnetization and the coercivity. Interestingly, a non-uniform core-shell nanoparticle structure was also observed in high-energy-milled LiNbO3 (Heitjans et al., 2007) and mechanochemically synthesized MgFe2O4 (Sˇepelák et al., 2006), MnFe2O4 (Muroi et al., 2001), PbTiO3 (Szafraniak et al., 2006) and BiFeO3 (Szafraniak et al., 2007).

6.3

Mechanochemical synthesis of complex oxides with various properties

6.3.1 Ferroelectric and related oxides Traditionally, ferroelectric oxides have been synthesized via a conventional solid-state synthesis route, which requires multiple calcinations at elevated temperatures with intermediate milling steps. In addition to the large number of processing steps, a homogeneous powder is difficult to obtain by this method, especially when complex compositions are being synthesized. Furthermore, owing to the coarse particles and low specific surface area, resulting in a poor sinterability of the obtained powders, increased sintering temperatures are required to obtain ferroelectric ceramics with the desired performance. For lead- and alkaline-containing ferroelectrics, for example, high temperatures can lead to an increased loss of highly volatile PbO or alkali oxides giving rise to stoichiometry variations and, eventually, the formation of secondary phases, which have a negative effect on the functional properties of the final ceramics. Because of these problems, extensive research has been done in recent decades to develop alternative synthesis routes. Recently, it has been realized that the main potential of the mechanochemical method lies in its ability to prepare highly homogeneous nanopowders, making it especially attractive for the preparation of materials with a rather complex chemical composition, which is the case for ferroelectric and related oxides. The method has several advantages over other processing routes. First, it uses cost effective and widely available starting compounds (typically, oxides, carbonates or hydroxides) in contrast to the often expensive and moisture-sensitive precursors used in wet chemistry-based methods, such as the sol–gel process. Second, by synthesizing the oxide directly using high-energy milling, the repeated calcination steps, encountered in conventional solid-state synthesis, can be skipped, leading to a simpler process. Furthermore, the mechanochemical reactions are believed © Woodhead Publishing Limited, 2010

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to take place at lower temperatures, which can reduce the loss of volatile components. This is particularly important since most ferroelectric oxides contain lead, bismuth or alkali elements (K, Na and Li). In the past few years, a large number of technologically important complex oxides with different structures (perovskite, Aurivillius, tungsten bronze) and properties (ferroelectric, antiferroelectric, piezoelectric, relaxor, multiferroic) have been prepared either directly by mechanochemical synthesis or by annealing mechanochemically activated powders. These oxides are collected in Table 6.1. Table 6.1 Complex oxides with various structures and properties prepared directly by mechanochemical synthesis or by annealing mechanochemically activated powders Ferroelectric and antiferroelectric oxides with perovskite structure

• BaTiO3, (Ba1−xSrx)TiO3 • PbTiO3, PbZrO3, Pb(ZrxTi1−x)O3, (Pb1−3x/2Lax)TiO3, (Pb1−yLay)(ZrxTi1−x)1−y/4O3, (Pb1−xZnx)(Zr0.53Ti0.47)O3,1 (Pb1−xGdx)(Zr0.53Ti0.47)1−x/4O3,2 (Pb0.97La0.02) (Zr0.62Sn0.31Ti0.04)O3, (Pb0.99Nb0.02)(Zr0.85Sn0.13Ti0.02)0.98O3, (Pb0.99La0.01)(Fe0.01Nb0.02Li0.007Ti0.453Zr0.53)O3.3 • LiNbO3, NaNbO3, (Na0.95Li0.05)NbO3,4 (K0.5Na0.5)NbO3.5 • (Na0.5Bi0.5)TiO3.6

Ferroelectric oxides with Aurivillius structure

• • • • • •

Ferroelectric oxides with tungstenbronze-type structure

• Ba2ANb5O15 (A = K,Na,Li).8

Relaxors and relaxorferroelectrics with perovskite structure

• PMN, PMN–PT, PMN–PZN • PZN, PZN–PT, PZN–BT, PZN–PFW, PZN–PMN–PT, PZN–PFN–PFW • PST

Multiferroic oxides

• BiFeO3,9 Bi(Fe1−xMnx)O3,10 0.7BiFeO3–0.3PbTiO3.11 • Pb(Fe1/2Nb1/2)O3, (Pb1−xBax)(Fe1/2Nb1/2)O3,12 0.8Pb(Fe1/2Nb1/2)O3–0.2Pb(Mg1/2W1/2)O3.13 • Pb(Fe2/3W1/3)O3

Bi4Ti3O12 Bi4Srn−3Ti4O3n+3 (n = 4,5),7 CaBi4Ti4O15 Bi3TiNbO9, (1 − x)Bi2SrNb2O9–xBi3TiNbO9 Bi2MoO6, Bi2(Mo0.25W0.75)O6 Bi2VO5.5 SrBi2Ta2O9

The table was constructed by updating the literature data given in the review of Kong et al. (2008). The reader can refer to the review article for references that are not given in the table. Notations: PMN is Pb(Mg1/3Nb2/3)O3, PZN is Pb(Zn1/3Nb2/3)O3, PST is Pb(Sc1/2Ta1/2) O3, PFN is Pb(Fe1/2Nb1/2)O3, PFW is Pb(Fe2/3W1/3)O3, PT is PbTiO3, BT is BaTiO3. References: 1 Parashar et al., 2005; 2 Parashar et al., 2004; 3 Miclea et al., 2004; 4 Pardo et al., 2004; 5 Rojac et al., 2005b; 6 Van Hal et al., 2001; 7 Ferrer et al., 2005; 8 Van Hal et al., 2001; Khachane et al., 2006; 9 Santos et al., 2006a; Szafraniak et al., 2007; 10 Santos et al., 2006a, 2006b; 11 Khan et al., 2005; 12 Varshney et al., 2007; 13 Levstik et al., 2007.

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The mechanochemical synthesis of a great majority of the oxides listed in Table 6.1 is discussed in the review by Kong et al. (2008). In the following section, some of the main results will be presented for selected materials or groups of materials, that is BaTiO3, Pb(Zr,Ti)O3, alkaline niobates, Bi4Ti3O12, Pb-based relaxor-ferroelectrics and multiferroic BiFeO3. Barium titanate (BaTiO3) is one of the most important materials among the ferroelectric oxides, with a variety of possible applications (Haertling, 1999). A lot of work has been done on its preparation, including mechanochemical synthesis. There are a large number of reports explaining the difficulties associated with synthesizing BaTiO3 directly by milling from a mixture of BaCO3 and TiO2 (Stojanovic´ et al., 1999; Gomez-Yanez et al., 2000; Brzozowski and Castro, 2000, 2003; Berbenni et al., 2001a; Pavlovic´ et al., 2002; Kong et al., 2002a; Goes et al., 2008). As explained by Berbenni et al. (2001a), and in agreement with the results obtained by other researchers (Gomez-Yanez et al., 2000; Brzozowski and Castro, 2000; Kong et al., 2002a; Pavlovic´ et al., 2002), the activation of the BaCO3–TiO2 mixture enhances its reactivity, resulting in a lowering of the temperature of BaTiO3 formation upon subsequent annealing in comparison with the non-activated mixture. However, in all cases, Ba2TiO4 formed as a transitional phase during the heating of the activated powders (Fig. 6.7). The formation of Ba2TiO4 is typically encountered in thermally activated synthesis as a result of the diffusion-controlled reaction between the initially formed BaTiO3 and the BaCO3 (Brzozowski and Castro, 2000; Pavlovic´ et al., 2002). This indicates that under the experimental conditions used in these studies the basic formation mechanism of BaTiO3 upon annealing was not changed by preactivating the BaCO3–TiO2 mixture. However, the amount of secondary phase formed upon annealing can be reduced by increasing the rotational frequency of a planetary mill and/or the activation time (Brzozowski and Castro, 2000). Lead zirconate titanate (Pb(ZrxTi1−x)O3 or PZT) is a technologically important material in electronics and microelectronics, owing to its outstanding ferroelectric, piezoelectric and other properties (Haertling, 1999). The mechanochemical synthesis of nano-sized PZT powders and the fabrication of PZT ceramics from these powders have been reported by several groups using different milling conditions (Lee et al., 1999; Xue et al., 1999b; Kong et al., 2000, 2001a, 2001b, 2002b; Brankovic´ et al., 2003; Beitollahi and Moravej, 2004). A combined approach consisting of mechanochemically activating a co-precipitated precursor has also been proposed (Xue et al., 2000b; Liu et al., 2005). Kong et al. (2001a) combined a mechanochemical synthesis with reaction sintering, using a partially mechanochemically reacted mixture comprising PZT (nominal composition: Pb(Zr0.52Ti0.48)O3), crystalline PbO and an amorphous phase, according to X-ray diffraction (XRD) analysis. Nearly

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(211)

(210)

(111) (d)

(200)

(110)

BaCO3 TiO2 Ba2TiO4 BaTiO3

(100)

Intensity (a.u.)

129

(c) (b) (a) 20

30

40 2θ (deg)

50

60

6.7 XRD patterns of mechanochemically activated BaCO3–TiO2 powder mixture annealed for 2 hours at different temperatures: (a) 600°C, (b) 700°C, (c) 800°C and (d) 900°C (from Kong et al., 2002a). The Miller indices in (d) refer to BaTiO3.

fully dense PZT ceramics (99.2% of theoretical density) sintered at 1100°C for 1 hour were obtained with dielectric and ferroelectric properties comparable to the values reported using other processing techniques. Similar results were also obtained from a mechanochemically synthesized PZT, which did not contain unreacted crystalline oxides (Lee et al., 1999; Xue et al., 1999b). Brankovic´ et al. (2003) synthesized PZT using a planetary mill in a much shorter time (1 hour) in comparison with the results reported by other authors, who typically report on the formation of almost pure PZT after several tens of hours (Kong et al., 2000; Lee et al., 1999; Xue et al., 1999b). The increased rate of PZT formation was attributed mainly to the higher rotational frequency of the disc (317 min−1) and the higher ball-to-powder mass ratio (40 : 1) than was used in other investigations. This could also cause increased contamination during milling, which was not discussed. In addition, the authors observed an interesting phenomenon related to the stability of PZT under high-energy impacts; with prolonged milling the newly formed PZT started to transform into an amorphous phase. However, no explanation was given for this. Alkaline niobates are of a particular interest as they are one of the most promising groups of materials that could replace Pb-containing piezoelectric oxides, such as, for example, PZT. Most of the work relating to

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the mechanochemical synthesis of alkaline niobates has been done on the antiferroelectric NaNbO3. Castro et al. (2004) reported on the mechanochemical synthesis of NaNbO3 using a starting mixture of Na2CO3 and Nb2O5. After 30 days of activation in a vibration mill, a powder mixture with increased reactivity was achieved. Single-phase NaNbO3 was obtained only after subsequent annealing at a lower temperature than usually required by conventional solid-state synthesis, that is 600°C. In order to decrease the activation time, more reactive sodium reagents were used, such as NaOH and Na2O (Hungría et al., 2005). However, because of its high sensitivity to moisture and, therefore, the difficulty of handling Na2O, an off-stoichiometry resulted and a secondary phase was formed upon heating the activated precursor. In addition, they showed that using a planetary mill instead of a vibration mill, which supplies more energy to the reagents, the milling time necessary for the activation of the powder mixture could be reduced to two days. In order to improve the synthesis of NaNbO3 further, a combined wet chemistry and mechanochemical activation technique was also proposed. In this case, single-phase NaNbO3 was prepared by heating the two-days-activated mixture at 700°C. In contrast, Rojac et al. (2005a) managed to prepare NaNbO3 with a crystallite size in the range of a few nanometres to 25 nm directly by a mechanochemical reaction between Na2CO3 and Nb2O5 using a planetary mill. Other studies related to alkaline niobates include the mechanochemical synthesis and activation of LiNbO3 (Figueiredo et al., 1998; Heitjans et al., 2007) and the synthesis of (Na0.95Li0.05)NbO3 (Pardo et al., 2004), which was prepared by annealing a mechanochemically activated powder mixture of alkali carbonates and Nb2O5. To the best of our knowledge there is only one publication on the synthesis of a (K0.5Na0.5)NbO3 (KNN) solid solution, in which the authors report the amorphization of the initial alkali carbonates (K2CO3 and Na2CO3) during milling, while the KNN formation was not realized under the milling conditions applied (Rojac et al., 2005b). A similar behaviour was observed for the synthesis of KNbO3 from a K2CO3–Nb2O5 mixture. Finally, mechanochemical activation was recently applied to the synthesis of a Li- and Ta-modified KNN solid solution (Rojac et al., 2008b). In general, however, poor literature data exist on the mechanochemical synthesis of alkaline niobates and further work is necessary. Bismuth titanate (Bi4Ti3O12) is the most famous Aurivillius-type ferroelectric material. It is distinguished by its high Curie temperature (675°C) and considered a good candidate for high-temperature piezoelectric applications (Subbarao, 1961; Megriche et al., 1999). In the work of Lisoni et al. (2001) it was shown that the mechanochemical activation of a Bi2O3–TiO2 mixture, leading to the formation of an amorphous phase, results in a lowering of the temperature of Bi4Ti3O12 formation upon subsequent annealing

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by about 300°C with respect to the traditional ceramic procedure (850°C). However, upon heating the powders activated in a vibration mill for a total of 72 and 168 hours or in a planetary mill for 19 hours they observed the formation of intermediate phases, that is, Bi12TiO20 and a fluorite-like Bi– Ti–O compound. In contrast, by increasing the activation time in the planetary mill from 19 to 72 hours, apparently, direct crystallization of Bi4Ti3O12, without the appearance of any intermediate phase, took place upon heating the amorphous precursor. This was explained by the increased homogeneity of the mixture, which was achieved by prolonging the activation time. Kong et al. (2001c) showed the feasibility of the synthesis of Bi4Ti3O12 directly by high-energy milling using a planetary mill and tungsten carbide milling media. By sintering the nanopowder, they obtained a dense ceramic (96% of theoretical density) at 750°C within 1 hour. According to the comparison made by the authors, the Bi4Ti3O12 powder obtained exhibited better sinterability than the one prepared by other synthesis methods, such as solid-state reaction and wet chemistry-based processes. Several other studies were published subsequently, reporting on the mechanochemical synthesis of Bi4Ti3O12 and its structural and electrical characterization (Zdujic´ et al., 2006; Stojanovic´ et al., 2006a, 2006b, 2008; Lazarevic´ et al., 2007a, 2007b, 2008). Mechanochemical synthesis also offers the possibility of synthesizing compositions with a rather complex nature, such as, for example, 0.54Pb(Zn1/3Nb2/3)O3–0.36Pb(Mg1/3Nb2/3)O3–0.1PbTiO3 (0.54PZN–0.36PMN– 0.1PT) (Wan et al., 1999). The mechanochemical reaction starting from PbO, ZnO, MgO, Nb2O5, and TiO2 resulted in the formation of the perovskite phase after 20 hours of milling in a shaker mill. The powder exhibited a surface area of 10.9 m2 g−1 and consisted of crystallites that were 10–15 nm in size. A dense single-phase ceramic (>95% of theoretical density) was obtained after sintering at temperatures as low as 900°C, which is considerably lower than those (in the range of 1200°C) generally required for Pb-based relaxors prepared by conventional solid-state synthesis. These powders could be used, for example, in the production of thick films, as proposed by Kosec et al. (2007). They used a mechanochemically synthesized 0.65PMN–0.35PT nanopowder together with the addition of 2 mol% of excess PbO to enable liquid-phase sintering and demonstrated that a homogeneous and single-phase PMN–PT thick film can be obtained after sintering at temperatures as low as 950°C with dielectric, ferroelectric and piezoelectric properties comparable to bulk ceramics (Fig. 6.8). BiFeO3 belongs to a rare group of materials, the multiferroics, in which ferroelectric and magnetic ordering coexist in a single phase. Among the multiferroic oxides, BiFeO3 is distinguished by its high Curie temperature (TC = 830°C) and Néel temperature (TN = 370°C) (Jian and JunHao, 2008). Because of the many problems related to the conventional solid-state

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950°C 900°C 850°C 10

20

30

40 2θ (deg)

50

60

70

(b)

PMN–PT

Pt

10 μm

6.8 (a) XRD patterns of 0.65PMN–0.35PT thick films, prepared using a mechanochemically processed powder, after sintering at 850, 900 and 950°C and (b) cross-section of the thick film on Pt/Al2O3 substrate sintered at 950°C for 2 hours (from Kosec et al., 2007).

synthesis of BiFeO3, such as the formation of undesirable secondary phases and the volatilization of Bi2O3 at higher temperatures, there is a need to develop alternative processing routes. The first steps towards the mechanochemical synthesis of BiFeO3 have already been taken. Santos et al. (2006a) report on the partial amorphization of a Bi2O3–Fe2O3 mixture after 24 hours of high-energy milling in a planetary mill. Annealing the activated mixture

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led to the formation of BiFeO3 together with Bi12Fe4O9 and Bi25FeO40 as secondary phases. In contrast, a successful mechanochemical synthesis of nanocrystalline BiFeO3 by prolonged milling (70 hours) in a shaker mill was achieved by Szafraniak et al. (2007). The impact of the mechanochemical synthesis on the final electrical and magnetic properties of the ferrite remains the subject of future studies.

6.3.2 Magnetic oxides Among various topics in the research on spinel ferrites of type MFe2O4 (M is a divalent metal cation), the magnetic properties of nanoparticles have recently attracted a lot of attention from a fundamental point of view, where an understanding of the basics of nanomagnetism is of major interest, as well as being of considerable practical interest. Because of this, much interest has been shown in mechanochemical synthesis as a way to synthesize ferrite nanoparticles. The idea of using mechanochemistry in order to prepare spinel ferrites dates back to the 1960s, when Kimura (1963) prepared zinc ferrite from ZnO and Fe2O3 by explosive compression. Later, several articles were published on the mechanochemical synthesis of various ferrites, such as ZnFe2O4 (Sˇepelák et al., 1999; Yang et al., 2004a; Verdier et al., 2005; Jean and Nachbaur, 2008), NiFe2O4 (Jovalekic´ et al., 1995; Yang et al., 2004b; Sˇepelák et al., 2005a, 2007a, 2007b), MgFe2O4 (Sˇepelák et al., 2005a, 2006; Bergmann et al., 2006), CoFe2O4 (Shi et al., 2000), MnFe2O4 (Muroi et al., 2001; Osmokrovic´ et al., 2006) and Mn1−xZnxFe2O4 (Arcos et al., 1998; Fatemi et al., 1999; Mozaffari et al., 2008). More data relating to this topic can be found in the review of the mechanochemical synthesis and mechanical activation of ferrites written by Sˇepelák et al. (2005b). Mechanochemical synthesis offers several advantages over traditional processing routes. Besides the low temperature retained during synthesis and the reduced number of processing steps, the method allows the production of nanocrystalline ferrites with particular structural and magnetic properties, as described at the end of Section 6.2. These properties can be tailored by controlling the crystallite size of the obtained mechanosynthesized ferrite, as shown by the example of NiFe2O4 in Fig. 6.9 (Sˇepelák et al., 2005b, 2007a, 2007b). Increasing the crystallite size by annealing at elevated temperatures (up to 1000°C) brings about a gradual increase in the magnetization and a decrease in the coercivity. A magnetic softening is observed and the hysteresis loop saturates at a larger crystallite size. On heating, the metastable mechanosynthesized NiFe2O4 relaxes to a structural and magnetic state similar to that of the bulk. While nanocrystalline mechanosynthesized NiFe2O4 showed a lower magnetization than the bulk ferrite (Fig. 6.9; compare ‘as-prepared’ with ‘1273 K’), the reverse effect was observed for MgFe2O4 (Sˇepelák et al.,

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60

1273 K

40

1073 K

20

973 K As-prepared

0 8.6 nm –20 16.7 nm

T=3K

–40 38.9 nm –60 82.8 nm –5 –4 –3 –2 –1 0

1

2

3

4

5

H (T)

6.9 Magnetization (M)–magnetic field (H) hysteresis loops for mechanosynthesized and subsequently annealed NiFe2O4. The annealing temperature and the corresponding crystallite size, determined by Rietveld refinement analysis, are shown (from ˇ epelák et al., 2007a). S

2006). This was explained in terms of a competition between two effects coming from the disordered shell of the mechanosynthesized nanoparticles (see Fig. 6.6): (1) spin canting, which tends to reduce the magnetization, and (2) site exchange of cations, resulting in an enhanced magnetization. It seems that the two effects play a crucial role in the macroscopic magnetic properties of the ferrite and, interestingly, give different results for NiFe2O4 and MgFe2O4. Recently, besides mechanochemical synthesis, the mechanical activation of spinel ferrites, prepared by a conventional ceramic method, has also become of interest (Sˇepelák et al., 2005b). Similar particularities as in the case of the mechanochemically synthesized ferrites were identified, that is structural disordering reflected in the cation redistribution in the spinel structure and magnetic disordering caused by spin canting (see also Fig. 6.6). Based on the literature data, it should be emphasized that the extent of structural disordering is influenced by the milling conditions applied, since the same nanomaterials from different laboratories exhibited different degrees of cation redistribution (represented by the degree of inversion). In addition, it has also been found that the degree of magnetic disordering (characterized by the spin-canting angle) increases with increasing milling time. These observations seem promising in view of the control of the magnetic properties of ferrites by mechanical activation.

6.3.3 Oxides with semiconducting and catalytic properties LaMnO3 and related perovskites, such as, for example, La1−xSrxMnO3, LaCrO3 and LaCoO3, have been studied primarily for utilization in solid oxide fuel cells (SOFCs) owing to their high electrical conductivity (Minh, © Woodhead Publishing Limited, 2010

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1993) and as catalysts for oxidation and reduction reactions associated with the regulation of waste-gas emissions from vehicles (Voorhoeve et al., 1977). A traditional way to synthesize such materials is by the solid-state reaction from constituent oxides, for which high temperatures are required, typically of the order of 1000°C (Shu et al., 2005; Cheikh-Rouhou Koubaa et al., 2008; Tseggai et al., 2008; Ito et al., 2004). This method results in a large particle size and a limited degree of chemical homogeneity. In order to decrease the synthesis temperature and to obtain fine powders with high specific surface areas, which is important with respect to catalytic applications, mechanochemical synthesis has been considered as a possible preparation method. Zhang and Saito (2000) reported the mechanochemical synthesis of LaMnO3 from a powder mixture of La2O3 and Mn2O3. According to the XRD data, the mechanochemical reaction, yielding LaMnO3, was realized in a planetary mill after just 3 hours of high-energy milling. The final LaMnO3 powder exhibited a specific surface area of 10 m2 g−1. However, the powder consisted of agglomerates with a size up to several micrometers, composed of nanoparticles (<100 nm). Similar results were obtained for the mechanochemical synthesis of La0.7Sr0.3MnO3 (Zhang et al., 2000). The powder also showed in this case strong agglomeration after milling, which prevented an even higher specific surface area being obtained. No methods for de-agglomerating the particles were proposed by the authors. In contrast to the above-presented results, the mechanochemical synthesis of LaCrO3 showed some particularities (Zhang et al., 2002). The synthesis was performed using the same planetary mill and similar milling conditions as in the previous cases for LaMnO3 and La0.7Sr0.3MnO3. The starting materials were La2O3 and different chromium reagents. They were prepared by heating amorphous hydrous chromium hydroxide (Cr(OH)3⋅nH2O) to several temperatures so as to obtain various types of Cr2O3, amorphous or crystalline with different amounts of H2O, depending on the temperature (Fig. 6.10a). According to the XRD data, single-phase LaCrO3 was obtained after 3 hours of milling in the case when amorphous Cr2O3⋅nH2O was used, which was prepared by heating Cr(OH)3⋅nH2O at 200 or 300°C (Fig. 6.10b). On the other hand, little or no reaction was observed when the original Cr(OH)3⋅nH2O or crystalline Cr2O3 (prepared by heating the hydroxide to 400°C, 600°C and 800°C) was used. The low reactivity of the Cr2O3 was explained in terms of its stable corundum-type structure. This point was later confirmed in other systems involving starting compounds with the corundum structure, such as α-Al2O3 and α-Fe2O3 (Zhang et al., 2004). However, no explanations were given for the case of Cr(OH)3⋅nH2O. The LaCrO3 prepared from La2O3 and Cr2O3⋅nH2O showed a lower specific surface area, 5 m2 g−1, than that from the previously reported LaMnO3 (Zhang and Saito, 2000) and La0.7Sr0.3MnO3 (Zhang et al., 2000), presumably owing to the presence of water in the chromium reagent. © Woodhead Publishing Limited, 2010

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(a)

(b) Cr2O3

800°C 600°C Intensity (a.u.)

Intensity (a.u.)

800°C (0.05%)

600°C (0.6%) 400°C (1.6%)

400°C LaCrO3 300°C

200°C

300°C (13.3%) Cr(OH)3·nH2O

200°C (20.3%) 20

40 2θ (deg, CuKα)

60

20

40 2θ (deg, CuKα)

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6.10 (a) XRD patterns of Cr(OH)3⋅nH2O heated to different temperatures and (b) XRD patterns of a mixture of La2O3 and different chromium reagents (see (a)) high-energy milled for 3 hours. In (a) the mass losses upon heating the as-prepared chromium reagents at 850°C for 3 hours are given (in wt%) and are related to the H2O content (for the original Cr(OH)3⋅nH2O, the H2O release upon heating was determined to be 43.7%) (from Zhang et al., 2002).

Sato et al. (2006) studied the mechanochemical synthesis of LaMnO3+δ from La2O3 and Mn3O4 under different humidity conditions. The reaction proceeded towards the formation of LaMnO3+δ only by milling in air with a relative humidity (RH) of 70%, while no reaction occurred after milling for 30 min under dry conditions (RH = 0.1%) (Fig. 6.11). This showed that the formation of La(OH)3 from La2O3, which was observed during milling in humid air (Fig. 6.11a), plays a crucial role in the synthesis of LaMnO3+δ. Further investigations of the milling experiment performed in humid conditions (RH = 70%) revealed an increase of the specific surface area of the powder mixture occurring simultaneously with the formation of La(OH)3; based on this observation, it was assumed that more efficient grinding took place due to the lower mechanical hardness of the hydroxide in comparison with La2O3. However, LaMnO3+δ did not form if the reaction was carried out starting from pure La(OH)3 and Mn3O4; in this case only amorphization of the mixture took place. During the milling of a La(OH)3–Mn3O4 mixture, a considerable decrease in the specific surface area was observed (from 10.5 to 4 m2 g−1 in 30 min of milling), suggesting strong agglomeration, which could influence the course of the reaction. The authors proposed that the agglomeration of the particles was caused

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(a) 30 min

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6.11 XRD patterns of La2O3–Mn3O4 mixture high-energy milled under (a) 70% and (b) 0.1% relative humidity (from Sato et al., 2006).

by the formation of oxygen bridging bonds through an acid/base reaction between the hydroxyl groups; however, no experimental evidence was given to support this conclusion (Ohara et al., 2008b). In general, the problem of contamination during milling is treated insufficiently in the literature. Bolarín et al. (2007) managed to synthesize LaMnO3 in a shaker mill and examined the level of contamination. The concentration of Fe coming from the hardened-steel milling balls and the inner walls of the vial increased with increasing milling time. A considerable amount of Fe was introduced into the powder after 20 hours of milling, 3.4% (in terms of mass fraction). In contrast, using an attrition-type mill (Fig. 6.3b), which enables LaMnO3+δ to be synthesized in a short time (30 min; Fig. 6.11a) without employing balls as the milling media, a much lower level of contamination can be preserved (<20 ppm Fe) (Sato et al., 2006). As an alternative to a mechanochemical synthesis from a mixture of oxides, Ito et al. (2004) proposed synthesizing LaCoO3 via a mechanochemical method from chlorides. Milling a LaCl3–CoCl2–NaOH mixture in a molar ratio of 1 : 1 : 6 for 3 hours in a planetary mill triggered a substitution reaction in which La(OH)3, Co(OH)2 and NaCl were formed. LaCoO3 was obtained after subsequent annealing at a temperature of 600°C, which is

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significantly lower than normally needed for a conventional synthesis from constituent oxides (930°C). In any case, the method requires an additional step, since the NaCl has to be removed after calcination by washing with water. The specific surface area of the final LaCoO3 prepared from the chloride method was 10 m2 g−1, which is three times larger than LaCoO3 obtained from a mechanochemical reaction with La2O3 and Co2O3, presumably owing to the fine dispersion of the hydroxides formed in the NaCl matrix after milling. More examples of the synthesis of perovskites, such as La1−xCaxMnO3+δ and La1−xSrxCoO3, by applying mechanochemical activation together with their catalytic activities can be found in Isupova et al. (2001a, 2001b). ZnFe2O4 has been studied because of its potential application as an absorber in the high-temperature desulphurization of coal gas (Ayala and Marsh, 1991). Both mechanical activation (high-energy milling of ZnFe2O4) and mechanochemical synthesis (preparation of ZnFe2O4 by high-energy milling) were considered as possible ways to produce nanocrystalline ZnFe2O4 with increased reactivity and, consequently, improved catalytic properties. Sˇepelák et al. (1997) showed that the mechanochemically synthesized ferrite exhibited a higher sulphur-absorption capacity, even with respect to the mechanically activated ZnFe2O4 (Fig. 6.12). The authors emphasized that the desulphurization tests were performed at a temperature at which the thermal relaxation of the mechanically induced defects had already taken place, that is at 400°C. Therefore, in addition to the high specific surface area, an important factor influencing the high-temperature reactivity of the mechanosynthesized ferrite arises from an increase in the ion mobility, which is connected to the thermally stimulated disappearance of the mechanically induced structural disorder. However, this also means

Mechanosynthesized ZnFe2O4 Unmilled mixture ZnO/α-Fe2O3 ZnFe2O4 milled for 18 min ZnFe2O4 milled for 8 min Unmilled ZnFe2O4 0

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6.12 Sulphur absorption capacity of various Zn–Fe–O sorbents ˇ epelák et al., 1997). (from S

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that the advantageous properties of the nanosized oxide are mostly lost after relaxation from a metastable (activated) state to a stable one. In this case, an understanding of the thermal stability and the relaxation mechanisms of mechanically induced metastable phases is essential and, unavoidably, further studies will be needed in the future. Other reports on the application of the mechanochemical method in the synthesis of ferrites possessing catalytic properties, such as LaFeO3, La1−xSrxFeO3−y, NiFe2O4 and Ca2Fe2O5, can be found in the literature (Zhang and Saito, 2001; Yakovleva et al., 2004; Manova et al., 2005; Isupova et al., 2002).

6.3.4 Miscellaneous Other oxides with dielectric, ionic conductive, optical and other properties that were synthesized by mechanochemical means can be found in the literature: CaTiO3 (Brankovic´ et al., 2007), SrTiO3 (Berbenni et al., 2001b), CaCu3Ti4O12 (Almeida et al., 2002), SrFeO2.5 (Schmidt and Kaczmarek, 1999), Bi3MO7 (M = Nb, Ta) (Castro and Palem, 2002), Y3Al5O12 (GuO and Sakurai, 2000), BaAl2Si2O8 (Bosˇkovic´ et al., 1999) and A2Ti2O7 (A = Y, Gd, Dy) (Fuentes et al., 2005), to give some examples. An extensive review of the mechanochemical synthesis of many oxides can be found in the work of Zyryanov (2008). Furthermore, the synthesis of materials based on lithium manganese oxide for Li-ion-battery applications was not included in this chapter. For this topic refer to Chapters 14 and 15 (part III). Recently, high-energy milling has also been applied in the preparation of ZnO-based varistor ceramics (Gomez-Yanez et al., 2004; Zˇunic´ et al., 2007; Liu et al., 2007; Bernik et al., 2008).

6.4

Future trends

Owing to an increased interest in the synthesis of complex ceramic oxides by the mechanochemical method, significant progress has been made in terms of the understanding of the basic mechanisms of mechanochemical reactions as well as in the preparation of several technologically important oxides. In our opinion there are some important points that should be explored in more detail in the future. First, when cross-checking the literature, one can see that a variety of milling machines are used to process the powders, from planetary, vibration, attrition and shaker mills, as examples of ball mills, to specially designed mills operating without employing any balls as the milling media (Fig. 6.3b). Irrespective of the type of mill employed, they basically differ in terms of the extent of the impact and the shear forces acting on the powders, which are responsible for the numerous phenomena occurring

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during the mechanochemical process. Presently, there are still several discrepancies between the results obtained by different laboratories using different or even the same milling equipment, as was illustrated by some of the examples in this chapter. In order to enable a comparison between them, further studies on a quantitative basis are necessary. The influence of the milling parameters, which are often described insufficiently in the literature, as well as of the mechanics of the different types of mills should be addressed systematically for various reaction systems. In relation to this, the effect of different parameters, such as the rotational frequency and powder load, on the extent of contamination from the milling media should also be studied in more detail, as it is an important subject that has not been sufficiently explored yet. The second point concerns the application of mechanochemically processed powders. There are many examples in the literature showing the benefit of using such powders in the preparation of bulk ceramics owing to their specific properties, such as, for example, nanocrystallinity and high homogeneity (see Section 6.3.1). However, other possibilities in the application of such powders have not been studied sufficiently. In the field of electronic ceramics, for example, there is a continuous interest in the development of thick films driven by the trend for miniaturizing electronic devices. Here, lowering the sintering temperature is important to avoid any undesirable chemical reactions between the active layer and the substrate. Mechanochemical synthesis could be an appropriate choice as it is capable of producing homogeneous nanopowders with improved sinterability, which could then be applied in the preparation of thick films by screen printing. The first attempts have already been made in this direction (Stojanovic´ et al., 2002; Hu et al., 2004, 2005; Jain and Patil, 2006; Jain et al., 2007; Kosec et al., 2007) and can be used as a starting point for future, more systematic studies. Even more challenging is to apply the mechanochemically synthesized nanoparticles to thin-film processing, as shown by Brankovic´ et al. (2004). However, as the thickness of the films reduces, well-dispersed nanoparticles would be necessary for the deposition process. In this case, one should think about the de-agglomeration of mechanochemically synthesized powders and/or the minimization of particle agglomeration during high-energy milling, to which more attention should be paid in the future. The third problem relates to contamination during milling. Presently, there are poor literature data on this subject. The first step is to establish clearly the relationship between the milling parameters and contamination from the milling media for various oxide systems. In addition to this, more attention should be also paid to the selection of appropriate material, from which the milling media is made, and on the influence of contamination, in view of the final functional properties of ceramics.

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rojac t, bencˇan a, ursˇicˇ h, malicˇ b and kosec m (2008b), ‘Synthesis of a Li- and Ta-modified (K,Na)NbO3 solid solution by mechanochemical activation’, J Am Ceram Soc, 11, 3789–3791. santos i a, cótica l f, de medeiros s n, paesano jr a, coelho a a, gama s, venet m z, garcia d and eiras j a (2006a), ‘Structural, microstructural and magnetic properties of the high-energy ball milled BiFeO3 and BiFe0.95Mn0.05O3 ferroelectromagnetic compounds’, Ferroelectrics, 338, 233–9. santos i a, grande h l c, freitas v f, de medeiros s n, paesano jr a, cótica l f and radanovic´ e (2006b), ‘Structural, microstructural and Mössbauer spectral study of the BiFe1−xMnxO3 mechanosynthesized system’, J Non-Cryst Solids, 352(32–5), 3721–4. sato k, chaichanawong j, abe h and naito m (2006), ‘Mechanochemical synthesis of LaMnO3+δ fine powder assisted with water vapor’, Mater Lett, 60(11), 1399–402. schmidt m and kaczmarek w a (1999), ‘Synthesis of SrFeO2.5 from mechanically activated reactants’, J Alloys Compd, 283(1), 117–21. sˇ epelák v, steinike u, uecker d c, trettin r, wißmann s and becker k d (1997), ‘High-temperature reactivity of mechanosynthesized zinc ferrite’, Solid State Ionics, 101–103, 1343–9. sˇ epelák v, wißmann s and becker k d (1999), ‘Magnetism of nanostructured mechanically activated and mechanosynthesized spinel ferrites, J Magn Magn Mater, 203, 135–7. sˇ epelák v, bergmann i, feldhoff a, litterst f j and becker k d (2005a), ‘Mechanosynthesis of spinel ferrite nanoparticles followed by Mössbauer spectroscopy’, Hyperfine Interact, 165, 81–7. sˇ epelák v, bergmann i, kipp s and becker k d (2005b), ‘Nanocrystalline ferrites prepared by mechanical activation and mechanosynthesis’, Z Anorg Allg Chem, 631(6–7), 993–1003. sˇ epelák v, feldhoff a, heitjans p, krumeich f, menzel d, litterst f j, bergmann i and becker k d (2006), ‘Nonequilibrium cation distribution, canted spin arrangement, and enhanced magnetization in nanosized MgFe2O4 prepared by a one-step mechanochemical route, Chem Mater, 18(3), 3057–67. sˇ epelák v, bergmann i, feldhoff a, heitjans p, krumeich f, menzel d, litterst f j, campbell s j and becker k d (2007a), ‘Nanocrystalline nickel ferrite, NiFe2O4: Mechanosynthesis, nonequilibrium cation distribution, canted spin arrangement, and magnetic behavior, J Phys Chem, 111(13), 5026–33. sˇ epelák v, heitjans p and becker k d (2007b), ‘Nanoscale spinel ferrites prepared by mechanochemical route: Thermal stability and size dependent magnetic properties’, J Therm Anal Calorim, 90(1), 93–7. shi y, ding j and yin h (2000), ‘CoFe2O4 nanoparticles prepared by the mechanochemical method’, J Alloys Compd, 308(1–2), 290–5. shu q f, zhang j y, liu j h and zhang m (2005), ‘Solid-state reaction for preparation of lanthanum manganite’, High Temp Mater Processes, 24(3), 169–73. stojanovic´ b d, pavlovic´ v b, pavlovic´ v p, djuric´ s, marinkovic´ b a and ristic´ m m (1999), ‘Dielectric properties of barium-titanate sintered from tribophysically activated powders’, J Eur Ceram Soc, 19(6–7), 1081–3. stojanovic´ b d, foschini c r, pavlovic´ v p, pavlovic´ v m, pejovic´ v b and varela j a (2002), ‘Barium titanate screen-printed thick films’, Ceram Int, 28(3), 293–8.

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stojanovic´ b d, simoes a z, paiva-santos c o, jovalekic´ cˇ, mitic´ v v and varela j a (2005), ‘Mechanochemical synthesis of barium titanate’, J Eur Ceram Soc, 25(12), 1985–9. stojanovic´ b d, paiva-santos c o, jovalekic´ cˇ, simoes a z, filho f m, lazarevic´ z and varela j a (2006a), ‘Mechanically activating formation of layered structured bismuth titanate’, Mater Chem Phys, 96(2–3), 471–6. stojanovic´ b d, simoes a z, paiva-santos c o, quinelato c, longo e and varela j a (2006b), ‘Effect of processing route on the phase formation and properties of Bi4Ti3O12 ceramics’, Ceram Int, 32(6), 707–12. stojanovic´ b d, paiva-santos c o, cilense m, jovalekic´ cˇ and lazarevic´ z zˇ (2008), ‘Structure study of Bi4Ti3O12 produced via mechanochemically assisted synthesis’, Mater Res Bull, 43(7), 1743–53. streletskii a n, borunova a b, shiryaev p a and shashkin d p (1995), ‘Kinetics and mechanism of mechanochemical synthesis in the PbO–Nb2O5 system’, Russ J Phys Chem, 69(6), 887–92. subbarao e c (1961), ‘Ferroelectricity in Bi4Ti3O12 and its solid solutions’, Phys Rev, 122(3), 804–7. szafraniak i, połomska m, hilczer b, talik e and kepin´ski l (2006), ‘Characterization of PbTiO3 nanopowders obtained by room temperature synthesis’, Ferroelectrics, 336, 279–87. szafraniak i, połomska m, hilczer b, pietraszko a and kepin´ski l (2007), ‘Characterization of BiFeO3 nanopowders obtained by mechanochemical synthesis’, J Eur Ceram Soc, 27(13–15), 4399–402. tseggai m, nordblad p, tellgren r, rundlöf h, andrè g and bourèe f (2008), ‘Synthesis, nuclear structure, and magnetic properties of LaCr1−yMnyO3 (y = 0, 0.1, 0.2, and 0.3)’, J Alloys Compd, 457(1–2), 532–40. van hal h a m, groen w a, maassen s and keur w c (2001), ‘Mechanochemical synthesis of BaTiO3, Bi0.5Na0.5TiO3 and Ba2NaNb5O15 dielectric ceramics’, J Eur Ceram Soc, 21(10–11), 1689–92. varshney d, choudhary r n p, rinaldi c and katiyar r s (2007), ‘Dielectric dispersion and magnetic properties of Ba-modified Pb(Fe1/2Nb1/2)O3’, Appl Phys A, 89(3), 793–8. verdier t, nachbaur v and jean m (2005), ‘Mechanosynthesis of zinc ferrite in hardened steel vials: Influence of ZnO on the appearance of Fe(II)’, J Solid State Chem, 178(11), 3243–50. voorhoeve r j h, johnson d w, remeika j p and gallagher p k (1977), ‘Perovskite oxides: Materials science in catalysis’, Science, 195(4281), 827–33. wan d m, xue j m and wang j (1999), ‘Synthesis of single phase 0.9Pb [(Zn0.6Mg0.4)1/3Nb2/3O3]–0.1PbTiO3 mechanically activating mixed oxides’, Acta Mater, 74(7), 2283–91. wang j, junmin x, dongmei w and weibeng n (1999), ‘Mechanochemical fabrication of single phase PMN of perovskite structure’, Solid State Ionics, 124(3), 271–9. wang j, xue j m, wan d m and gan b k (2000a), ‘Mechanically activating nucleation and growth of complex perovskites’, J Solid State Chem, 154(2), 321–8. wang j, xue j m and wan d (2000b), ‘How different is mechanical activation from thermal activation? A case study with PZN and PZN-based relaxors’, Solid State Ionics, 127(1–2), 169–75.

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welham n j (1998), ‘Mechanically induced reaction between alkaline earth metal oxides and TiO2’, J Mater Res, 13(6), 1607–13. xue j, wan d and wang j (1998), ‘Mechanochemical synthesis of nanosized lead titanate powders from mixed oxides’, Mater Lett, 39(6), 364–9. xue j, tan y, wan d and wang j (1999a), ‘Synthesizing 0.9PZN–0.1BT by mechanically activating mixed oxides’, Solid State Ionics, 120(1), 183–8. xue j, wan d, lee s e and wang j (1999b), ‘Mechanochemical synthesis of lead zirconate titanate from mixed oxides’, J Am Ceram Soc, 82(7), 1687–92. xue j, wang j and wan d (2000a), ‘Nanosized barium titanate powder by mechanical activation’, J Am Ceram Soc, 83(1), 232–4. xue j, wang j and weiseng t (2000b), ‘Synthesis of lead zirconate titanate from an amorphous precursor by mechanical activation’, J Alloys Compd, 308(1–2), 139–46. xue j m, wan d m and wang j (2002), ‘Functional ceramics of nanocrystallinity by mechanical activation’, Solid State Ionics, 151(1–4), 403–12. yakovleva i s, isupova l a, tsybulya s v, chernysh a v, boldyreva n n, alikina g m and sadykov v a (2004), ‘Mechanochemical synthesis and reactivity of La1−xSrxFeO3−y perovskites (0 ≤ x ≤ 1)’, J Mater Sci, 39(16–17), 5517–21. yang h, zhang x, huang c, yang w and qiu g (2004a), ‘Synthesis of ZnFe2O4 nanocrystallites by mechanochemical reaction’, J Phys Chem Solids, 65(7), 1329–32. yang h, zhang x, ao w and qiu g (2004b), ‘Formation of NiFe2O4 nanoparticles by mechanochemical reaction’, Mater Res Bull, 39(6), 833–7. zdujic´ m, poleti d, jovalekic´ cˇ and karanovic´ l (2006), ‘The evolution of structure induced by intensive milling in the system 2Bi2O3⋅3TiO2’, J Non-Cryst Solids, 352(28–9), 3058–68. zhang q and saito f (2000), ‘Mechanochemical synthesis of LaMnO3 from La2O3 and Mn2O3 powders’, J Alloys Compd, 297(1–2), 99–103. zhang q and saito f (2001), ‘Effect of Fe2O3 crystallite size on its mechanochemical reaction with La2O3 to form LaFeO3’, J Mater Sci, 36(9), 2287–90. zhang q, nakagawa t and saito f (2000), ‘Mechanochemical synthesis of La0.7Sr0.3MnO3 by grinding constituent oxides’, J Alloys Compd, 308(1–2), 121–5. zhang q, lu j and saito f (2002), ‘Mechanochemical synthesis of LaCrO3 by grinding constituent oxides’, Powder Technol, 122(2–3), 145–9. zhang q, lu j, wang j and saito f (2004), ‘Dependency of mechanochemical reactions forming complex oxides on the crystal structures of starting oxides’, J Mater Sci, 39(16–17), 5527–30. zˇunic´ m, brankovic´ z, bernik s, goes m s and brankovic´ g (2007), ‘ZnO varistors from intesively milled powders’, J Eur Ceram Soc, 27(13–15), 3897–900. zyryanov v v (2008), ‘Mechanochemical synthesis of complex oxides’, Russ Chem Rev, 77(2), 105–35.

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7 Production of intermetallic compound powders by a mechanochemical approach: solid–liquid reaction ball milling C H E N D I N G, J I A N G YO N G, C A I J I A N - G U O and C H E N Z H E N - H UA, Hunan University, People’s Republic of China; and H UA N G P E I - Y U N, Central South University, People’s Republic of China

Abstract: A novel mechanochemical approach, known as solid–liquid reaction ball milling, was developed in order to prepare solid state intermetallic compound powders by the reaction of the milling mediums with a liquid state metal over a particular temperature range. Numerous ultrafine elevated temperature intermetallic compound powders were obtained by this newly developed process. In contrast to the conventional mechanical milling technique, nanometer-sized elevated temperature intermetallic powders and nearly pure intermetallic phases can be formed directly at relatively low temperatures. This short review will introduce the main results and reaction mechanism for this mechanochemical approach as well as its advantages and potential applications. Key words: mechanochemical approach, solid–liquid reaction ball milling, intermetallic compound powders, reaction mechanism, nanoparticles.

7.1

Introduction

Since the term mechanochemistry was first proposed by W. Ostwald and added to the types of reactions initiated by mechanical force, it has been widely used in the fields of materials synthesis, the chemical industry, extractive metallurgy, minerals engineering, bioengineering and so on (Gilman, 1996; Fernandez-Bertran, 1999; Bálaz, 2000, 2008; Todres, 2006; Boldyrev, 2006; Chen and Chen, 2008). Many methods have been used to initiate or improve mechanochemical effects. The ball-milling process was considered to be one of the most effective approaches for causing and accelerating various chemistry reactions (Murty and Ranganathan, 1998; Lu and Lai, 1998; Soni, 1999; El-Eskandarany, 2001; Suryanarayana, 2001a, b; Takacs, 2002; Zhang, 2004; Chen and Chen, 2006). The terms ‘mechanical alloying’ (MA) and/or ‘reaction milling’ (RM) were used for this mechanochemical 149 © Woodhead Publishing Limited, 2010

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process in the field of advanced material synthesis, although they are not precise descriptions. Intermetallic compounds, a type of metallic material, have attracted intense study by materials scientists and engineers. Their ordered nature leads to them having attractive elevated temperature properties such as high strength, increased stiffness and excellent corrosion/oxidation resistance. However, their low ambient temperature ductility and fracture toughness preclude their large-scale industrial application. Over the past few decades, the ball milling/mechanical alloying process has proved to be a successful method of improving the room temperature ductility of intermetallics, because it can reduce the grain size, disorder the lattice and modify the crystal structure of the phase into a more symmetrical one (Koch and Whittenberger, 1996; Koch, 1998; Morris, 1998; Koch et al., 1999). Thus, the mechanochemical approach (reaction milling) became an important way to obtain high-performance intermetallic compound powders. Although the chemical reactions of intermetallic synthesis caused by solid–solid and solid–gas reactions have been very widely reported, there have been few investigations and reports about solid–liquid reactions (SLRs) (Lu and Lai, 1998; Soni, 1999; El-Eskandarany, 2001; Suryanarayana, 2001a, b; Chen and Chen, 2008). The most probable reason for this is that almost all metals (except for mercury) and alloys are in a solid state at room temperature and external effects such as magnetic, electric and ultrasonic fields on the solidification behaviour of metal melts have attracted much attention in the last half century or so (Flemings, 1974; Lane et al., 1960; Asai, 1978; Campbell, 1981; Misra, 1985; Abramov, 1987; Asai, 1989; Nakada et al., 1990; Uhlmann et al., 1966; Conrad, 2000). In the process of milling a melt, the portion of melt in direct contact with the balls is under high pressure, which enhances atomic reactivity, and this can induce solid–liquid mechanochemical reactions. It has been recognized that these kinds of reaction can occur at room temperature or at a lower temperature than that normally required for the preparation of pure metal particles, nanocomposite powders and other useful functional materials in material chemistry and related fields (Chen and Chen, 2008). However, much more work is still required to achieve the synthesis of high-purity elevated temperature intermetallic compounds by the solid–liquid mechanochemical reaction method. A novel solid–liquid reaction (SLR) ball milling technique, based on the combined effects of mechanochemistry and thermochemistry, has recently been developed by our group (Chen and Chen, 2005, 2006; Chen et al., 2004a, b, 2007a, b, c, 2008a, b, 2009). Numerous binary and ternary intermetallic compound powders have been successfully obtained and most of them were single-phase powders and nanoparticles. Compared with tradi-

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tional mechanical alloying (high-energy ball milling) processes, the advantages of this mechanochemical approach are a finer particle size, higher purity of production, a faster reactivity speed and a different micro-reaction mechanism. So far, binary intermetallic powders of X–Fe (X = Sn, Pb, Zn, Sb) and Al–X (X = Ti, Ni, Fe, Cu) systems such as Fe–Sn, Fe3Sn2, Ti3Al, Ni2Al3, Fe3Al, and ternary intermetallics of Al–Fe–X (X = Cu, Si), Al– Cu–X (X = Co, Ni) alloy systems such as Al7Cu2Fe, Al13Cu4Fe3, Al65Cu20Fe15, Al8Fe2Si, Al3FeSi, Al65Co15Cu20, Al69Co25Cu6, Al17Cu4Ni and Al0.28Cu0.69Ni, Cu–Sn and Cu–Ni–Sn have been successfully prepared. In this chapter, the equipment and methods used in the experiment are briefly described, together with the main results obtained and the reaction mechanism of the solid–liquid reaction ball milling technology. The advantages and potential applications of this mechanochemical approach will also be discussed.

7.2

Experimental equipment and methods

The solid–liquid ball milling equipment developed by the authors is shown in Fig. 7.1 (Chen and Chen, 2006), in which a sealed milling cylinder with a diameter of 300 mm rotates in a resistance-heated furnace equipped with a thermostatic system. Both the balls and the milling cylinder (or inner liner of the milling cylinder) are made of the same starting material to prevent contamination of the melt. The milling cylinder was first evacuated and then filled with pure inert gas to avoid oxidation of the metals during milling at the chosen

2 1

3

4

5

6

7

8

7.1 Schematic diagram of solid–liquid reaction ball milling equipment: 1, thermoelectric couple; 2, thermostatic system; 3, electric furnace; 4, milling cylinder; 5, vacuum valve; 6, shaft coupling; 7, motor; 8, milling ball.

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temperature. The metal with the higher melting point is called the start material, while the metal with the lower melting point is the reactant. The weight ratio of the balls to the reactant was 20 : 1 to 30 : 1 for different alloy systems and the rotational speed was kept at 80 rpm. Here, it must be pointed out that the balls need to be replaced periodically because of their reduction during the milling process. The milling temperature was chosen to be higher than the melting point of the reactant (about 50 K) and therefore it was in a molten state during the milling process. For comparison, a powdery metal with the same composition as the starting material was also added to investigate its effect on the phase formation process during the solid–liquid reaction ball milling.

7.3

As-milled products obtained from solid–liquid reaction ball milling

7.3.1 Binary alloy systems The as-milled products of the binary alloy systems that were synthesized by solid–liquid reaction milling under different conditions are summarized in Table 7.1. According to the X-ray diffraction results obtained, an almost single phase was successfully obtained from the binary intermetallic phase for various systems (except for the Ti–Al and Ni–Al systems) by this mechanochemical approach, a result which is generally obtained by ingot metallurgy at relatively high temperatures. Some details of the results are given for different milling-medium systems. Table 7.1 As-milled products of the binary alloy systems for solid–liquid reaction ball milling Milling medium

Melting metals/alloys

Reactive production

Fe

Sn Zn Sb Al

FeSn2, FeSn, Fe3Sn2, Fe1.3Sn FeZn13, FeZn10.98, FeZn8.87, Fe11Zn40 FeSb2, FeSb FeAl2 + Fe2Al5 (mixture)

Ni

Al Sn Sb Zn

Ni2Al3 + NiAl3 (mixture) Ni3Sn2, Ni3Sn NiSb, Ni5Sb2 Ni5Zn21, NiZn

Cu

Sn Al

Cu6Sn5, Cu3Sn Cu2Al

Ti

Al

TiAl3

W

Al

W12Al, W4Al, W5Al

Mo

Al

MoAl5, Mo4Al17

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Fe–X (X = Sn, Zn, Sb, Al) systems As one of the most important industrial metal elements, Fe was a prime candidate for the milling medium and solid reactant, and Sn, Zn, Sb and Al were chosen as the liquid reactants for their low melting points. For the Fe–Sn system, the FeSn2, FeSn, Fe1.3Sn and Fe3Sn single-phased intermetallic powders were synthesized after milling for 48 hours, 64 hours, 72 hours and 72 hours at 597 K, 873 K, 1033 K and 1073 K, respectively (Chen et al., 2004a). Of these intermetallic compound phases, Fe1.3Sn is a high-temperature phase and thus its formation needs a higher milling temperature and a much longer milling time than the low-temperature phase FeSn2. However, it was obtained at a relatively lower temperature, by about 200 K, in our experiment than by the traditional ingot metallurgy method. If this were not the case, it would be extremely difficult to control the process parameters to obtain a single phase in the casting process. All the as-prepared powders were characteristically layered and agglomerated with an overall particle size less than 100 nm in diameter, demonstrated by typical scanning electron microscope (SEM) and transmission electron microscope (TEM) images of Fe1.3Sn intermetallic compound particles, as shown in Fig. 7.2. Furthermore, FeZn13, FeZn10.89, FeZn8.9 and Fe11Zn40 were formed at 723 K, 823 K, 923 K and 993 K, respectively, after milling for 48 hours in the Fe–Zn system. FeSb2 and FeSb were obtained after milling for 48 hours at 923 K and 1043 K, respectively, in the Fe–Sb system (Chen et al., 2007a). However, only a two-phase mixture of Fe2Al5 and FeAl2 was formed after milling at 973 K for 48 hours for the Fe–Al system in our experiment (Chen et al., 2008a). In addition, the results show that the addition of powdery iron to the molten metal expanded the surface area and hence accelerated the reaction rate. For instance, the addition of 20 wt% Fe powder (below 74 μm) halved the time needed to obtain the same milled product with the same milling parameters (Chen et al., 2004a, 2007a). Ni–X (X = Sn, Sb, Zn, Al) systems Ni is another element that was used as a grinding medium in our research. From the Ni–Sn system, pure Ni3Sn and Ni3Sn2 were obtained after milling for 48 hours at 573 K and 1073 K, respectively. According to the XRD data of the as-obtained powders from the Ni–Sb system, NiSb and Ni5Sb2 powders were produced after milling for 24 hours at 923 and 1023 K, respectively. From the Ni–Zn system, Ni5Zn21 and NiZn powders were produced after milling for 24 hours at 723 K and 923 K, respectively (Chen and Chen, 2006). From the Ni–Al system, as with the Fe–Al system, the as-milled

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100 nm (b)

50 nm

7.2 (a) SEM and (b) TEM images of the Fe1.3Sn powder particle obtained by solid–liquid reaction ball milling.

product was a two-phase mixture of Ni2Al3 + NiAl3 after 12 hours at 973 K (Chen et al., 2008a).

Other binary alloy systems Apart from Fe and Ni, Cu, Ti, W and Mo were used as the milling medium (milling ball and inner liner of the milling cylinder). Thus, a series of singlephase intermetallic compounds were obtained after different milling times and at different temperatures, that is CuAl2 from the Cu–Al system, Cu6Sn5 and Cu3Sn from the Cu–Sn system (Chen et al., 2007c), TiAl3 from the Ti–Al

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system (Chen et al., 2008a), W12Al, W4Al and W5Al from the W–Al system (Chen et al., 2008b) and MoAl5 and Mo4Al17 from the Mo–Al system (Chen et al., 2009). In addition, it must noted that although most of the binary intermetallic compound powders in Table 7.1 can usually be obtained by mechanical alloying and other chemical/metallurgy approaches, phase-pure products are difficult to achieve and the particle size and morphology are essentially uncontrollable. On the other hand, nanoparticle and single-phase intermetallic compound powders can easily be obtained by solid–liquid reaction ball milling at a relatively low cost. Moreover, some of the binary intermetallic compound powders listed in Table 7.1 have not been listed in existing reference material as being produced by the conventional mechanical alloying process, such as Fe1.3Sn, W12Al, Mo4Al17. Other intermetallic compound powders, such as TiAl3 W4Al, MoAl5, are not obtained directly from mechanical alloying and need an additional heat-treatment process (Zhang et al., 2000; Zhu et al., 2004, 2005; Zdujic´ et al., 1994; Monagheddu et al., 1999).

7.3.2 Ternary alloy systems Al–Fe–X (X = Cu and Si) systems In the Al–Fe–X system, the iron in the balls and the milling cylinder acted as the start material and pure binary alloy ingots, that is Al–33.2wt%Cu, Al–54wt%Cu, Al–7wt%Si (hypoeutectic composition), Al–12.6wt%Si (eutectic composition) and Al–30wt%Si were used as the reactants, respectively (Chen et al., 2004b). Completely solid state Al13Cu4Fe3 powders were obtained after milling the Al–33.2wt%Cu melt for 48 hours, while Al7Cu2Fe was obtained after milling for 12 hours when 20 wt% iron powder was added into the melt. Al13Cu4Fe3 and Al65Cu20Fe15 phases were obtained after milling the Al– 45wt%Cu melt blended with 20 wt% iron powder at 943 K for 12 hours and 48 hours, respectively. Obviously, the iron powder also accelerated the reaction rate. Al8Fe2Si was the final as-milled product from milling the Al–7wt%Si melt blended with 20 wt% iron powder, or without the iron powder addition for a different period of time, at 973 K. The mixture of Al8Fe2Si and Al3FeSi was obtained after milling the Al–12.6wt%Si alloy melt blended with the iron powder for 48 hours at 963 K. Iron powder not only accelerated the formation rate of Al8Fe2Si, but also contributed to the formation of Al3FeSi during a further solid state high-temperature milling process. For the Al– 30wt%Si melt, Al3FeSi was finally obtained after milling the Al–30wt%Si alloy melt blended with the iron powder for 24 hours at 1133 K. Iron

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powder also promoted the reaction rate in a similar way to the Al–Cu–Fe system. Using a similar approach, Al65Co15Cu20, Al69Co25Cu6, Al0.28Cu0.69Ni0.02 and Al7Cu4Ni ternary alloy powders were obtained after milling for different times at different temperatures by milling the Al–Cu melt with a Co and Ni medium, respectively (Chen et al., 2007b). The TEM images of the asmilled products for the Al–Cu–Co and Al–Cu–Ni systems from solid–liquid reaction milling are shown in Fig. 7.3, which shows that these intermetallic

(a)

50 nm (b)

50 nm

7.3 TEM images of ternary alloy powders obtained by solid–liquid reaction ball. (a) Al–Cu–Co system, (b) Al–Cu–Ni system.

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powders are composed of agglomerated particles with an average grain size of less than 100 nm.

7.3.3 Effect of process control agents on solid–liquid reaction ball milling A process control agent (PCA) is usually used in the process of mechanical alloying to avoid cold welding and bonding between the powder particles and the balls and the agglomeration of powder during milling, which is also noticeable for some alloy systems in the solid–liquid reaction ball milling process (Chen and Chen, 2006). Thus, an investigation into the effect of PCAs on solid–liquid reaction ball milling was necessary and important. Some results of the effect of PCAs on the solid–liquid reaction ball milling process are given here using the Ni–Al alloy system as an example. PCAs can be solid, liquid or gas and most of them are organic compounds. The usual addition of PCA is 1–5% (Suryanarayana, 2001a). Based on previous research results, NaCl, KCl, Na2SO4 and sodium stearate were selected as PCAs to reduce bonding between the balls and the melt during solid–liquid reaction milling (Chen and Chen, 2006; Chen et al., 2008a). It should be noted that the addition of NaCl, KCl, Na2SO4 and sodium stearate into the Ni–Al system can effectively reduce coarsening and agglomeration. Owing to the volatility of inorganic substances at high temperatures and the corrosive effect of saline material on an aluminum melt, choosing the right PCAs for the solid–liquid reaction milling is important. The XRD and SEM results indicate that different PCAs showed different anti-adhesive properties during milling and the same PCA even showed different effects at different concentrations. The grain size, powder output and purity of the product depend on the properties of the PCA. The specific results are listed and shown in Table 7.2 and Fig. 7.4. Compared with the Table 7.2 XRD results for as-milled products when different PCAs were added Concentration PCA

1 wt%

2 wt%

4 wt%

NaCl

Ni2Al3, NiAl3

Ni2Al3, NiAl3

KCl

Ni2Al3, NiAl3

Na2SO4

Ni2Al3, NiAl3

Ni2Al3, NiAl3, NiAl3– AlCl3 (impurity) Ni2Al3, NiAl3

Sodium stearate

Ni2Al3, NiAl3 (higher formed speed)

Ni2Al3, NiAl3, AlCl3 (impurity) Ni2Al3, NiAl3, NiAl3– AlCl3 (impurity) Ni2Al3, NiAl3, some impurity peaks Ni2Al3, NiAl3

Ni2Al3, NiAl3

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1 μm (b)

1 μm

7.4 SEM photographs of milled product with addition of different PCAs immersed in (a) 1%HCl solution, (b) 1% stearic acid, (c) 2% stearic acid and (d) 2% NaCl.

milling products without the addition of PCA, it can also be concluded that the right PCAs are not only able to decrease bonding and agglomeration, but can also accelerate the reaction to some extent (Chen et al., 2008a).

7.4

Reaction mechanism of solid–liquid reaction ball milling

7.4.1 Reaction mechanism Traditionally, intermetallics have been prepared by ingot metallurgy technology, which requires high-temperature processing, and powder metal-

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(c)

1 μm (d)

1 μm

7.4 Continued

lurgy, which also requires high-temperature processing or complex sintering processing. But during the solid–liquid reaction milling process, pure intermetallics such as Fe1.3Sn, FeSb2, TiAl3, Al13Cu4Fe3, Al8Fe2Si, and so on, were successfully prepared at temperatures much lower than their respective melting points. With regard to mechanochemical effects, ball milling can cause crystal lattice distortions, crystal defects and electron emissions, resulting in improved chemical activity and an increased reaction rate. As a result, the activity of the starting material, that is iron during the solid–liquid reaction milling process, is high enough to accelerate the reaction and shorten the milling time for the formation of pure intermetallics. Additionally, collisions between the milling balls lead to a local temperature increase above the

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melting point of the ternary alloy, which is conducive to its formation. Moreover, the final ternary alloy product has the same elemental molar ratio characteristic as its corresponding original binary alloy. The reaction milling process shows distinct differences from conventional mechanical alloying such as high-energy ball milling and tumbler milling. The former involves a direct reaction, while the latter involves repeated flattening, cold-welding, fracturing and rewelding and a solid state diffusion process at the atomic scale of the powders. Consequently, solid–liquid reaction ball milling has a higher reaction rate than mechanical alloying. For example, some investigation results show that the intermetallic compound phase of FeAl appeared after high-energy ball milling for 25–30 hours (Matsumoto et al., 2000) while this can be obtained using solid–liquid reaction ball milling after 6 hours. Furthermore, when alloy systems with intermetallic phases in their phase diagrams are treated using high-energy ball milling, in most cases an amorphous phase, rather than an intermetallic, is obtained after milling for a long time; intermetallics form directly in relatively few alloy systems (Chen and Chen, 2006). In other words, the solid– liquid reaction milling process can result in an intensive solid–liquid reaction. In the next section, a more detailed comparison between solid– liquid reaction ball milling and mechanical alloying (high-energy ball milling) will be given. One characteristic of our novel solid–liquid reaction milling technique is that it can produce nano-sized intermetallic particles. The formation mechanism is explained in Fig. 7.5. During this process, the newly formed reaction

Ball

Melt

Powder

Ball Collision

Reaction product Solid–liquid reaction Fresh surface Reaction product peels off from the balls and breaks into fine particles

Intermetalllic compound powder Melt

7.5 Formation mechanism for a nano-size intermetallic compound powder by solid–liquid reaction ball milling.

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product on the surface of the balls peels off and breaks into fine particles during the continuous collisions between the balls. The fresh surface helps to accelerate the solid–liquid reaction rate to form an intermetallic layer. These cycles will not stop until the melt is consumed. The added metal powders provide more reaction surface area to accelerate the reaction rate further. The particles formed are deposited in the melt in a solid state. These nanoparticle powders will obviously not amalgamate or increase in size owing to continuous impact from the balls. Finally, obtaining high-temperature intermetallic compound phases at a relatively lower temperature has a particular advantage over the usual casting process. It is well known that in the casting process, the melting temperature must be higher than the melting points of all the alloy elements, in other words, the temperature where the intermetallic compound phases form must be higher than the highest melting point of the alloy elements. However, in the solid–liquid reaction ball milling process, the temperature at which intermetallic compound phases are obtained only needs to be higher than the lowest melting point of the alloy elements. Obviously, the latter approach is simpler, more convenient and less costly.

7.4.2 Comparisons with conventional mechanical alloying (high-energy ball milling) In this section, a detailed comparison is made of the milling mechanisms used in solid–liquid reaction ball milling and conventional mechanical alloying (high-energy ball milling). First, solid–liquid reaction ball milling utilizes the combined effects of the mechanical force field and the temperature field and induces a chemical reaction between the solid ball milling medium and the liquid metal melt. Mechanical alloying involves lengthy intense impacting of the powder particles and the balls, after which the powders flatten, fracture and cold-weld; these steps are repeated until the alloy is produced. Some studies have shown that local intense impacting during the mechanical alloying process produced high temperature and pressure, resulting in partial melting and a chemical reaction. But this reaction is thermopositive and alloying occurs through the self-propagating reaction. The present studies show that this is only one model of mechanical alloying mechanisms, not the universal and essential principle. It mainly applies to some metal–oxide systems, and the constituents of the reactant, the reaction type, the intensity of the milling, the reaction thermodynamics, the additive component and the microstructure, and so on, all have an effect on it. The mechanism for the formation of mechanical alloying is based on the ball-milling process: the impact of the mechanical force, the intense plastic deformation of the powder particles and the large amount of dislocations

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and sub-grains generated in the particles. The whole system is in an unsteady state, according to minimum energy theory, and in every reaction system, the corresponding product is the alloy phase, amorphous phase, nano-crystal, quasi-crystal and so on. But this mechanism does not exist in the solid– liquid reaction ball-milling process, as the deformation of grains and the increase in dislocations only occur in the solid phase. There are no dislocations and sub-grains in the alloy melt. So the mechanisms of the two processes have to be essentially different. In addition, the chemical reaction of the solid–liquid reaction ball-milling process is different to that of mechanical alloying. 1. Based on the description above, the diffusion rate of the solid–liquid reaction is much (often several orders of magnitude) larger than that of the solid–solid reaction. So the reaction rate of solid–liquid reaction ball milling is not analogous to that of mechanical alloying, which has already been proved by experiment. 2. According to the reaction model suggested by the authors, intermetallic compound powders are produced by a solid–liquid reaction from the beginning of the ball milling process and what follows is only the accumulation of the production of the reaction until the liquid state is fully exhausted. But during mechanical alloying, the chemical reaction process must have an induction period, which is a period of accumulation of the reaction energy, before the mechanical alloying begins and intermetallic compound powders are obtained. 3. Similarly, based on the reaction model mentioned in Section 7.4.1, the products of the solid phase are often finer or nano-size compound powders, because of the small area of the solid–liquid reaction and the effect of the impact from continuous ball milling. One feature of mechanical alloying is the production of ultrafine powders and most powders are nanocrystalline. However, nano-sized powders are difficult to produce directly through solid state mechanical alloying. 4. In comparison with mechanical alloying, the temperature field of the solid–liquid reaction ball milling process decreases the reaction activation energy significantly. The solid–liquid reaction ball milling process induces some reactions and produces some alloys and compounds which could not be obtained during the mechanical alloying process. 5. As seen from the experimental results of solid–liquid reaction ball milling and high-energy ball milling (commonly called the MA process), the products of the two are not exactly the same. For example, most of the products of the Fe system are not the phases obtained during mechanical alloying, except for FeSn2. But FeSn2, FeSn, Fe1.3Sn and Fe3Sn2 were obtained through solid–liquid reaction ball milling. In the W–Al and Mo–Al systems, W12Al, W4Al, Mo4Al17, W5Al and MoAl5

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were obtained, while the products of mechanical alloying were some amorphous or oversaturated solid solution powders, under the same experimental conditions. The products of the Al–Cu–Fe and Al–Si–Fe systems were ternary intermetallic powders, but the results of mechanical alloying were amorphous and oversaturated solid solution powders. The above results obtained from mechanical alloying are similar to those obtained by others (Tsai et al., 1993; Barua et al., 2001). In their research, the quasi-crystal phase and amorphous phase were obtained by the mechanical alloying process, but the intermetallic compound phase was not produced by ball milling. Obviously, the difference in products is further proof of the different mechanisms of these two milling processes. 6. Finally, although the solid–liquid reaction ball milling process and mechanical alloying are both based on the mechanical force chemical reaction theory, they are quite different. The mechanical force chemical reaction theory was not used at its birth to explain the chemical reactions that happen in the mechanical alloying process, because the appearance of this theory and the establishment of the subject both occurred half a century earlier than the mechanical alloying process. Therefore it was used to explain some of the experimental results found during mechanical alloying, while solid–liquid reaction ball milling is a new process based on this theory.

7.4.3 Main advantages of solid–liquid reaction ball milling compared to mechanical alloying Following on from the discussions above, the two main advantages of solid– liquid reaction ball milling compared to mechanical alloying are given below. First, the reaction rate and the productivity of solid–liquid reaction ball milling are high. The rotation speed used by the current authors for solid– liquid reaction ball milling was 80 rpm and the ball-to-reactant ratio was 30 : 1. After 48 hours milling time, 250 g of FeSn2 powder were produced with an average diameter of 100 nm. But when using a high-energy planetary ball milling machine as the device for mechanical alloying, with the same ball-to-reactant ratio and a rotation speed of 400 rpm, 120 hours were needed to obtain 100 g of FeSn2 powder and the powder was not as fine as that from solid–liquid reaction ball milling. As a result, solid–liquid reaction ball milling has better prospects for industrial use than the mechanical alloying process. Second, compared with the mechanical force field of mechanical alloying, the temperature field induced during milling caused the effects of fields external to the material to be stronger, so the ability to produce new alloy

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compounds increased. Solid–liquid reaction ball milling provides a new method of producing and studying new materials.

7.5

Conclusions

As a new approach to preparing intermetallic compound powders, solid– liquid reaction ball milling shows some promising applications for extending the industrial scale application of intermetallic compounds. In fact, the industrialization of this technology has been carried out in some companies in China. Moreover, if some intermetallic compound phases exist at certain high temperatures or some unusual phase transformation processes occur at high temperatures, the solid–liquid reaction ball milling process may be able to help in obtaining these high-temperature intermetallic compound phases and phase transformation processes. Furthermore, some new alloy phases, which cannot be obtained by traditional means, may be discovered and produced in this way, which may be beneficial in understanding some dark or questionable regions (especially in some high temperature range) of the existing phase diagram. However, some limitations of this new ball milling method can be foreseen. Improvements need to be made in order to control the reaction process and production more precisely and the detailed micro-reaction mechanism for particular alloy systems is not very clear as yet.

7.6

Acknowledgement

The authors are grateful to the National Natural Science Foundation of China (Project 50574039) and the Doctorate Fund of the Education Ministry of China (20050532004, 20070532016) for their support of the present research. Chen Ding gratefully acknowledges the helpful discussion and important information received from Prof. Dr Yan Hong-ge (Hunan University, Changsha) and Prof. Dr Chen Gang (Hunan University, Changsha).

7.7

References

abramov o v (1987), ‘Action of high-intensity ultrasound on solidifying metal’, Ultrasonic, 25, 73–82. asai s (1978), ‘Effects of electromagnetic forces on solidified structure of metal’, ISIJ International, 18, 754–60. asai s (1989), ‘Birth and recent activities of electromagnetic processing of materials’, ISIJ International, 29, 981–92. bálaz p (2000), Extractive Metallurgy of Activated Minerals, Elsevier Press, Oxford. bálaz p (2008), Mechanochemistry in Nanoscience and Minerals Engineering, Springer Press, Berlin. barua p, murty b s and srinivas v (2001), ‘Mechanical alloying of Al–Cu–Fe elemental powders’, Materials Science and Engineering A, 304–6, 863–6.

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boldyrev v v (2006), ‘Mechanochemistry and mechanical activation of solids’, Uspekhi Khimii, 75, 203–16. campbell j (1981), ‘Effects of vibration during solidification’, International Materials Review, 26, 71–104. chen z h and chen d (2005), ‘Preparation of binary and ternary intermetallic powders via a novel reaction ball milling technique’, Journal of Metastable and Nanocrystalline Materials, 24/25, 719–22. chen z h and chen d (2006), Mechanical Alloying and Solid–liquid Reaction Ball Milling, Chemistry Industry Press, Beijing. chen d and chen z h (2008), Mechanochemistry, Chemistry Industry Press, Beijing. chen z h, chen d, chen g, yan h g and huang p y (2004a), ‘Preparation of elevatedtemperature intermetallic powders via a novel reaction ball milling technique’, Journal of Alloys and Compounds, 370, 43–6. chen d, chen z h, chen j h and huang p y (2004b), ‘Preparation of the Al–Cu–Fe and Al–Fe–Si ternary intermetallic powders via a novel reaction ball milling technique’, Journal of Alloys and Compounds, 376, 89–94. chen d, chen j h, yan h g and chen z h (2007a), ‘Synthesis of binary and ternary intermetallic powders via a novel reaction ball milling technique’, Materials Science and Engineering A, 444, 1–5. chen d, cai j g, wu w, chen g and yan h g (2007b), ‘Preparation of the Al–Cu–Co and Al–Cu–Ni ternary intermetallic powders via a solid–liquid reaction ball milling technique’, Material Science Forum, 561–5, 363–6. chen d, wu w, chen z h, fu d f and chen g (2007c), ‘Preparation for intermetallic powders Cu–Sn and Cu–Ni–Sn systems via solid–liquid reaction milling technology’, Transactions of Nonferrous Metals Society of China, 17s1, 594–8. chen d, chen g l, ni s, chen g, yan h g and chen z h (2008a), ‘Phase formation regularities of ultrafine TiAl, NiAl and FeAl intermetallic compound powders during solid–liquid reaction milling’, Journal of Alloys and Compounds, 457, 292–5. chen d, chen z, cai j g and chen z h (2008b), ‘Preparation of W–Al intermetallic compound powders by a mechanochemical approach’, Journal of Alloys and Compounds, 461, L23–25. chen d, cai j g, fan j j and chen z h (2009), ‘Preparation of Al–Mo intermetallic powders by solid–liquid reaction ball milling’, Journal of Alloys and Compounds, 485, L9–L11. conrad h (2000), ‘Influence of an electric or magnetic field on the liquid–solid transformation in materials and on the microstructure of the solid’, Materials Science and Engineering A, 287, 205–12. el-eskandarany m s (2001), Mechanical Alloying for Fabrication of Advanced Engineering Materials, Noyes Publications/William Andrew Publishing, New York. fernandez-bertran j f (1999), ‘Mechanochemistry: an overview’, Pure and Applied Chemistry, 71, 581–6. flemings m c (1974), Solidification processing, McGraw-Hill, New York. gilman j j (1996), ‘Mechanochemistry’, Science, 274(5284), 65. koch c c (1998), ‘Intermetallic matrix composites prepared by mechanical alloying – a review’, Materials Science and Engineering A, 244, 39–48. koch c c and whittenberger j d (1996), ‘Mechanical milling/alloying of intermetallic’, Intermetallics, 4, 339–55. koch c c, morris d g, lu k and inoue a (1999), ‘Ductility of nanostructured materials’, MRS Bulletin, 24, 54–8.

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lane d h, cunningham j w and tiller w a (1960), ‘The application of ultrasonic energy to ingot solidification’, Transactions of The Metallurgical Society of AIME, 218, 985–94. lu l and lai m o (1998), Mechanical Alloying, MA Kluwer Academic Publishers, Hingham. matsumoto a, kobayashi k, nishio t, ozaki k and sugiyama a (2000), ‘Microstructure and mechanical properties of FeAl compacts prepared by pulse current sintering of mechanically alloyed powders’, Journal of the Japan Society of Powder and Powder Metallurgy, 47, 1253–7. misra a k (1985), ‘A novel solidification technique of metals and alloys – under the influence of applied potential’, Metallurgical Transactions A, 16, 1354–5. monagheddu m, delogu f, schiffini l, frattini r and enzo s (1999), ‘Local icosahedral symmetries in Al–Mo alloys prepared by ball milling’, NanoStructured Materials, 11, 1253–61. morris d g (1998), Mechanical Behavior of Nanostructured Materials, Trans Tech Publications, Zurich, Switzerland. murty b s and ranganathan s (1998), ‘Novel materials synthesis by mechanical alloying / milling’, International Materials Reviews, 43, 101–41. nakada m, shiohara y and flemmings m c (1990), ‘Modification of solidification structures by pulse electric discharging’, ISIJ International, 30, 27–33. soni p r (1999), Mechanical Alloying: Fundamentals and Applications, Cambridge International Science Publishing, Cambridge. suryanarayana c (2001a), ‘Mechanical alloying and milling’, Progress in Materials Science, 46, 1–181. suryanarayana c (2001b), Mechanical alloying and Milling, Elsevier Press, Oxford. takacs l (2002), ‘Self-sustaining reactions induced by ball milling’, Progress in Materials Science, 47, 355–416. todres z v (2006), Organic Mechanochemistry and its Practical Application, CRC Press, Boca Raton. tsai a p, tsurui t and memezawa a (1993), ‘Formation of quasi-crystals through peritectoid reactions from a nonequilibrium cubic phase prepared by mechanical alloying’, Philosophical Magazine Letters, 67, 393–8. uhlmann d r, seward t p and chalmers b (1966), ‘Effect of magnetic fields on structure of metal alloy castings’, Transactions of The Metallurgical Society of AIME, 236, 527–31. zdujic´ m, poleti d, karanovi l, kobayashi k f and shingu p h (1994), ‘Intermetallic phases produced by the heat-treatment of mechanically alloyed Al–Mo powders’, Materials Science and Engineering A, 185, 77–86. zhang d l (2004), ‘Processing of advanced materials using high-energy mechanical milling’, Progress in Materials Science, 49, 537–60. zhang f, lu l and lai m o (2000), ‘Study of thermal stability of mechanically alloyed Ti–75% Al powders’, Journal of Alloys and Compounds, 297, 211–18. zhu c j, ma x f, zhao w, tang h g, yan j m and cai s g (2004), ‘Processing, microstructure and mechanical properties of W50Al50 bulk alloy obtained by mechanical alloying and hot-pressing’, Scripta Materialia, 51, 993–7. zhu c j, ma x f, zhao w, tang h g, yan j m and cai s g (2005), ‘Synthesis and thermal stability of Al75W25 alloy obtained by mechanically alloying’, Journal of Alloys and Compounds, 393, 248–51.

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8 Mechanochemical processing of non-oxide systems with highly covalent bonds M A Ł G O R Z ATA S O P I C K A - L I Z E R, Silesian University of Technology, Poland

Abstract: This chapter discusses mechanochemical processing in nonoxide compounds with covalent bonds and explains how the technique could overcome processing complexity in the system of interest. Two types of processing are described, in situ and ex situ. The former is limited to mechanochemically induced exothermal synthesis of powder particles from elements, followed, if necessary by self-sustaining reaction. The latter is applied to as-synthesized non-oxide powders in order to improve their reactivity/sinterability or to stimulate the reaction during the subsequent heat treatment or consolidation. Titanium or silicon carbides and nitrides, sialon (nano)powder synthesis and processing by high-energy milling are described. Key words: carbides, nitrides, silicon, titanium, self-propagating synthesis, carbothermal reduction, mechanochemical process.

8.1

Introduction

There is a growing demand for ultra-wear-resistant and tough structural refractory materials in applications such as high-speed cutting tools, hard disc drivers, engine parts and so on and structural non-oxide materials such as nitrides, carbides and silicides are of special interest because of their exceptional hardness, extraordinary corrosion resistance, low density and high melting point. Covalent bonds between the elements in groups IV and V in the periodic table (C, Si, N) and transition metals in groups III, IV and V (B, Si, Ti, Zr, Mo, W) bring exceptional resistance to deformation under applied stress, that is high elastic modulus and outermost hardness. However, the resulting covalent bond response of the crystal to applied stress is modified by its crystal structure, that is the density and arrangement of the packing of its atoms. In addition, it is just this covalent bonding nature of the compounds in question that directly results in inherent brittleness, complex processing and difficulties in machining. There are several approaches for modifying the microstructure of these structural ceramic products through careful microstructural design of the (nano)composites or laminates, combining 167 © Woodhead Publishing Limited, 2010

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the properties that result from the crystal structure, the most interesting feature of nanocrystallinity, but obstacles to successful and financially acceptable manufacturing routes still remain. Despite the enormous progress that has been made in the processing of monolithic engineering ceramics in the last two to three decades in the field of densification, microstructural control and grain boundary refinement, reduction of the grain sizes in the final product to the nanometer level is still a challenge. The primary fabrication route for bulk advanced ceramic products involves manufacturing a green compact followed by densification at high temperature by pressure/pressureless sintering, hotpressing (HP), gas pressure sintering (GPS), hot isostatic pressing (HIP) or, recently, spark plasma sintering (SPS). Thus a nanometer-sized or nanocrystalline/amorphous powder that has superior quality in terms of nanocrystallinity, the ability to be compacted and is free of agglomerates becomes of a special interest. Nano-sized powder is therefore essential in manufacturing nanostructured ceramic products. Moreover, nano-sized particles shorten the diffusion distance during densification of the covalent bonded compound or could be applied for reinforcing particulate composites, manufacturing nanostructured protective wear resistant coatings and so on. The major drawbacks of gas phase-derived nanopowders are their high agglomeration, low packing density and the enormous problems that arise if homogeneous mixing with a minor additive is required. Additionally, synthesis of non-oxide compounds in the form of nanocrystalline powders encounters a problem of the high melting temperature that makes formation of a coarse-crystalline powder unavoidable. Subsequent milling of these powders suffers from low efficiency and heavy contamination because of the great hardness of the powder particles involved. Therefore, mechanochemical processing (MCP) appears to be an attractive and commercially promising method of direct synthesis of these compounds, of reducing the reaction temperature of the prealloyed composition or of the effective manufacture of nanocomposites in the form of nanostructured particles. This chapter describes the results of mechanochemical processing those covalently bonded compounds that turn out to be resistant to the technique because of their inherent hardness and oxidation ability and thus only mills with high energy density can be applied and a protective atmosphere is necessary for processing. The chapter is divided into two main sections. The first devoted to the direct synthesis of the compound from the elements accompanied by exothermal reactions (in situ reactions). The second is dedicated to a variety of reactions in the non-oxide systems which require substantial activation energy to form carbides, nitrides, borides, silicides or nanocomposite particles. Frequently this type of process requires subse-

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Table 8.1 Basic properties of non-oxide compounds

Compound

Density (g cm−3)

Elastic modulus (GPa)

Hardness HV (GPa)

Melting temperature (°C)

Si3N4 (HP) β-sialon SiC TiC WC B4C Ti3SiC2 MoSi2

3.29 3.24 3.21 4.92 15.8 2.25 4.55 6.31

210–310 280–290 400–450 497* 620–720 450–470 320–400 441–439*

14–16 13–15 18–24 28–34 17–24 29–35 8–11 13

1900 2500 3070 2780 2445 2930 2030

* Refers to monocrystals.

quent thermal treatment (ex situ reaction). The primary choice of relevant compounds to be included in this chapter was their intrinsic resistance to applied external stress, although it was modified by the industrial significance of the compounds and extraordinary features that it is possible to achieve by MCP. The most important properties of these compounds/polycrystalline ceramics are given in Table 8.1.

8.2

Manufacturing non-oxide powders by mechanochemical processing

The combustion process is assumed to be an effective powder processing technique capable of synthesizing a variety of equilibrium and non-equilibrium phases directly from blended elements. This type of direct high exothermic synthesis is described as self-propagating high temperature synthesis (SHS), and sometimes referred to as combustion synthesis (CS) (Munir and Tamburini, 1989). It utilizes the heat generated by an exothermic reaction, sustaining the reaction along a combustion wave after external ignition, although the feasibility of SHS is limited to reaction systems with an adiabatic temperature exceeding 1500°C. The combined effect of MCP and CS leads to a process called mechanically activated combustion synthesis (MACS) or a mechanically induced self-sustaining reaction (MSR). Alternatively, in some systems, the kinetics of metal combustion is fast and the combustion temperature in the combustion wave is much higher than the relevant metal melting point. As a result, melting and coalescence of metal powder particles can inhibit compound formation as well as inducing rapid grain growth at high temperature. The preliminary mechanochemical processing of the blended elements can improve the reaction activity of the materials and affect the combustion synthesis. Details of

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these combined reactions in various systems are discussed below for nitrides, carbides and silicides.

8.2.1 Nitrides Silicon nitride and its large family of solid solutions containing the principal element (Si, Al, O and N (sialons))-based ceramics have earned a premier place among engineering ceramics because of their superior properties and excellent performance in applications such as cutting tools, ball bearings and seals (Oyama, 1971; Jack and Wilson, 1972). Silicon nitride is a highly covalent bonded compound with a hexagonal structure and two basic crystallographic modifications designated as α and β (Grun, 1979; Kohatsu and McCauley, 1974). Silicon nitride powders are usually produced via direct nitridation, carbothermal reduction and vapour phase reactions or by chemical wet methods. Direct formation of nitride powder particles from its elements implies the heterogeneous reaction of metal powder particles in a nitrogen-rich atmosphere. However the attraction of this type of reaction is weakened by the very low self-diffusion coefficient of nitrogen atoms throughout the newly formed nitride product layer thus involving a long reaction time. It was obvious that a high-energy milling could overcome this problem since repeated collisions between the milling media and metal particles with the newly formed nitride film would destroy the film and remove the nitride product, thus making available fresh surface for subsequent reaction with nitrogen. The first reports showed that mechanical activation of silicon by ball milling leads to formation of nanocrystalline or amorphous particles accompanied by an increase in the free energy of 19.9 kJ mol−1, when the volume fraction of amorphous Si reaches 15% (Shen et al., 1995). The author’s own studies showed that ball milling of silicon powder with alumina (rotational speed of 250 rpm, ball-to-powder (BP) ratio 20 : 1) brought about a reduction of silicon crystallite size to below 100 nm after 4 hours milling in argon. Significant progress has been shown to have been achieved after examination of the reaction milling of silicon or boron in ammonia (Nikolov et al., 1995). Combustion analysis of the milled powders (200 hours dry ammonia under pressure of 600 kPa in a specially designed Uni-Ball-Mill (Calka and Radlinski, 1991)) showed a substantial amount of acquired nitrogen in the milled silicon powder (32 wt% N2 and 1.07 wt% H2) but besides Fe contamination (7 wt%) from the milling media no crystalline phase was detected. However, subsequent annealing in argon at 800°C for 2 hours resulted in the crystallization of α-Si3N4. Too lengthy ball milling and serious contamination from the milling media did nothing to enhance research into manufacturing silicon nitride

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powders by MCP until a new idea was suggested to combine the combustion synthesis (CS) with the mechanochemical pretreatment of the powders of interest. The CS of silicon nitride suffers from heavily agglomerated coarse powders if the high nitrogen pressure of several megapascals combined with significant dilution of the starting silicon powder by silicon nitride is not applied (Kata et al., 1998). Moreover, the product of CS appears mostly as β-Si3N4, while α modification is of commercial interest. This hindrance was resolved by the use of CS with MCP of the initial silicon powder together with the addition of ammonium salts (NH4Cl or NH4F) that are able to enhance the nitridation reaction (Cano et al., 2001). Consequently, α-Si3N4 powder was manufactured by the application of a nitrogen pressure as low as 1 MPa (Jin et al., 2006). Since an addition of ammonia salts clearly facilitates transition of silicon to gaseous phase, thus reducing the nitridation temperature to 1000°C was observed at this pressure (Zakorzhevskii and Borovinskaya, 2000). The clear effect of MCP on combustion synthesis was shown by the measurement of combustion temperature in relation to the milling time, as shown in Fig. 8.1 (Jin et al., 2008). Further studies showed, however, an interesting feature of the mechanochemically induced effects in amorphous silicon-based powder where the path of the subsequent crystallization in CS depends on the extent of the applied pressure, that is the duration of the milling process or altitude of nitrogen pressure during CS (Jin et al., 2006, 2008). The longer milling time or higher nitrogen pressure promotes formation of β-Si3N4 in contrast to mild conditions where α-Si3N4 was favoured. A similar phenomenon was observed after high-energy milling of boron in ammonia followed by annealing the resultant powder at 1000°C. It has been reported that only

Temperature (°C)

1800 1780 1760 1740 1720 1700

0

5

10 15 Milling time (h)

20

25

8.1 Maximum combustion temperature at nitrogen pressure of 6 MPa of silicon-based mixture obtained after ball milling for different times, recorded by a thermocouple (after Jin et al., 2008, with kind permission of Elsevier).

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low-energy milling limited to shearing actions resulted in subsequent crystallization of the amorphous powder to h-BN (Nikolov et al., 1995). MCP-assised combustion synthesis could also be utilized in the manufacture of sialon powders as there is a great demand for a good quality nanocrystalline sialon powder. In contrast to the commonly operating highpressure reaction sintering of the mixture of the relevant nitrides and oxides, pressureless fabrication of β- or α/β-sialon ceramics directly from the powders with a controlled sialon solid solution could reduce production costs and broaden the range of application. First reports in this field demonstrate the advantages of direct combustion synthesis of β- or α-sialon powders from their elements if the blended batches were mechanically activated first. High-energy milling of mixtures of silicon, aluminium and alumina for only 18 min enabled a necessary dilutent in the CS process to be eliminated and the nitrogen pressure to be reduced to 1 MPa and improved the yield of the resultant β-sialon powder in a broad range of solid solutions, as evidenced by Sivakumar et al. (2007). This technique offers pure β-sialon powder with a designed composition but rather coarse particles because of the sudden rise in temperature during synthesis. Application of planetary milling to the additive-free resultant product and SPS for powder consolidation gave fully dense specimens after heat treatment at 1600°C for 10 min (Sivakumar et al., 2009). It is expected that the limited content of glassy phase in the manufactured product will raise the hightemperature properties of the material. Preparation of α-sialon powder by MCP is a challenging task as the nitrogen-to-oxygen ratio strictly limits the range of solid solubility in the α-Si3N4 crystal lattice and precise design of the cation concentration is necessary (Ekström, 1989). The first attempt to produce Y-α-sialon powder by combustion synthesis in air preceded by high-energy milling was demonstrated by Liu et al. (2007). A mixture of silicon and aluminium powder with addition of yttria and silica, diluted by 3 wt% α-Si3N4 powder and enriched by 3 wt% stearic acid was mechanically activated in a vibratory mill (BP = 15 : 1) for 15–120 min with zirconia balls in a sealed jar. The combustion synthesis was performed in a graphite die after triggering the combustion wave at the one end of the sample and was finished after approximately 8 min. Despite significant particle size reduction below 500 nm after MCP, no changes in the reactant crystallinity were evidenced by X-ray diffraction (XRD) but the CS was possible only in the specimens with the longest milling time. Moreover, the resultant phase composition was dependent on their position along the combustion wave, starting from oxides in the top through Y-α-sialon in the centre as well as non-reacted silicon. It is supposed that after the initial oxidation of metal particles at the launch of the combustion wave, the subsequent drop in the oxygen partial pressure inside the specimen caused formation of the oxynitride

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liquid followed by supersaturation and crystallization of Y-α-sialon (Liu et al., 2007).

8.2.2 Carbides Carbon forms several compounds with transition metals from groups IV– VI and with boron and silicon where the covalent bonds dominate their mechanical properties. The former (TiC, WC, MoC, ZrC), which have excellent properties such as high melting temperature, high hardness, good corrosion resistance combined with good electrical and thermal conductivity, are widely used as hard coatings and reinforcement for metal matrix composites (MMC), while the latter (SiC, B4C) with a diamond-like structure, apart from application as exceptional hard powder particles, are widely used for manufacturing bulk engineering products where their electrical/ thermal conductivity is utilized. Carbide powders are manufactured by application of solid state reactions (synthesis from elements, carbothermal reduction, self-propagating high temperature synthesis), chemical vapour deposition (CVD) or via thermolysis of organometallic polymers. All methods suffer either from the low reactivity of powders (solid state methods) or the low yield (gas phase methods). High purity amorphous nanopowders can be obtained by the organometallic route but they are limited only to silicon-based compounds as only organosilicon chemistry offers the formation of a variety of polysilanes. Thus MCP of the blended elements offers a powerful technique for the preparation of nanocrystalline carbide powders or serves as a preliminary step in the preparation of the mixtures for the subsequent SHS. Titanium carbide Titanium carbide attains a high rank among other carbides because of its exceptional hardness, high melting temperature (3100°C) and thermal stability, as it does not undergo phase transformations. On the other hand, preparation of TiC powders by carbothermal reaction or synthesis from its elements requires a high temperature over 1800°C and thus the resultant product is coarse with a broad particle size distribution which deteriorates the service properties of TiC-containing materials. Mechanochemical processing of the blended fine titanium powders and different sources of carbon (carbon black, activated carbon, fullerene, carbon nanotubes) appear to be a very attractive method for the in situ formation of TiC and were studied by several authors who show the possibility of manufacturing nanocrystalline TiC powders. Attempts to manufacture TiC powders by MCP dates back to the 1990s (Ye and Quan, 1995; Teresiak and Kubsch, 1995; Liu et al., 1994). Milling

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fine titanium powder (<74 μm) with graphite (<250 μm) in a stoichiometric composition of 50 : 50 with a high ball-to-powder (BP) ratio of 40 : 1 in an argon atmosphere led to the gradual formation of nanosized TiC (9 nm) particles after 10–27 hours of milling (Ye and Quan, 1995). Further studies showed that in situ TiC formation strongly depends on the type of carbon used for the reaction (Jia et al., 2009), the atmosphere in the reaction vial (Yuan et al., 2009), the starting composition (Lohse et al., 2005), type of the mill (Jia et al., 2009) and so on. A clue to the reaction is that it shows a highly negative free energy of formation at room temperature, as shown in Equation [8.1]: Ti + C = TiC

ΔG298 = −182.96 J mol−1

[8.1]

thus a highly exothermic reaction can be initiated by high-energy ball milling. However, a mechanically induced self-propagating reaction (MSR) would not occur unless the adiabatic reaction temperature is over 1800K (Gras et al., 2007). Since the adiabatic temperature of TiC formation from its elements is about 3290K (Grebe et al., 1992) MSR in this system is likely to occur. Thorough research, however, showed that formation of TiC under MCP could occur by two types of reaction mechanism (Lu and Li, 2005): • •

gradual reaction through elemental diffusion a self-sustaining reaction ignited after a specific period of activation.

The gradual reaction involves formation of non-stoichiometric Ti(C) as a result of carbon atoms diffusing into the Ti lattice during ball milling as evidenced by XRD (Tang et al., 2006), Raman spectroscopy (Lohse et al., 2007) and microscopic studies (El-Eskandarany, 1996). The ductile Ti particles are flattered, welded and form a lamella structure, while the fragmented graphite particles are dispersed among them (Jia et al., 2009) or both constituent powder particles form thick layers between the diffusion couples (El-Eskandarany, 1996). The Ti(C) particles are then gradually refined and take on more defects to the level of the threshold energy where transformation from Ti(C) to TiC occurs, as schematically illustrated in Fig. 8.2 (Jia et al., 2009). It has been observed that the reactant layers disappear just after formation of the new phase (El-Eskandarany, 1996). At this point in the reaction, the kinetic energy of the collisions between powder particles and the milling balls could control the transition between gradual reaction kinetics and the formation of TiC via MSR. If the milling was performed in a planetary mill at a low operating speed (Jia et al., 2009) or the low BP ratio (Tang et al., 2006) or with low density agate milling balls (Schaffer and Forrester, 1997), transformation from Ti(C) to TiC occurred gradually and a substoichiometric composition of TiC was produced, even if an elongated milling time was applied.

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Ti C

C Ti

(a)

(b)

TiC C TiC Ti (c)

(d)

8.2 Scheme showing the reaction progress of Ti and activated carbon during MCP. (a) Initial stage, no TiC; (b) second stage, the refinement of Ti and carbon powders, no TiC; (c) third stage, a small amount of TiC formed on the surface of the Ti particles; (d) final stage, the as-milled products composed completely of nanocrystalline TiC (after Jia et al., 2009, with kind permission of Elsevier).

In contrast, if the higher energy milling was applied in the form of a high BP ratio (Yuan et al., 2009), or milling was performed in a magneto ball mill (Uni-Ball-Mill 5) with the possibility of controlling the type of acting forces (Lohse et al., 2005), or other types of mill were used (a GN-2 ball mill) (Xinkun et al., 2001) or a high-energy SPEX 8000 ml mixer mill (Calos et al., 2001), the self-sustaining reaction could have been triggered. As the local temperature was raised above the ignition temperature, the reaction proceeded rapidly until the synthesis had been thoroughly completed. The development of Ti–C powder particles in the various stages of the TiC formation is presented in Fig. 8.3 (Yuan et al., 2009). The prolongation of milling after completion of the self-sustained TiC formation led to significant grain refinement down to 4–5 nm (El-Eskandarany, 2000; Jia et al., 2009). However, some obstacles to the self-sustained reaction in the Ti–C system were reported. Transformation of Ti(C) particles into TiC could be effectively slowed down if there is a carbon deficiency below the stoichiometric composition (Jia et al., 2009) or if the reaction is performed in a hydrogen atmosphere (Yuan et al., 2009). This could happen if intermediate TiH2

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100 nm (c)

100 nm

(b)

100 nm (d)

100 nm

8.3 Transmission electron microscopy (TEM) micrographs (bright field image) of Ti–C powder milled in a vacuum for (a) 20 min, (b) 40 min, (c) 1 hour and (d) 6 hours. XRD results showed full transition to TiC after 1 hour milling and TiC crystallite size refinement after 6 hours milling (after Yuan et al., 2009, with kind permission of Elsevier).

appears before TiC formation thus releasing the previously stored energy in the defective particles and alternating the reaction mechanism. Silicon carbide Silicon carbide has wide applications in the field of composite structural materials and semiconducting devices operating in a high temperature, high frequency and high power environment. Its manufacture dates back to 19th century in Acheson’s carbothermal reduction although this requires a high temperature of synthesis (2600°C) and the commercial product is coarse and contaminated with oxygen. Other manufacturing methods of ultra fine and pure SiC powders (polysilane pyrolysis) are less suitable for mass production because of the low yield and high agglomeration; thus alternative methods are of interest. Preliminary studies of the in situ formation of SiC during MCP showed it to be a time consuming procedure as the reported milling time exceeded

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300 hours (El-Eskandarany et al., 1995). In comparison to TiC formation, the reaction heat ΔH298 of silicon carbide is much lower, −73 kJ mol−1, and the Gibbs energy ΔG of the critical radius for nucleus formation, which is regarded as an energy barrier to the generation of a new phase, seems to be high (Semchenko et al., 2004). Nevertheless, after the threshold activation time, as the refinement of particles occurs, followed by thorough mixing and an increased number of chemically active defect sites, the energy barrier of the SiC nuclei formation is significantly decreased and the reaction proceeds. The further development of the milling equipment (a dual-drive planetary mill, Chaira et al., 2007) opens up new possibilities for the mechanochemical synthesis of carbides and the time for SiC formation to take place is shortened. Detailed studies of β-SiC formation from a mixture of graphite and silicon in the dual-drive planetary mill (BP = 20) showed that nearly full conversion to β-SiC was possible after 40 hours milling time (Chaira et al., 2007). XRD of the milled powder (Fig. 8.4, Chaira et al., 2007) at the selected milling times reveals the gradual formation of β-SiC via amorphization of carbon after 10 hours milling followed by the disappearance of silicon peaks and formation of a transitional bonded state between Si and C (20 hours milling), finally terminating in the incomplete ordering of the crystal lattice into defective crystallites of SiC. Calculation of the lattice strain and the crystallite size from the XRD data (Fig. 8.5, Chaira et al., 2007) clearly shows that the onset of SiC nucleation occurred after 20 hours

β SiC

Si5C3

β SiC

400

SiC

350

40 h β SiC

Intensity (a u)

300 20 h

250 200

10 h

150 Si

100

Si

C

0

0h

Si

50

Si 20

30

40

50

60

70

Si 80

2θ (degree)

8.4 XRD patterns of milled silicon and graphite powders at selected milling times (after Chaira et al., 2007, with kind permission of Elsevier).

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(b) 200

0.8

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Crystallite size (nm)

Lattice strain (%)

(a)

High-energy ball milling

0.6 0.4 6 mm ball 12 mm ball

0.2 0

0

20 40 Milling time (h)

60

6 mm ball 12 mm ball

120 80 40 0

0

20 40 Milling time (h)

60

8.5 Effect of ball size on lattice strain and the crystallite size of the milled powder calculated from X-ray diffraction patterns: (a) lattice strain; (b) crystal size (after Chaira et al., 2007, with kind permission of Elsevier).

milling and further crystallization was accompanied by a gradual decrease in the lattice strain. The final powder contained 20% of particles with a diameter below 1 μm. MCP of elemental silicon and carbon powder mixtures could be more effective if combined with self-propagating high temperature synthesis. Since the calculated adiabatic temperature of SiC formation is in the range of 1350–1450°C and because of the high partial pressure of Si(g) above 1500°C in this system, SHS requires an additional energy source and a nitrogen pressure above 3 MPa. However, strong agglomeration of the resultant particles is unavoidable because of the low melting temperature of silicon (1414°C, Pampuch et al., 1987). A milder environment for the manufacture of SiC by the SHS method could be achieved by application of polytetrafluoroethylene (PTFE). The nitrogen atmosphere could be replaced by argon and the ignition temperature would be significantly reduced because of the heat evolved in the reaction of PTFE with silicon (Nersisyan et al., 1991; Zhang et al., 2002a). However, PTFE reacts with silicon and generates poisonous fluoride according to the reaction: Si(s) + (–CF2CF2–)n(s) → SiF4(g) + C (s)

[8.2]

which should be avoided or the amount of PTFE reduced significantly. Yang et al. (2007) combined addition of PTFE with NH4Cl powders to a blended mixture of a stoichiometric amount of silicon and carbon black. Addition of 1 wt% of NH4Cl with 6 or 10 wt% of PTFE resulted in a self-sustaining reaction in argon (0.5 MPa) after ignition using a tungsten coil (current 15 A) if the mixture was previously milled in a vibration mill operating at 23.2 Hz (2MZS-3, China, BP = 10) for at least 2 hours. It is believed that

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the decomposition products of NH4Cl during MCP prevented silicon particles from oxidation and resulted in better refinement and partial amorphization of Si. Consequently, the onset temperature of SiC formation was reduced and less PTFE was necessary for SHS to start. High-energy ball milling in the ternary Ti–Si–C system involves a more complicated reaction path dependent on the starting composition, as well as the milling parameters, as formation heat in the various phases differs significantly from −72 kJ mol−1 for SiC to −579 kJ mol−1 for Ti5Si3 (Lu and Li, 2005). Ti3SiC2 powder particles with high purity and ability to sinter at low temperature are attractive for manufacturing nanostructured ceramics as potential nuclear materials for use in a lead-cooled fast reactor (Barnes et al., 2008). The slightly lower mechanical properties of Ti3SiC2 in comparison to TiC are compensated by its excellent oxidation resistance up to 1100°C if it is not contaminated by TiC precipitates. TiC is a common by-product of Ti3SiC2 irrespective of the manufacturing method (Zhang et al., 2002b) but the formation kinetics of TiC during MPC strongly depend on the applied energy as described above. Thus the precise control of the milling energy could alternate the phase formation sequence and avoid crystallization of TiC. Earlier studies of mechanochemical processing of 3Ti/Si/2C powder (10 hours) showed that Ti3SiC2 with 95.8 vol% purity was obtained after subsequent vacuum heat-treatment at 1000°C (Li et al., 2002b). Better results were obtained if more gentle milling parameters were applied (2 hours) and excess silicon was used (3Ti/1.3Si/2C) because of the unavoidable silica film on the silicon particles (Orther et al., 2002). The final heat treatment must have been conducted at a higher temperature (1350°C) but the single-phase Ti3SiC2 was reported instead (Orther et al., 2002). The crucial role of the silica film in formation of non-contaminated Ti3SiC2 was shown by an approach described by Meng et al. (2009): gentle milling conditions in the vibrating frame grinder (Pulverisette 0 Fritsch), a lower silicon excess (1.2), replacement of silicon powder by Si bars and degassing of graphite powder (1000°C, vacuum). As a result, Ti3SiC2 of good quality was produced after heat treatment (1350°C) of the as-milled powder (Fig. 8.6).

8.2.3 Silicides Silicides of molybdenum and tungsten are an important group of attractive intermetallic systems for high-temperature applications because of their excellent oxidation resistance and outstanding thermal shock resistance. Preparation of these compounds in the form of ultragrained powders is of the utmost necessity if the brittle nature of these polycrystalline products is to be overcome. The in situ synthesis from the elements does not yield a monophase powder while formation of molybdenum silicides by MCP is

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g – TiSi2 m – Ti5Si3Cx

h

Intensity (a u)

h h

(c)

h

h – Ti3 SiC2 e – TiC

h

e h m

g

h e g

(b)

g

g

(a)

30

h

35

40

45 2θ (degree)

50

g

g

55

60

8.6 X-ray diffraction patterns of milled and heat-treated powders: (a) three-element powder; (b) Ti, graphite powder and Si bar and (c) Ti powder, Si bar and degassed graphite (after Meng et al., 2009 with a kind permission of Elsevier).

sensitive to the energy of milling. Ball milling with a lower energy promotes formation of MoSi2 by diffusion in a solution of Mo-Si, but high-energy milling involves a self-sustaining reaction as it occurs in the Ti–C system (Zhang and Liu, 2001). Application of a vibratory mill (BP = 20, frequency 1000 per min) to the blended mixture of molybdenum and silicon led to formation of a solid solution of Mo(Si) and significant strain was released by crystallization of MoSi2 in an exothermic reaction followed by amorphization of the reaction product and Mo (Feng et al., 2008). In order to obtain monophase MoSi2 it is necessary to find a threshold time for milling with a given energy that is necessary to attain a sufficient and critical solid solution of Mo(Si). The subsequent annealing at higher temperature (i.e. 1000°C, 1 hour) allows crystallization of the single phase MoSi2 and grain coarsening. A similar procedure was applied to manufacture WSi2 but synthesis of Mo5Si3, Mo3Si was less sucsessful (Feng et al., 2008).

8.3

Mechanochemical processing of reactive systems

Once synthesized, non-oxide compounds such as carbides or nitrides are very resistant to any plastic deformation under mechanical stress because of their intrinsic stiffness and hardness and thus application of mechanochemical treatment to the relevant powders seems less attractive. Neverthe-

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less, there was some research on MCP in this system. First, carbide-based or nitride powders suffer from a low densification rate because of their high melting temperatures and the very low self-diffusion coefficient of the nitrogen or carbon atoms: the self diffusion coefficient of nitrogen is close to 6.3 × 10−20 cm2 s−1 at 1400°C while that of carbon is 10 times lower than for silicon in SiC (Ghoshtagore and Coble, 1966). Thus serious damage to the crystalline lattice, in addition to powder particle refinement would be beneficial for lowering the densification temperature as diffusion paths are enhanced by damage to the lattice and shortened by particle refinement, respectively. Lowering the densification temperature in any system of interest opens up new perspectives for manufacturing nanostructured ceramics, as grain growth at the final stage of sintering strongly depends on the grain boundary mobility. Moreover, any sintering additives could be perfectly distributed among the carbide/nitride-based particles under MCP resulting in more homogeneous distribution of liquid during densification. A second attractive field of MCP utilization in the non-oxide systems is activation of the blended powders before high-temperature synthesis if the endothermic reaction can be applied. These reactions involve carbothermic or alumothermic reduction/displacement as widely applied in manufacturing silicon or titanium nitride/carbide powders or several other types of composite powders in the systems of interest. It is expected that mechanochemically treated precursors will react at lower temperature and/or in the shorter reaction time. Decline of one or both above-mentioned reaction parameters is advantageous in the silicon-based compounds because of its high partial pressure at elevated temperature. The next section deals with research results in both fields of application of MCP: enhancing densification of silicon nitride/sialon powders in engineering ceramics and activation of the blended precursors in carbothermal or displacement reactions.

8.3.1 Activation of precursor in subsequent densification Silicon nitride ceramics The results of the mechanochemical treatment of silicon nitride powder have been published by Xu (Xu et al., 2005a) and showed the greater ability of this compound for densification. A mixture of ultrafine β-Si3N4 with alumina and yttria sintering additives was milled in a planetary mill (BP = 20, 475 rpm) for 4 hours and the resultant powder was further densified by SPS at 1600°C for 5 min under a pressure of 30 MPa. XRD results for the milled powders showed evidence of a highly defective β-Si3N4 crystal structure and amorphization of the sintering additives (alumina and yttria). As a result of significant damage to the powder particles the maximal shrinkage

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rate values (−d(ΔL/L0)/dt) increased as much as 20 times from 0.004 s−1 (the reference material) to 0.08 s−1 for the milled powders. Consequently a fully dense nanoceramic with a grain size around 70 nm (Xu et al., 2005a) was produced after densification at temperature 150°C lower than the reference specimen. Nowadays, however, it is widely accepted that commercially produced silicon nitride ceramics with excellent mechanical properties owe their fracture toughness to the specific microstructure with elongated β-Si3N4 grains with a high aspect ratio. This microstructure can be produced both from the commercial powders and from plasmachemical synthesis derived nanopowders if sufficient additives and temperature greater than 1700°C are applied (Herrmann et al., 2004). In contrast to the coarse microstructure of elongated β-Si3N4 grains, isometric but nanostructured silicon nitride ceramics from mechanochemically processed powders show higher hardness (18.3 GPa) and better ductility than reported elsewhere (Xu et al., 2006). Sialon ceramics Preparation of β-sialon precursor powder by mechanochemical treatment seems to be attractive for subsequent reaction sintering as more oxygen can enter the β-sialon crystal lattice and preliminary formation of a solid solution in the course of MCP is not excluded. The first report of the effect of mechanochemical treatment of various sialon precursor mixtures on the subsequent heat treatment was given by MacKenzie et al. (2003). It was shown that planetary milling (BP = 50, 400 rpm, 48 hours) enhanced ex situ the formation of various sialons (β-sialon, O-sialon, X-sialon) in comparison to the unground mixture since the nitrogen uptake took place at a temperature 130°C lower than before and crystallization of various sialons was detected by XRD after heat treatment at 1200°C regardless of the starting composition and the source of alumina (γ-Al2O3, Al(OH)3). Medium milling parameters (BP = 20, 475 rpm, 4 hours) in the planetary mill applied to a mixture of β-Si3N4 with alumina and AlN resulted in significant amorphization of the precursor component (Xu et al., 2005b). Our research showed that planetary milling of the precursor mixture (Si3N4 + AlN + Y2O3) results in a significant increase in the specific surface area Fig. 8.7) and strongly depends on the milling time in addition to the applied energy (rotational speed of the vial, ball-to-powder ratio, acceleration in the chambers, etc.). The effects of this MCP on the component particles depends on the mechanical properties of the particles involved and on the thermodynamics of the reactions concerned. Refinement of yttria and aluminium nitride crystallites is observed after a short milling time, while the α- or β-Si3N4 crystal lattice remains more resistant to the applied mechanochemical treatment (Xu et al., 2005a; Sopicka-Lizer et al., 2007a). Examination by transmission microscope © Woodhead Publishing Limited, 2010

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30 Specific surface area (m2 g–1)

750 rpm 25 500 rpm

20 15 10

250 rpm

5 0

Attritor 0

20

40

60

80

100

120

140

Milling time (min)

8.7 Specific surface area of sialon precursor powder versus milling time for a given rotational speed of the vial.

showed evidence of several types of defective crystal lattice: dislocations, damaged surface of individual particles and amorphization (Fig. 8.8). The increasing amorphization of the β- or α-Si3N4 crystal lattice was indicated by the swelling unit cell volume (Fig. 8.9). This process was accompanied by reduction of the crystallite size as measured by XRD peak broadening by Sherrer’s formula. As the mechanochemical process does not change the crystallite orientation inside the treated powder particles nor are new crystallite boundaries formed, the observed peak broadening must come from advanced lattice disordering and diminishing ordered areas as observed by transmission electron microscopy (TEM) or high resolution transmission electron microscopy (HRTEM) (Xu et al., 2005b, Sopicka-Lizer et al., 2007a). NMR studies revealed that Si–O and Al–O bonds were much more prone to deformation under MCP compared with those of Al–N or Si–N, but supplementary FT-IR results showed formation of the new Al–O bonds in a tetrahedral coordination and the creation of new Si–O–Al bonds. Thus it could be concluded that mechanical activation of the powders from the Si–Al–O–N system under a controlled atmosphere could lead to breaking and fixing the new bonds, that is solid solution formation (Sopicka-Lizer et al., 2008). The true effects of MCP in the Si–Al–O–N–(Me) system are visible during heat treatment of the mixtures as crystallization/ordering of the amorphous spots is supported by an exothermic peak on the differential thermal analysis (DTA) curve (MacKenzie et al., 2003; Xu et al., 2005b). The subsequent reactions are robustly influenced by oxygen in the system and © Woodhead Publishing Limited, 2010

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7.00 nm

8.8 TEM image of a defective β-Si3N4 particle after high-energy milling in a planetary mill. (a) TEM image, (b) HRTEM image.

292.75 Volume of the unit cell (106 pm3)

184

292.7 292.65 292.6 292.55

α-Si3N4

292.5 292.45 292.4

0

10

20

30

40

50

60

70

Milling time (min)

8.9 Volume of α-Si3N4 unit cell versus milling time in a planetary mill (750 rpm).

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the thermodynamics of the related compounds. Excess oxygen in the form of a layer of oxide on the Si3N4 particles and hydroxyl groups in the form of a hydrolysed layer on AlN particles predisposes them to form oxidebased compounds like mullite or calcian albite, if Ca-α-sialon was the intended product (MacKenzie et al., 2003). This crystallization relaxes the free energy of the system and reduces MCP effects on non-oxide crystallization or powder consolidation. The last, but not least benefit of application of MCP to the Si–Al–O–N system arises from changes in the consolidation process. Mechanochemically treated powders show good compactibility after cold isostatic pressing and green densities over 60% could be attained (Sopicka-Lizer et al., 2007b). In contrast, silicon nitride nanopowders from plasma-chemical synthesis were reported to show a lower green density of around 40% (Bulic´ et al., 2004). The very low self-diffusion coefficient of the nitrogen atoms requires application of a suitable sintering aid in order to promote liquid phase formation followed by a dissolution–transport–precipitation process resulting in the α-Si3N4 → β-sialon transformation. The latter leads to the formation of a final microstructure tailored by the careful design of the type and amount of sintering additives and thermal treatment of the manufactured objects. Precipitation of β-sialon in the mechanochemically treated precursor can be moved to a lower temperature because of precursor particles amorphization and occurs in the numerous β-Si3N4 nuclei inside defective particles. Thus a β-Si3N4 precursor is recommended for MCP instead of the standard route with α-Si3N4. Since high viscosity at lower temperature prevents substantial β-Si3N4 or β-sialon grain growth, the resultant microstructure is in the nanometer or submicrometer range regardless of the sintering technique used: SPS (Xu et al., 2005a) or pressureless (Sopicka-Lizer et al., 2009). Consolidation at higher temperature (1800°C) leads to the well-known solution–precipitation process followed by elongated grain growth. A similar microstructure is obtained from nanosized powders regardless of their origin: plasma chemical synthesis or MCP if densified at high temperature. But MCPderived specimens exhibit a finer distribution of yttria-bearing liquid phase and thus better high-temperature properties are expected for these materials (Fig. 8.10).

8.3.2 Carbothermal reduction Manufacturing of non-oxide powders by a carbothermal reaction (CR) is one of the oldest and most attractive commercial methods. However, high temperature and the low reaction rate in the solid state prompt a search for an improvement. Manufacturing SiC or TiC via carbothermal reduction could be significantly enhanced if the precursor mixture is

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1 μm (b)

1 μm

8.10 Microstructure of gas-pressure sintered (1800°C, 5 bar) β-sialon from an unground mixture (a) and from a mechanochemically processed precursor milled for 1 hour (500 rpm) in a planetary mill (b) (SEM).

preliminary processed by mechanochemical milling. The former was successfully performed from a preliminary milled (2 hours) mixture of graphite and SiO2, resulting in a decrease in the reduction temperature of 500°C and a shortening of the reaction time from 30 hours to 1 hour (Shaw, 1998). Competitive methods for SiC nanopowders production via polysilane routes further limit research into mechanochemical processing in this system. In contrast to the availability of several methods of SiC nanopowder production, manufacturing titanium carbide nanopowder remains a point of interest. Application of attrition milling (BP = 60, 600 rpm, 2 hours) to a mixture of anatase–TiO2 and graphite resulted in a decrease in the reaction temperature of 500°C and a shortening of the reaction time to 1–2 hours (Ren et al., 1998). Similar results were obtained if a mixture of rutile– graphite was milled in a planetary mill (BP = 10, 50 hours) and subsequently reacted in argon. The resultant powder consisted of non-stoichiometric TiC powder with 80 nm crystallites if heat treatment was conducted at 1250– 1500°C for 1 hour (Razavi et al., 2008).

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Mechanochemical treatment of the starting precursor for subsequent carbothermal reduction/nitridation was also applied to β-sialon manufacture. A reaction mixture of halloysite clay and fine carbon was milled in a planetary mill (BP = 50, 400 rpm, 4–48 hours) under flowing nitrogen and afterwards reacted for 4 hours in the temperature range 1200–1400°C (MacKenzie and Barneveld, 2006). The results of XRD complemented by NMR studies revealed an interesting feature of β-sialon formation and the accompanying phases: mullite, corundum and 21R AlN polytypoid. Preliminary milling reduces the formation temperature of the β-sialon phase during CR by 100°C but the reaction sequence and intermediate phase formation depend on the milling time. The longer the milling time, the higher the ratio of corundum to mullite formed, thus transformation from mullite to β-sialon was hindered. As standard carbothermal reduction of kaolinite in nitrogen engages crystallization of transient mullite followed by transformation to β-sialon (Lee and Cutler, 1979), MCP deformation of the starting clay leads to amorphous silica and alumina. Consequently, mullite formation is prevented and the amount of β-sialon is reduced at the end of carbothermal reduction. This shows that the imperfect structure of mullite after clay decomposition is crucial in the successive formation of β-sialon. A carbothermic or alumothermic reaction can also be applied to the synthesis of intermetallic compounds by preliminary mechanochemical treatment of the relevant oxide and the reducing agent (carbon, aluminium). There are reports on manufacturing carbides, silicides and borides or nanocomposite particles by MCP of the blended components followed by appropriate high temperature annealing (Li et al., 2001; 2002a). Moreover, powders like mechanochemically processed MoSi2 or those where Si3N4 has been added could be consolidated to fully dense monolithic ceramics by vacuum hot-pressing at 1400°C and their microstructure was retained in the submicrometer range (Suryanarayana, 2008). Application of carbon or aluminium to reduction of the mixture of MoO3 and SiO2 led to formation of molybdenum disilicide after ball milling (700 rpm, BP = 10) for 3–100 hours without a following heat treatment (Zakeri et al., 2007). Reduction by aluminium was more effective, proceeded by a self-sustaining reaction and resulted in development of nearly spherical composite MoSi2/Al2O3 particles with a diameter of around 500 nm and an alumina crystallite size below 15 nm. Prolongation of the milling time from 40 to 70 hours led to a phase transformation of α-MoSi2 to hightemperature β-MoSi2 accompanied by crystallite coarsening. Reduction of MoO3 by carbon in the presence of quartz was incomplete even after 100 hours milling because of the low reactivity of quartz. These examples show the versatility of the tested non-oxide systems and prospects for further research.

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8.4

Conclusion

Application of mechanochemical treatment to manufacturing or processing of nitrides/carbides fine powders is shown to be an attractive and effective method in various systems. The main advantages lie in lowering the temperature of the synthesis reaction itself or the subsequent powder reaction/ consolidation. This feature is enormously attractive as carbides of interest melt at high temperatures, usually over 2000°C. On the other hand, siliconbearing compounds suffer from high volatility of silicon at elevated temperatures, so that reduction of the consolidation temperature eliminates high-pressure densification. Submicrometre particles of mechanochemically treated powders attain an ability for compaction and simultaneously offer the possibility of preserving the nanostructure in the resultant ceramic product. The commercial significance of the described methods is seen among the compounds, which cannot be produced by alternatively methods (TiC, Ti3SiC2, composite particles, etc.). Further research and development is essential for a better understanding of the effects of MCP in the treated powder particles. The application of mechanochemical treatment to the above systems has some limitations. The results described showed that reactions induced by mechanochemical treatment could follow a different path from the heatactivated ones. This is not necessarily advantageous in terms of the expected product, as shown in the formation of β-sialon under carbothermal reduction–nitridation preceded by mechanochemical processing of the clay– carbon mixture. The formation of transient phases during the heat treatment of mechanochemically activated powders is always inconvenient, as crystallization consumes the excess energy located in the activated powders. However, the main drawback of the method is contamination from the milling media and the vial walls. These aspects are not made clear in the published literature, although they are mostly unavoidable and could have a detrimental effect on the properties of the final product, that is contamination of silicon nitride or carbide by iron from stainless steel milling balls could lead to the formation of low melting iron silicides during further heat treatment. The choice of the proper lining of the vial and milling media could be crucial for the properties of the final product although replacement of stainless balls by ceramic balls decreases the input of energy needed during milling. Nevertheless, mechanochemical processing of non-oxide covalent compounds has bright prospects for future development.

8.5

References

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bulic´ f.i., zalite i. and zhilinska n. (2004), ‘Comparison of plasma-chemical synthesised SiAlON nano-powder and conventional prepared SiAlON powder’, J. Eur. Ceram. Soc., 24, 3303–6. calka a. and radlinski a.p. (1991), ‘Formation of amorphous Fe–B alloys by mechanical alloying’, J. Appl. Phys. Lett., 58, 119–21. calos n.j., forrester j.s. and schaffer g.b. (2001), ‘The mechanisms of combustion and continous reactions during mechanical alloying’, J. Solid State Chem., 158, 268. cano i.g., perez baelo s., rodriguez m.a. and aza s. (2001), ‘Self-propagating high temperature synthesis of Si3N4: role of ammonium salt addition’, J. Eur. Ceram. Soc., 21, 291–5. chaira d., mishra b.k. and sangal s. (2007), ‘Synthesis and characterization of silicon carbide by reaction milling in a dual-drive planetary mill’, Mater. Sci. Eng. A, 460–1, 111–20. ekström t. (1989), ‘Effect of composition, phase content and microstructure on the performance of Yttrium–Si–Al–O–N ceramics’, Mater. Sci. Eng., A109, 341–9. el-eskandarany m.s. (1996), ‘Synthesis of full density nanocrystalline tungsten carbide by reduction of tungstic oxide at room temperature’, Metall. Trans. A, 27, 2374–82. el-eskandarany m.s. (2000), ‘Structure and properties of nanocrystalline TiC fulldensity bulk alloy consolidated from mechanically reacted powders’, J. Alloys Compd., 305, 225–38. el-eskandarany m.s., sumyiama k. and suzuki k. (1995), ‘Mechanical solid state reaction for synthesis of β-SiC powders’, J. Mater. Res., 10(3), 659–67. feng p., farid a., wang x., humail i.s. and xuanhui q. (2008), ‘Mechanically activated reactive synthesis of refractory molybdenum and tungsten silicides’, Int. J. Refract. Met. Hard Mater., 26, 173–8. ghoshtagore r.n. and coble r.l. (1966), ‘Self-diffusion in silicon carbide’, Phys. Rev., 143, 623–6. gras ch., zink n., bernard f. and gaffet e. (2007), ‘Assisted self-sustaining combustion reaction in the Fe-Si system: mechanical and chemical activation’, Mater. Sci. Eng. A, 456, 270–7. grebe h.a., advani a., thadhani n.n. and kottke t. (1992), ‘Combustion synthesis and subsequent explosive densification of titanium carbide ceramics’, Metall Trans., 23(9), 2365–72. grun r. (1979), ‘The crystal structure of β-Si3N4’, Acta Crystallogr., B35, 800–4. herrmann m., schultz i. and zalite i. (2004), ‘Materials based on nanosized β-Si3N4 composite powders’, J. Eur. Ceram. Soc., 24, 3327–35. jack k.h. and wilson w.i. (1972), ‘Ceramic based on the Si–Al–O–N and related system’, Nature (London), Phys. Sci., 238, 28–9. jia h., zhang z., qi z., liu g. and bian x. (2009), ‘Formation of nanocrystalline TiC from titanium and different carbon sources by mechanical alloying’, J. Alloys Compd., 472, 97–103. jin h.b., yang y., chen y.x., lin z.m. and li j.t. (2006), ‘Mechanochemical activationassisted combustion synthesis of alpha-Si3N4’, J. Am. Ceram. Soc., 89, 1099–102. jin h.b., cao m.s., chen y.x., li j.t. and agathopoulos s. (2008), ‘The influence of mechanochemical activation on combustion synthesis of Si3N4’, Ceram. Int., 34, 1267–71. kata d., lis j., pampuch r. and sobierski l. (1998), ‘Preparation of fine powders in the Si–C–N system using SHS method’, Intern. J. SHS, 7(4), 475–85.

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kohatsu i. and mccauley j.w. (1974), ‘Re-examination of the crystal structure of α-Si3N4’, Mater. Res. Bull., 9, 917–20. lee g.l. and cutler b. (1979), ‘Sinterable sialon powder by reaction of clay with carbon and nitrogen’, Am. Ceram. Soc. Bull., 58(9), 859–71. li j., li f., hu k. and zhou y. (2001), ‘TiB2/TiC nanocomposite powder fabricated via high energy ball milling’, J. Eur. Ceram. Soc., 21, 2829–33. li j., li f., hu k. and zhou y. (2002a), ‘Formation of TiB2/TiN/Ti (CxN1-x) nanocomposite powder via high-energy ball milling and subsequent heat treatment’, J. Alloys Compd., 334, 253–60. li j.f., matsuki t. and watanabe r. (2002b), ‘Mechanical-alloying-assisted synthesis of Ti3SiC2 powder’, J. Am. Ceram. Soc., 85, 1004. liu g., chen k., zhou h., li j., pereira c. and ferreira j.m.f. (2007), ‘Mechanicalactivation-assisted combustion synthesis of α-SiAlON in air’, Mater. Res. Bull, 42, 989–95. liu z.g., guo j.t., ye l.l., li g.s. and hu z.q. (1994), ‘Formation mechanism of TiC by mechanical alloying’, Apply. Phys. Lett., 65, 2666. lu c.j. and li z.q. (2005), ‘Structural evolution of the Ti–Si–C system during mechanical alloying’, J. Alloys Compd., 395, 88–92. lohse b.h., calka a. and wexler d. (2005), ‘Effect of starting composition on the synthesis of nanocrystalline TiC during milling of titanium and carbon’, J. Alloys Compd., 394, 148–51. lohse b.h., calka a. and wexler d. (2007), ‘Raman spectroscopy sheds new light on TiC formation during the controlled milling of titanium and carbon’, J. Alloys Compd., 434–5, 405–9. mackenzie k.j.d., temuujin j., smith m.e., okada k. and kameshima y. (2003), ‘Mechanochemical processing of sialon compositions’, J. Eur. Ceram. Soc., 23, 1069–82. mackenzie k.j.d. and barneveld d. (2006), ‘Carbothermal synthesis of β-sialon from mechanochemically activated precursors’, J. Eur. Ceram. Soc., 26, 209–15. meng f., chaffron l. and zhou y. (2009), ‘Synthesis of Ti3SiC2 by high energy ball milling and reactive sintering from Ti, Si, and C elements’, J. Nucl. Mater., 386–8, 647–9. munir z.a. and tamburini a. (1989), ‘Self-propagation exothermic reactions: the synthesis of high-temperature materials by combustion’, Mater. Sci. Rep., 3, 277–365. nersisyan g.a., nikogosov v.n., kharatyan s.l. and merzhanov a.g. (1991), ‘Chemical transformation mechanism and combustion regimes in the system siliconcarbon-fluoroplastic’, Combust. Expl. Shock Waves, 27, 720. nikolov j.i., calka a. and williams j.s. (1995), ‘The effect of milling condition on micro/macrostructural transformations during milling of B and Si in nitriding environment: formation of BN and Si3N4’, NanoStructured Mater., 6, 401–4. orther h.r., tomasi r. and botta w.j. (2002), ‘Reaction sintering of titanium carbide and titanium silicide prepared by high energy milling’, Mater. Sci. Eng. A, 336, 202–8. oyama y. (1971), ‘Solid solution in the ternary system Si3N4-AlN-Al2O3’, Jpn. J. Appl. Phys., 10, 1687. pampuch r., stobierski l., lis j. and ra˛czka m. (1987), ‘Solid combustion synthesis of SiC’, Mater. Res. Bull., 22, 1225.

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razavi m., rahimipour m.r. and kaboli r. (2008), ‘Synthesis of TiC nanocomposite powder from impure TiO2 and carbon black by mechanically activated sintering’, J. Alloys Compd., 460, 694–8. ren r.-m., yang z.-g. and shaw l.l. (1998), ‘Synthesis of nanostructured TiC via carbothermic reduction enhanced by mechanical milling’, Scripta Mater., 38(5), 735–41. schaffer g.b. and forrester j.s. (1997), ‘The influence of collision energy and strain accumulation on the kinetics of mechanical alloying’, J. Mater. Sci., 32, 3157–62. semchenko g.d., opryshko i.n., angolenko l.a., vernigora k.p. and kalin v.v. (2004), ‘Mechanochemical nucleation and stability of newly-formed phases’, Refractories Industrial Ceram., 45(3), 181–4. shaw l.l. (1998), ‘Processing of nanostructured carbides, nitrides and their composites’, Adv. Eng. Mater., 21(11), 721–3. shen t.d., koch c.c., mccornick t.l., nemanich r.j., huang j.y. and huang j.g. (1995), ‘The structure and property characteristics of amorphous/nanocrystalline silicon produced by ball milling’, J. Mater. Res., 10, 139–48. sivakumar r., aoyagi k. and akiyama t. (2007), ‘Effect of mechanically activated raw materials on β-sialon formation by combustion synthesis’, J. Mater. Res., 22, 2863–7. sivakumar r., aoyagi k., watanabe t. and akiyama t. (2009), ‘Thermal conductivities of β-sialons by mechanically activated combustions synthesis’, Key Eng. Mater., 403, 139–40. sopicka-lizer m., tan´ cula m., pawlik t., kochnev v. and fokina e. (2007a), ‘The new top-to-bottom method of sialon precursor preparation by activation in a planetary mill with a high acceleration’, Mater. Sci. Forum, 554, 59–64. sopicka-lizer m., pawlik t., włodek t., tan´cula m. and chernik g. (2007b), ‘The effect of sialon precursor nanostructurization in a planetary mill on the properties of sintered ceramics’, Key Eng. Mater., 352, 179–84. sopicka-lizer m., tan´cula m., włodek t., rodak k., hüller m., kochnev v. and fokina e. (2008), ‘The effect of mechanical activation on the properties of β-sialon precursor’, J. Eur. Ceram. Soc., 28, 279–88. sopicka-lizer m., pawlik t., włodek t. and tan´ cula m. (2009), ‘The phase evolution in the Si3N4-AlN system after high-energy mechanical treatment of the precursor powder in SiAlONs and non-oxides’, Key Eng. Mater., 403, 7–10. suryanarayana. (2008), ‘Structure and properties of ultrafine-grained MoSi2 + Si3N4 composites synthesized by mechanical alloying’, Mater. Sci. Eng. A, 479, 23–30. tang w.m., zheng z.x., wu w.c., lü j., liu j.w. and wang j.m. (2006), ‘Structural evolutions of mechnically alloyed and heat treated Ti50C50 and Ti33B67 powders’, Mater. Chem. Phys., 99, 144–9. teresiak a. and kubsch h. (1995), ‘X-ray investigations of high energy ball milled transition metal carbides’, NanoStructured Mater., 6, 671–4. xinkun z., kunyu z., baochang ch., qiushi l., xiuquin z., tieli ch. and yunscheng s. (2001), ‘Synthesis of nanocrystalline TiC powder by mechanical alloying’, Mater. Sci. Eng. C, 16, 103–5. xu x., nishimura t., hirosaki n., xie r-j., zhu y., yamamoto y. and tanaka h. (2005a), ‘New strategies for preparing nanosized silicon nitride ceramics’, J. Am. Ceram. Soc., 88(4), 934–7.

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xu x., nishimura t., hirosaki n., xie r-j., yamamoto y. and tanaka h. (2005b), ‘Fabrication of β-sialon ceramics by high-energy mechanical milling and spark plasma sintering’, Nanotechnology, 16, 1569–73. xu x., nishimura t., hirosaki n., xie r-j., yamamoto y. and tanaka h. (2006), ‘Superplastic deformation of nano-sized silicon nitride ceramics’, Acta Mater., 54(1), 255–62. yang k., yang y., lin z-m., li j-t., du j-s. (2007), ‘Mechanical-activation-assisted combustion synthesis of SiC powders with polytetrafluoroethylene as promoter’, Mater. Res. Bull., 42, 1625–32. ye l.l. and quan m.x. (1995), ‘Synthesis of nanocrystalline TiC powders by mechanical alloying’, NanoStructured Mater., 5(1), 25–31. yuan q., zheng y. and yu h. (2009), ‘Mechanism of synthesizing nanocrystalline TiC in different milling atmospheres’, Int. J. Refract. Metal. Hard Mater., 27, 696–700. zakeri m., yazdani-rad r., enayati m.h., rahimipour m.r. and mobasherpour i. (2007), ‘Mechanochemical reduction of MoO3/SiO2 powder mixtures by Al and carbon for the synthesis of nanocrystalline MoSi2’, J. Alloys Compd., 430, 170–4. zakorzhevskii v.v. and borovinskaya i.p. (2000), ‘Some regularities of α-Si3N4 synthesis in a commercial SHS reactor’, Int. J. SHS, 9, 171–91. zhang h. and liu x. (2001), ‘Analysis of milling energy in synthesis and formation mechanisms of molybdenum disilicide by mechanical alloying’, Int. J. Refract. Metal. Hard. Mater., 19, 203–8. zhang j., jeong j.c., lee j.h., won c.w., kim d.j. and kim c.o. (2002a), ‘The effect of carbon sources and activative additive on the formation of SiC powder in combustion reaction’, Mater. Res. Bull., 37, 319. zhang z.f., sun z.m. and hashimoto h. (2002b), ‘Rapid synthesis of ternary carbide Ti3SiC2 through pulse-discharge sintering technique from Ti/Si/C powders’, Metall. Mater. Trans. A, 33(11), 3321.

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9 Mechanochemical synthesis of metallic– ceramic composite powders K. W I E C Z O R E K - C I U R O WA, Cracow University of Technology, Poland

Abstract: The chapter discusses several variants of mechanosyntheses for composite powder formation of metal alloy matrices with ceramic particles. The necessity of following progress in mechanochemical processes using different analytical methods is shown. Based on the results of experimental studies on Cu-Al/Al2O3 and Ni-Al/Al2O3 nanocomposites powder formation through mechanochemical synthesis, the method of transforming combustive processes to progressive ones by applying hydroxosalts instead of metal oxides as precursors of composites is described. Key words: metallic–ceramic composites, high-energy ball milling, metallothermic reactions, combustive self-propagating reactions, progressive reactions, soft mechanochemistry, matrices of Cu–Al and Ni–Al alloys.

9.1

Introduction

Technical progress is dependent on the development of new materials. Recently a lot of attention has been given to technologies which create and test functional materials, both constructive and tools. Materials properties, especially on the nanometric scale, are specific, comparing well with poly- or monocrystalline ones for materials with the same chemical composition (e.g. Fernández-Bertran, 1999; Ivanov and Suryanarayana, 2000; Suryanarayana, 2001a; Sˇepelak, 2002; Kelsall et al., 2005; Balázˇ, 2008). From a practical point of view, the desired features of nanocrystalline powders can be intensified by creating multiphase composites from them. Therefore, bulk composites are engineering materials made from two or more components with significantly different physical and/or chemical properties that remain separate and distinct on the macroscopic level within the finished structures. One of the components is a matrix and the other provides structural support. Three types of composite categories can be defined based on the characteristics of the matrix: • • •

metallic matrix composites (MMCs) ceramic matrix composites (CMCs) polymer matrix composites (PMCs). 193 © Woodhead Publishing Limited, 2010

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Application of the relevant composite powders to produce the final bulk product or any other composite product is one of the most attractive features in their fabrication. These composite powders need to be consolidated at elevated temperatures by thermomechanical processing and then their microstructure stability and microhardness should be determined in order to estimate the optimal parameters for the consolidation process. Actual metal–ceramic composite powders usually incorporate light alloys such as Al, Mg and Ti. Other metals commonly used are Cu, Ni, Fe, Co, Zn and Mo. Metals, ceramics or polymers can be used to support MMC powders. Metal matrix reinforced phases can exist as dispersed particles of (Kaczmar, 2000; Koch, 2001; Agrawal and Sun, 2004): • metal oxides (ZrO2, Al2O3, ThO2,Cr2O3) • carbides (SiC, TaC, WC, B4C) • nitrides (Si3N4, TaN, ZrN, TiN), • borides (TaB2, ZrB2, TiB2, WB). Cermetals are comprised of metallic and ceramic components, which combine the physical properties of metals and ceramics. The properties of metals include: • • •

ductility tensile strength thermal and electrical conductivity.

Ceramics possess physical properties such as: • • •

high melting point chemical stability resistance to oxidation.

Milling of soft metals/alloys and hard ceramic materials leads to the formation of valuable composites powders with important new features. Changing the ratio of the cermetal components produces final composite products with widely different properties. These materials have much higher mechanical, temperature and aggressive media resistance properties compared to ceramics properties. The hardness of cermetals increases with decreasing particle size (Chang et al., 1999; Zhang et al., 2000; Portnoy et al., 2002).

9.2

Composite powder formation: bottom-up and top-down techniques

Composite powders can be produced by two techniques: bottom-up and top-down. The first method is based on building nanostructures atom-by-atom, layer-by-layer. These chemical processes could take place in liquid, solid as well as in gas phases (Tjong and Chen, 2004). Examples of these are the © Woodhead Publishing Limited, 2010

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sol–gel method, melt spinning–melt quenching (MQ), chemical vapour deposition (CVD), physical vapour deposition (PVD), plasma ablation, laser pyrolysis, and so on. Unfortunately, there are often some disadvantages, for example multi-stage processes, the high cost of the alkoxide precursors, solvent evaporation and necessity for thermal treatment at high temperatures to provide coarse-grained products. The top-down technique begins with macrostructured materials and uses mechanical, chemical or other forms of energy to ‘break’ them into smaller pieces. However, it is very important to develop methods which minimize damage to the environment. One promising candidate is mechanochemistry, often referred to as ‘green’ processes. Mechanical energy can be easily explored for chemical syntheses of new functional materials like composite powders. No doubt, this method is fast, economical and gives high purity products, which can often take nanostructure forms (Heegn et al., 2003). In mechanochemical syntheses, the course of solid state reactions can be effectively controlled and regulated by choosing the precursors of syntheses and/ or milling variables.

9.2.1 Mechanical treatment Mechanical treatment of solids uses mechanical energy through high-energy ball milling, abrasion, fracture or welding to produce new materials or generate products with desired features, for example powder mixtures consisting of metal and ceramics (Courtney, 1994; Boldyrev et al., 1996). Thus, mechanical treatment of reactants can form metallic–ceramic composite powder particles. The structure quality, that is whether such composite particles are micro or nanosized depends on the type of reagents and milling conditions.

9.2.2 Mechanochemical synthesis: mechanical activation, mechanical alloying, reactive milling Mechanochemical synthesis includes mechanical alloying together with reactive milling which follows from mechanical activation of reagents. This is also done by high-energy ball milling (Murty and Ranganathan, 1998; Suryanarayana, 2001; Suryanarayana et al., 2001; Avvakumov et al., 2001; Boldyrev, 2006; Smolyakov et al., 2007; Balázˇ, 2008) which is a more complex but very effective process of metallic–ceramic composites formation. There are several variants of mechanochemical syntheses, which depend on the nature of the initial reagents nature (e.g. two different metals with ceramics, metal oxide with active metal or metal salt with active metal, and others). Mechanical activation is responsible for enhancing the reactivity of solids by enlarging the surface area or accelerating the reaction by correct mixing © Woodhead Publishing Limited, 2010

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of the reagents. In the processes of plastic deformation and fracture and friction during ball collisions, the impact energy is converted into other forms of energy. These induce structural defects, broken bonds and other forms of excess energy. These accumulate and a new, active state of the substances is produced. Consequently, the chemical reactivity of the solids increases significantly (e.g. Chen et al., 1997). Mechanical alloying is the process in which mixtures of powders are milled to achieve alloying at the atomic level. During this process, when high-energy impulses are used, metallic powders can form alloys like different solid solutions, intermetallic phases, mixtures of metals or amorphic materials with properties that are very often different from those found when using traditional methods (McCormick et al., 1989; Maurice and Courtney, 1990; Schwarz, 1996; Urakaev and Boldyrev, 2000). It is important to note that any form of metal alloy which comprises composite components, has better physical and mechanical properties than pure metals. For example, consolidated materials based on intermetallics of nickel, iron, titanium with aluminium are treated as functional materials that have specific physical properties and as constructive materials that have unique mechanical features such as structural stability at high temperature. Especially important are their high melting temperature, high mechanical resistance and low density. These properties allow them to be used in the automobile industry (turbo compressor rotors, valves, combustion chamber and exhauster details), in the aircraft industry, in nuclear power and in metallurgy. However, the low yield point at room temperature (what causes difficulties in mechanical treatment) is one of the main hindrances to using them more widely. Modification of features of the intermetallic phase can be made, for example by adding ceramic particles into their structure. Inserting ceramic particles (type: Al2O3, ZrO2 or TiC) into the intermetallic matrix (Ni, Al) leads to improvements in some mechanical properties like hardness and abrasion resistance at high temperatures and a significant rise of utility features, for example an increase in the corrosion temperature and in oxidation and erosion resistance. The mechanisms of metal alloy formation by mechanical treatment of metal powders are categorized into three different systems depending on the mechanical features: ductile–ductile, ductile–brittle and brittle–brittle. In the first case in the early stages of milling, the components become flatted to platelets by cold welding. In the next stage, they are cold-welded together and form a composite lamellar geometry of the constituent metals. Further true alloying occurs at the atomic level until the homogeneous structure of the powders is attained. The mechanisms of metallic–ceramics composites formation fall into the ductile–brittle components category because the brittle oxide particles are dispersed in a ductile metallic matrix. The ceramics reagent is closely spaced

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along the interlamellar spacing. With further milling, the ductile powder of alloys particles becomes work hardened and the lamellae become convoluted and refined. With continued milling the lamellae are further refined and the brittle particles are uniformly dispersed in the ductile metallic matrix. The ductile components are flattened by a micro-forging process while the brittle ones are fragmented. One can assume that the brittle particles of the materials are dispersed in the ductile matrix. However, if both milled materials are brittle, this phenomenon was not observed. A typical example of ductile–ductile microstructures is Cu/Al–Al2O3 composite powder formed during milling of the Cu-hydroxocarbonate and Al mixture, described in Section 9.5 (Wieczorek-Ciurowa et al., 2000, 2003a, 2003b). The microstructure of the products after the first stage of the mechanochemical synthesis is shown in Fig. 9.1 which presents the SEM microphotograph. The material reveals a lamellar microstructure. Based on quantitative energy dispersive X-ray elemental analysis (EDX) (Table 9.1) it is possible to estimate that mechanically alloyed composite particles consist of Cu–Al intermetallic phases as a matrix, and aluminium oxide.

+2

+5

+4

+1

+3

20.0 µm

9.1 SEM microphotograph (BSE image) illustrating a lamellar structure of Cu–Al/Al2O3 composite powder following aluminothermic reaction with Cu2(OH)2CO3 at the first stage of reactive milling. The darker network corresponds to a higher amount of Al2O3, the brighter one to a higher content of Cu–Al intermetallics (SEM Hitachi S-4700 with EDX Noran Vantage) (Wieczorek-Ciurowa et al., 2005).

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Table 9.1 Results of the EDX elemental microanalysis of the composite powder synthesized by milling a Cu2(OH)2CO3–Al2O3 mixture, from the corresponding SEM microphotograph (BSE image) in Fig. 9.1 (Wieczorek-Ciurowa et al., 2005) Concentration/mass (%) Point

Cu

O

Al

BSE image

1 2 3 4 5

32.7 39.0 47.2 38.4 23.3

30.1 25.5 22.2 25.3 32.6

37.2 35.5 30.6 35.5 44.1

Dark grey Grey Light grey Grey Dark grey

The darker network corresponds to the higher amount of Al2O3 and the brighter phase to the higher amount of metallic phases. The final form of Cu-Al/Al2O3 composite particles shown in Fig. 9.2 reveals its homogeneous microstructure. Reactive milling of solids relates to the process in which chemical reactions occur. This is a stage of mechanochemical synthesis that can occur in two different kinetic ways: • •

a reaction which develops slowly with each collision, which results in a gradual transformation of the substrates; a self-propagating reaction initiated when the reaction enthalpy is sufficiently high.

These kinetics models depend on the nature, type of reactants and their properties (e.g. hardness), thermodynamics of the reactions and technical parameters of milling (Butyagin, 2000, 2003; Takacs et al., 2001; Takacs and Mandal, 2001; Takacs, 2002; Lyakhov et al., 2008). The first type of mechanical syntheses concerns reactions, which proceed more slowly up to the point at which processes become a function of the milling time. If ignition does not occur, the collisions between milled material and the grinding medium contribute to comminution, mixing and defect formation. The formation of the final product occurs gradually. Figure 9.3 shows the first and final step of mechanochemical synthesis of Cu–Al/Al2O3 composite powder as a function of temperature and pressure inside the mill (Wieczorek-Ciurowa et al., 2007b). The second type requires a critical time for ignition of the reaction. When the temperature of the vial was recorded during milling, it was observed that initially the temperature increased slowly with time. After a certain period of milling the temperature increased abruptly, which confirmed the fact that ignition has occurred. The time at which a sudden increase in temperature occurs is referred to as the ignition time. Beyond this time, the

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AI2O3

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5.00 µm

9.2 SEM microphotograph (BSE image) illustrating a high level of structure homogeneity in a Cu–Al/Al2O3 composite powder following aluminothermic reaction with Cu2(OH)2CO3 at the final stage of reactive milling. The darker spots are finely dispersed Al2O3 particles whose average size is no larger then 1 μm. The brighter spots correspond to the higher content of Cu–Al intermetallics (see Fig. 9.1) (SEM Hitachi S-4700 with EDX Noran Vantage) (Wieczorek-Ciurowa and Gamrat, 2007b).

reaction takes place within seconds (see Fig. 9.4). The example in the figure is an illustration of Cu–Al/Al2O3 composite powder formation in the CuO– Al system (Wieczorek-Ciurowa and Gamrat, 2007a; Wieczorek-Ciurowa et al., 2007a). The ignition temperature is a function of the enthalpy change and the microstructure parameters of the particles, for example the interfacial area between the reactants. Continuing milling is necessary to form the correct desired structure of synthesis products. Generally, a self-propagating high-temperature synthesis (SHS) reaction consists of three stages: component activation, chemical reaction initiation and a synthesis process in bulk with product formation. Schaffer and McCormick (1992) stated that SHS reactions occur when the temperature created by the colliding balls in a ball mill is higher than initiation temperature, which depends on enthalpy changes in the process. The decreasing of this temperature can be the effect of mechanical activation by diminution of particle size and increasing the surface between interfacial contacts. The explanation of this phenomenon is as follows. It was assumed that after a period of comminution, mixing and activation, agglomerates begin

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20

Milling time (h)

9.3 Variations in temperature and pressure (GTM results) during mechanosynthesis of Cu–Al/Al2O3 composite powder from Cu2(OH)2CO3–Al mixture: (a) beginning of mechanical treatment and (b) final steps of mechanosynthesis (planetary ball mill Fritsch GmbH Pulverisette 6) (Wieczorek-Ciurowa et al., 2007b).

to form and increase in size. The reaction starts in a single agglomerate or in the powder layer coating a milling ball or the wall of the vial. One reaction front propagates into other parts of the powder layer. The powder layer can be attached to the surface of a milling ball or the inner wall of the container. When a ball hits this layer, part of the kinetic energy is

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140 Temperature

28

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Pressure (kPa)

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120 26

Pressure

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24 10

20

30

40

50

Milling time (min)

9.4 Variations in temperature and pressure (GTM results) during selfpropagating high-temperature synthesis of Cu–Al/Al2O3 composite powder from a CuO–Al mixture. The milling conditions are the same as used for a Cu2(OH)2CO3–Al mixture treatment (see Fig. 9.3) (planetary ball mill Fritsch GmbH Pulverisette 6) (Wieczorek-Ciurowa et al., 2007a).

transferred to the powder as heat and a local increase of temperature can occur. The stresses inside the powder are not uniform but are concentrated at a few points. This results in the formation of ‘hot spots’, where the reaction can start even if the average temperature of the powder is not sufficient to initiate a reaction front. Intimate contact between the reactant phases is an essential requirement for self-propagating synthesis. This condition is easy to achieve when mechanical activation is conducted in a system of ductile-brittle substances. The temperature of initiation of a metallothermic reaction can be recognized by differential thermal analysis (DTA) (strong exoeffects). This is illustrated by the DTA curve (Fig. 9.5) for the product of the CuO–Al powder after the first step of mechanochemical synthesis. The CuO formed by the mechanical decomposition of Cu2(OH)2CO3 reacts with Al forming copper in a combustible manner at about 600°C. This procedure can be considered a simulation of a mechanochemical reaction (WieczorekCiurowa and Gamrat, 2007a). An example of self-propagation reactions are processes occurring during milling aluminium and magnetite powders in argon, given by Botta et al. (2000) as follows:

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Exo

Endo

Aluminothermic reduction of CuO

Eutectic reaction CuAI2 + AI→L

Decomposition of Cu2(OH)2CO3 100

300

500

700

Temperature (°C)

9.5 DTA curve (non-oxidizing atmosphere) of a Cu2(OH)2CO3–Al mixture after the first step in mechanical treatment. The product contains an undecomposed part of Cu2(OH)2CO3 and unreacted Al. This also represents a simulation of the reactions in the system that may occur under mechanical action (SDT 2960 TA Instrument) (Wieczorek-Ciurowa et al., 2005).

3Fe3O4 + 8Al → 9Fe + 4Al2O3

[9.1]

There are two different effects which are due to the time period of mechanical activation. After 30 min milling, a decrease in the starting reactant crystallinity could be observed as well as a small amount of product, whereas after 37 min milling time the aluminothermic reaction was completed. Traditionally, this reaction would be thermally initiated during high-temperature treatment. However, under mechanical activation conditions a ‘flash point’ is the direct consequence of previous energy accumulation in the crystalline structure of powder particles. In the case of milling time less than 30 min a certain amount of active product precursor forms from the energy accumulated in the solid phase. The precursor releases this energy by reacting at lower temperatures during thermal treatment. This effect is close to the one observed in the TiO2–Al system, but in this case, the aluminothermic reaction does not occur during milling. This phenomenon of solid activation, because of milling, can be seen in syntheses other than in metallic–ceramic composites ones. The kinetics of mechanochemical reactions depends not only on the reactant properties but also on the synthesis conditions. The synthesis can be explosive or occur in a controlled steady way. Changing the conditions of the activation process can prevent the reaction taking place at all. As well as the ball-to-powder ratio (BPR), the type of ball material is also

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Table 9.2 Examples of induction time intervals for selected SHS mechanochemical reactions occurring under different milling conditions (literature data)

Reaction

Type of mill

ZnO + Mg → Zn + MgO

Vibratory

7:1

45

2FeO + Ti → 2Fe + TiO2 8CuO + 3Fe → 4Cu2O + Fe3O4 2CuO + Zn → Cu2O + ZnO Cu2S + Fe → 2Cu + FeS

Vibratory 8:1 Planetary 30 : 1 Vibratory – Vibratory 8:1

20 153 44 –

BPR

Induction time (min) Reference Yang and McCormick, 1993 Takacs, 2002 Shen et al., 1992 Takacs, 2002 McCormick et al., 1989

important. Moreover, too low heat conductivity of the milling medium can hinder formation of hot spots in which the reaction is initiated. Table 9.2 gives induction times (in minutes) for SHS reactions carried out under the given conditions of mechanochemical syntheses. On the other hand, in order to control explosive effects during selfpropagating reactions, that is to suppress their kinetics, it is possible for example to use a process control agent (PCA) such as toluene which does not stop the reduction reaction but allows it to run in a controlled manner (Murty and Ranganathan, 1998; Takacs et al., 2001; Takacs and Mandal, 2001; Takacs, 2002). The mechanochemical syntheses of metallic–ceramic nanocomposite powders carried out in a non-explosive way based on copper–aluminium and nickel–aluminium alloys reinforced by aluminium oxide using their salts instead of oxides are described in Sections 9.5 and 9.6. For comparison, the explosive reactions in metal oxide with active metal systems are also presented.

9.3

Monitoring mechanochemical processes

A comprehensive study of the physical and chemical processes that occur during mechanical treatment by high-energy ball milling appears only to be possible if a reliable identification of solids and the quantitative phase analysis of activated products is made. Because of the complexity of mechanochemical reactions, the nature of the solids obtained closely relates to the milling conditions and so they should be well defined. Moreover, it is very important to determine the factors that influence the activation effects. Another difficulty arises from the fact that the reactions are composed of many successive stages which vary in different

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cases. The experimental methods required to identify and characterize materials synthesized mechanochemically involve not only techniques which are applicable to solids, but also others, more particularly adapted to the nanostructured character of the milling products. Thus, different types of analytical methods must be applied (e.g. Wieczorek-Ciurowa et al., 2000, 2001; Balázˇ, 2008). The best way to study the kinetics of mechanochemical transformations would be to analyse continuously the milled products in situ. However, until now this has been a difficult task. The only method that has been used in some laboratories is the gas pressure–temperature measuring system (GTM), to acquire in situ data during planetary ball milling, for example Fritsch GmbH mills, which enable indirectly control and provision of synthesis of chemical compounds by selection of appropriate parameters for milling (http://www.fritsch.com). Recording from the beginning of the milling process is especially useful when gaseous products evolve from the system caused by mechanical decomposition of the compounds. It is also useful when the self-propagating high-temperature syntheses occur in an explosive manner owing to local overheating, for example by two impacting balls (e.g. Murty and Ranganathan, 1998; Kwon et al., 2002). In such cases applying less intensive milling could be considered, for example by using lower values of rotation per minute (rpm) and/or ball-to-powder mass ratio and/or the size and material of grinding balls, as well as controlling the milling atmosphere and the type of control agent. Optimization of the milling parameters chosen experimentally must be taken into account when selecting mechanochemical syntheses conditions because not all these variables are completely independent. Solid phase characterization is realized ex situ when a sample of milled powder is picked up from the vial after a defined milling period. Thus, in the case of a material in crystalline form, the X-ray diffraction (XRD) method is advantageous; it is fast and allows the progress of the syntheses to be followed. However, this is not convenient after some stages of milling where the mechanical treatment caused comminution and prolonged activation destroyed the crystallinity of the solid powder particles leading to their complete or partial amorphization and simultaneously to consumption of the initial mixture components, making them undetectable by XRD because the amounts were too small. Thermal analysis (TA) with registration of thermogravimetric (TG/ DTG) and differential thermal analysis (DTA) curves facilitates the determination of an amount of undecomposed and/or unreacted reagents, oxidation/reduction processes and mechanically induced physical transformations in/or between the reagents. The example shown in Fig. 9.5 allows us readily to see that the analysed sample consists of mechanically undecomposed salt, that is Cu–hydroxocarbonate with the remaining amount

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being aluminium. The endothermic effect, at 350°C reveals that the Cu2(OH)2CO3–Al mixture decomposed into CuO, which was then reduced by Al to metallic Cu generating a large amount of heat at about 600°C. This heat may accelerate the alloying of Cu with Al, for example into intermetallic phases. The effect at 520°C can be related to the eutectoid transition of CuAl2 with Al →L at 528°C, according to the phase diagram of the binary Cu–Al system (Massalski, 1992), because the sample used in this thermoanalytical experiment was a product of the early stage of salt–metal mixture milling. In addition to the above-mentioned analytical techniques for solids, versatile scanning and transmission electron microscopy (SEM and TEM) with backscattered electron imaging and quantitative energy dispersive X-ray elemental microanalysis (EDS) are very helpful in estimating the microstructure of synthesized composite particles for a relatively coarse powder and for fine powders. By combining the grey tone levels with the results of EDS some compounds can be identified and localized (Yakowitz, 1975; Wieczorek-Ciurowa et al., 2003c). To illustrate the utility of transmission electron microscopy in estimating the microstructures of mechanochemically synthesized composite powders, Fig. 9.6 shows a set of TEM microphotographs of Cu2(OH)2CO3–Al powder after mechanochemical synthesis showing nanocrystalline copper. Electron diffraction patterns unquestionably confirm that the detected Cu, CuAl2 phases are in the nanocrystalline form. Moreover, it is worthwhile adding methods such as neutron diffraction, magic angle spinning solid-state nuclear magnetic resonance (MAS NMR), Fourier transform infrared (FT-IR), ultraviolet (UV) or X-ray photoelectron spectroscopy (XPS). All these techniques have been described in standard textbooks on solid state chemistry and they are not described in detail here. To determine the real phase constitution with characterization of milled powders, a combination of some methods is necessary.

9.4

Examples of applied high-energy milling in the synthesis of selected metallic–ceramic composite powders

Mechanochemical synthesis is almost an ideal method for preparing nanosized metal matrix composite powders because of its simplicity and the possibility of forming composite powder particles with a uniform distribution of grain sizes. Moreover, such in situ routes of synthesis (Sections 9.5 and 9.6) result in the production of powdered materials that have microstructures that are more homogeneous than those synthesized using conventional ex situ techniques (Section 9.7).

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(b)

250 nm

(c)

1-10 002 2-20

9.6 Set of TEM microphotographs of the Cu2(OH)2CO3–Al system after mechanochemical synthesis showing nanocrystalline copper: (a) bright field, (b) dark field, (c) electron diffraction patterns (TEM Philips CM 20 with EDX) (Wieczorek-Ciurowa and Gamrat, 2007a).

In metallic–ceramic composite powder formation, the preparation of metals and alloys by reducing their salts or oxides with more reactive metals is commonly known as metallothermic reduction. This reaction is expressed in general by the equation: MAX + MB > MA + MBX

[9.2]

where a metal MA is reduced by a more reactive metal MB (reductant) to the pure metal MA. MAX and MBX are oxides, chlorides, sulphides and other salts. Metallothermic reactions are characterized by a large negative freeenergy change and therefore they are thermodynamically feasible at room temperature (see for example reactions [9.4] and [9.5]. Moreover, the

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mechanical activation of reagents significantly increases the synthesis kinetics. A salt such as MA hydroxocarbonate can be used as a ‘source’ of MAO, as described in Sections 9.5 and 9.6.

9.5

Copper-based composite powders with Al2O3

The formation of a copper-based composite powder with aluminium is a consequence of many complex, simultaneous and subsequent chemical reactions that occur in the milling of a Cu2(OH)2CO3 and Al mixture. They are schematically reported below (equations [9.3]–[9.5]) together with the enthalpy values (ΔH298/kJ mol−1; Barin et al., 1977): •

mechanochemical decomposition of salt: +82; Cu2(OH)2CO3(S) → 2CuO(S) + CO2(G) + H2O(G)

•

aluminothermic reduction of CuO to Cu with Al2O3 formation: −1179; 3CuO(S) + 2Al0(S) → 3Cu0(S) + Al2O3(S)

•

[9.3]

[9.4]

mechanical alloying of Cu and Al (in the presence of Al2O3): −8 (Cu9Al4) and −4 (CuAl2); Cu0(S) + Al2O3(S) + Al0(S) → Cu0−Al0 + Al2O3

[9.5]

(S) indicates solid and subscript (G) indicates gas. The product is a metallic–ceramic composite powder. Cu0–Al0 alloy can exist as intermetallics and/or a solid solution. The formation enthalpy of oxide MBX should be higher than that for MAX (see Equation [9.2]). The values for Cu, Ni and Al oxides are given in Table 9.3. From a thermodynamic point of view, the aluminothermic reactions in the systems studied are significantly favoured. In Sections 9.5.1 and 9.5.2, the in situ formation of a nanostructural metallic–ceramics nanocomposite Cu–Al/Al2O3 powder is described starting from two different mixtures, Cu2(OH)2CO3 and CuO with Al, respectively (Wieczorek-Ciurowa et al., 2002a; Wieczorek-Ciurowa and Gamrat, 2007a,b; Wieczorek-Ciurowa et al., 2007a,b; 2008). Composite powders were synthesized by high-energy ball milling in a laboratory planetary mill with vial and balls made of hardened steel. The component systems were prepared as physical mixtures in which the amount of Al was calculated assuming the formation Cu9Al4 with Al2O3 and CuAl2 with Al2O3 as components of the synthesis products. Among the intermetallics in the binary Al–Cu phase diagram (Hansen, 1958) are CuAl2, Cu9Al4, Cu(Al) and Al(Cu) solid solutions. It was found that if Cu9Al4 was the expected

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High-energy ball milling Table 9.3 Formation enthalpy values of selected metal oxides (Kubashewski et al., 1993) Oxide

−ΔH (kJ mol−1)

Al2O3 NiO Cu2O CuO

1675.7 239.7 173.2 161.9

phase, mechanosynthesis in both reagents used brings about the formation of a composite material consisting of Cu(Al) solid solution and Al2O3. The amount of Al dissolved into Cu matrix is considerably higher for the CuO–Al system than for Cu2(OH)2CO3–Al. This is probably caused by the different kinetics of mechanochemical reactions in these two systems. According to Avvakumov et al. (2001), compounds containing groups of atoms with oxygen and/or hydrogen (like acidic or basic salts, hydrates) take part in reactions evolving water and/or volatile compounds. These compounds are characterized by 3–4 times lower hardness in comparison to their anhydrous oxides, probably providing the mechanochemical process with lower mechanical loading. Processes taking place with these types of substances are known as soft mechanochemical reactions. Mechanochemical syntheses where the mixtures are provided with stoichiometric proportions of reagents to form CuAl2 result in the formation of two intermetallic phases, CuAl2 and Cu9Al4, as well as Al2O3. It is evident that in the systems with CuO as a reagent for aluminothermic reaction CuAl2 appears after only 1 hour of milling. This indicates the rapid alloying process of both metals: copper and aluminium. The experimental results are described in detail in Sections 9.5.1 and 9.5.2.

9.5.1 Effects of mechanosynthesis in the Cu2(OH)2CO3–Al reagent system Figure 9.7 shows the XRD patterns (Philips X’Pert diffractometer Cu Kα) of the Cu2(OH)2CO3–Al system mechanically treated at different time intervals (Pulverissette 6 Fritsch GmbH). The results indicate that the initial components of the mixture studied still exist in the system after up to 10 hours of mechanical treatment. However, their intensities reduced significantly in the fourth hour of milling. This indicates that Cu2(OH)2CO3 started to decompose and metallothermic reduction of CuO with Al proceeded gradually with formation of metallic copper. Furthermore, a broad diffrac-

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20 h 6

5

2

5

7 15 h 6 6

3 5

2

2

5

2

1

1

1

4 6

3

4 6

1 1

10 h

1

5

2

4

1 4 1 4h 4

1 1

3

4

3

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1

1

1

Intensity (a u)

2000

4

1500

1

1

1000

4

11 1

4

1

1 1

1 Cu2(OH)2CO3 2 CuO 3 Cu2O 4 AI 5 AI2O3 6 Cu 7 Cu(AI) 11 1 4

4

0h

500 0 10

20

30

40 50 60 2θ (degree)

70

80

90

9.7 X-ray diffraction patterns of the Cu2(OH)2CO3–Al system after different mechanochemical treatment times in a Pulverisette 6 Fritsch GmbH mill (intensities are on the same scale) (Philips X’Pert diffractometer, CuKα) (Wieczorek-Ciurowa et al., 2007b).

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tion peak appears in the range of 2θ = 42° to 45° after 10 hours of milling. Its intensity increases significantly with a lengthening of the milling time. The shape of this X-ray indicates its complexity and the deconvolution confirms the presence of copper in addition to part of the Al2O3 phase. After 15 hours and 20 hours of milling the position of the Cu(1 1 1) peak shifted from 2θ = 43.32° to 43.25°, respectively, while Al peaks completely disappeared. This may suggest that a solid solution of Cu(Al) has been formed as a result of mechanical alloying of two metals, Cu and Al. This is confirmed by an increase of metallic copper lattice parameter from 3.6151 Å to 3.6203 Å (Pearson, 1958). According to Vegard’s law, which states that the lattice parameter of the solid solution is a function of the amount of solute and using parameters for pure Cu and Al, the contents of the Al solute in the Cu matrix for the mechanically treated Cu2(OH)2CO3–Al system are estimated on a level of about 5%. From a thermodynamic point of view, the reactions occurring in the tested system are highly exothermic, for example the enthalpy of aluminothermic reduction of CuO with Al equals −1179 kJ mol−1. However, in the hydroxosalt–active metal system, processes in the mill occur in a controlled manner. Coming back to Fig. 9.3 illustrating the temperature and pressure in a vial during milling process it can be seen that in the first hours of milling the temperature and pressure increase simultaneously. The rise in pressure is due to mechanical decomposition of Cu2(OH)2CO3 to CuO with H2O and CO2 evolving, while the temperature increase is caused by the friction and impact of the balls. When Cu–hydroxocarbonate is completely decomposed, any rise in pressure is observed, whereas only a transitory rise in temperature occurs (Fig. 9.3b). On the T versus milling time curve, one can observe several small rises in the gas temperature. However, all of them take place below 40°C. This indicates that the aluminothermic reduction occurs locally, in a repeatable fashion. Based on these data we can conclude that in this case mechanochemical synthesis takes place in a controlled way, gradually, so CuO is reduced by Al in many steps proceeding locally, in isolated parts of treated powders. This process of aluminothermic reduction may be explained by the fact that the occurrence of a highly exothermic reaction via combustion depends not only on thermodynamics but also on other factors such as the nature of reactants, their mechanical properties, crystalline structures and their stability during milling. The thermoanalytical curves, TG–DTA, presented in Fig. 9.8 (WieczorekCiurowa and Gamrat, 2006) for an untreated mixture of Cu2(OH)2CO3 with Al, and after milling for 20 hours, respectively, indicate that Cu–hydroxocarbonate completely decomposes during mechanosynthesis and aluminium is consumed because of the lack of an endo effect at 640°C. Moreover, any sharp exothermic effect caused by the reaction of CuO with Al is

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Mass (%)

100

TG

Aluminothermic CuO reduction

DTA 90

Thermal decomposition of Cu2(OH)2CO3

AI melting

Temperature change (°C mg–1)

(a)

Mass (%)

(b) TG 100 DTA

Rearrangements in Cu–AI system

90 200

400 Temperature (°C)

600

800

Temperature change (°C mg–1)

80

9.8 TG–DTA curves of the Cu2(OH)2CO3–Al system after mechanochemical synthesis for: (a) 0 hours, (b) 20 hours (SDT 2960 TA Instrument) (Wieczorek-Ciurowa and Gamrat, 2006).

detected on the DTA curve. This means that the aluminothermic reaction is completed during that time. It is interesting that there is no endothermic peak before 450°C, which normally corresponds to the transformation of Cu9Al4 into a CuAl2 phase. This is also confirmed by the results of XRD patterns, indicating that this phase did not form.

9.5.2 Effects of mechanosynthesis in the CuO–Al reagent system Mechanosynthesis in a copper oxide–aluminium system under the same conditions as in the case where Cu2(OH)2CO3 was used as a source of CuO caused, within the first hour of milling, the formation of two intermetallic phases, CuAl2 and Cu9Al4 (Fig. 9.9) (Wieczorek-Ciurowa et al., 2007a, 2008). Lengthening the milling time results in the disappearance of both intermetallics, although with different kinetics. The first, CuAl2, decays

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5 2

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2 5

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4 1 2h

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2 2 1

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Intensity (a u)

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1

2

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4 2 1

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1 Cu 2 AI2O3 3 CuAI2 4 Cu9AI4 5 Cu(AI)

2000 1500 2

2

1000 3 500

4

2 3

4

2

4

12

4

2 14

1h 4

0 10

20

30

40 50 60 2θ (degree)

70

80

90

9.9 X-ray diffraction patterns for a CuO–Al system after different mechanochemical treatment times in a Pulverisette 6 Fritsch GmbH mill (Philips X’Pert diffractometer, CuKα, intensities are on the same scale) (Wieczorek-Ciurowa et al., 2008).

after 1 hour of milling, while Cu9Al4 remains in the system up to 2 hours of milling. Further milling caused the consumption of aluminium from these phases to form a Cu(Al) solid solution. This is confirmed by the shift of the Cu(1 1 1) peak from 2θ = 42.57° to 43.33° observed in the system treated for 20 hours. The calculated lattice parameter of this phase is equal to 3.6564 Å. Comparison of this parameter with its value for pure Cu suggests expansion of the copper lattice by aluminium substitution, forming a solid solution of Al in the Cu matrix. In fact, the content of the Al solute was estimated to be at a level close to 10%. Conventionally Al dissolves in Cu at the temperature of eutectoid transformation (565°C) at 9.4%.

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An increase in the temperature in the milling vial registered by the GTM system indicates that the combustion process occurs during the first hour of milling (see Fig. 9.4). In this case, the reaction between CuO and Al is completed, confirmed by the absence of peaks corresponding to the initial components. Considering phase evolution in the CuO–Al mixed powders, one can suggest that intermetallic phases are formed only as intermediate products under the applied conditions. During further milling they transform to the Cu(Al) solid solution, that is the expected Cu9Al4 phase was not formed in final products. The characteristics of the milling product microstructure were provided using a scanning electron microscopy. Figure 9.10 shows the typical morphology of composite powder after milling for 20 hours. EDS analysis showed that the dark grey phase is Al2O3 while the bright grey one is a Cu(Al) matrix. The microphotograph confirmed that alumina particles are evenly embedded in the Cu–Al matrix, in size range from 100 nm to 500 nm. This indicates that milling the CuO–Al system with a stoichiometry close to Cu9Al4 brings about formation of metal matrix composite powder particles in which Cu(Al) forms a matrix while alumina grains act as its reinforcement.

AI2O3

Cu(AI)

10.0 µm

9.10 SEM microphotograph (BSE image) illustrating that Al2O3 particles (ranging from 100 nm to 500 nm) are embedded in the Cu–Al matrix formed during reactive milling of the CuO–Al system. The dark spots correspond to Al2O3, the brighter grey ones to Cu(Al) (SEM Hitachi S-4700 with EDX Noran Vantage) (Wieczorek-Ciurowa et al., 2007a).

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9.6

Nickel-based composite powders with Al2O3

Formation in situ of composite powders of nickel–aluminium metals strengthened by alumina is very important from both a theoretical and a practical point of view. Analogous to the mechanochemical synthesis of Cu–Al/Al2O3 composite described in Section 9.5, which is based on the type of reactants, three modifications of the process are possible. Nickel in the form of its salt, oxide or simply elemental metal with aluminium can be used. Because nickel and its alloys have good plastic fatigue features but are not sufficiently resistant to high temperature oxidation and abrasive wear, intermetallics such as NiAl and Ni3Al are a better choice for a composite matrix. They are resistant to very high temperature, corrosion, mechanical and abrasive actions and they have a relatively low density. However, their low intergranular cohesion is responsible for their great brittleness. For this reason, a ceramic component is often present in the composite. In Sections 9.6.1 and 9.6.2, in situ mechanochemical syntheses of Ni–Al/ Al2O3 are described, starting from two different mixtures, Ni2(OH)2CO3. xH2O or NiO with Al, respectively (Wieczorek-Ciurowa and Oleszak, 2008). The milling procedure was similar to the one described in Section 9.5. The component systems were prepared as physical mixtures in which the amount of Al was calculated assuming that NiAl with Al2O3 or Ni3Al with Al2O3 formed as final products of the synthesis. Among the intermetallics in the binary Ni–Al phase diagram (Hansen, 1958), the main ones are Ni3Al, NiAl, Ni2Al3, NiAl3 and solid solutions. It was found that if Ni3Al was the expected phase, mechanosynthesis brings about the formation of composite powder particles consisting of Ni3Al matrix and Al2O3. In mechanochemical syntheses in the NiO–Al system where reagent proportions were calculated for the NiAl phase, the expected NiAl and additionally the Ni3Al intermetallic phase were formed as well as Al2O3. However during the mechanical treatment of Ni2(OH)2CO3⋅xH2O with aluminium, the alloying process does not form NiAl but mainly proceeds to Ni3Al formation. Presumably, the reason for this is the different nature and structure of the nickel salt compared with copper salt. Ni-hydroxocarbonate belongs to hydrated salts with low crystallinity (Wieczorek-Ciurowa et al., 2002b; Wieczorek-Ciurowa and Gamrat, 2005). Moreover, crystalline water stabilizes the structure of the nickel salt, making its mechanical decomposition more difficult, in effect slowing down the delivery of NiO in the aluminothermic reaction [9.7]. The examples of two kinds of salts presented here, Ni2(OH)2CO3⋅xH2O, almost amorphous, and Cu2(OH)2CO3 with high crystallinity showed that crystalline structures are more vulnerable to mechanical treatment because their activity becomes higher, for example creating structural defects and new surfaces, and undergoing mass transfer (mixing).

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Special attention should be paid to the important part being played by water vapour in the milling system of Ni2(OH)2CO3.xH2O–Al during mechanical decomposition of salt. The compound contains both crystalline water (x = about three molecules) and OH groups. The evolved water reacting rapidly (ΔH298 = −949 kJ mol−1) with mechanically activated aluminium powder, see Equation [9.6] decreases the amount of Al needed to synthesise the intermetallics (see Section 9.6.1). 2Al + 6H2O → 2Al2O3 + 3H2

[9.6]

Too low a concentration of Al in the desired NiAl phase of the Ni2(OH)2CO3. xH2O–Al system gave Ni3Al intermetallics in the composite powder and a phase with a lower amount of aluminium. Similar to the formation of copper–aluminium matrix/Al2O3 composite particles which are the result of strong exothermic reactions, formation of composite particles of nickel–aluminium/Al2O3 is also typical of a selfpropagating high-temperature synthesis (with a combustible nature). The enthalpy changes for reactions [9.7] and [9.8] are negative (Barin et al., 1977; Kubaschewski et al., 1993; Takacs, 2002): 3NiO + 2Al → Ni + Al2O3, ΔH298 = −955 kJ mol−1

[9.7]

xNi + Al → NixAl where x = 1 or 3, ΔH298 = −117 kJ mol−1 (NiAl) and −153 kJ mol−1 (Ni3Al)

[9.8]

9.6.1 Effects of mechanosynthesis in the Ni2(OH)2CO3.xH2O–Al reagent system The XRD patterns of the Ni2(OH)2CO3.xH2O–Al systems with reagent proportions calculated for NiAl or Ni3Al intermetallics mechanically treated for 20 hours are shown in Fig. 9.11. It is seen that instead of the desired NiAl, the Ni3Al phase is formed with traces of NiAl. Furthermore, up to the fourth hour of milling, reagents such as NiO, Ni and Al still remain and after 10 hours of milling the reaction was still incomplete, about 15% salt was undecomposed and mechanical alloying of Ni3Al was in its first stage. The results of thermal analysis measurements also confirmed these effects (Wieczorek-Ciurowa and Oleszak, 2008). After 20 hours of milling the composite powder is fully formed. Figure 9.11(b) shows the synthesis of Ni3Al/Al2O3 composite powder formed after 20 hours of milling in the system with reactant proportions for the Ni3Al phase. However, after up to 4 hours of mechanical treatment, Ni and Al were present and after 10 hours milling, NiO was still detectable. It is assumed that the slower decomposition kinetics of the hydrate of nickel

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1 NiAI 2 Ni3AI

2

300

3 AI2O3

200 1 2

100

3

Intensity (a u)

3

2

23

1

2 1

1

0 400 (b)

2

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9.11 X-ray diffraction patterns (Philips X’Pert diffractometer, CuKα) for a Ni2(OH)2CO3.xH2O–Al system after 20 hours of mechanochemical treatment (Pulverisette 6 Fritsch GmbH mill) with reagent proportions calculated for intermetallic phase formation of (a) NiAl and (b) Ni3Al (Wieczorek-Ciurowa and Oleszak, 2008).

hydroxocarbonate in comparison to anhydrous Cu2(OH)2CO3 is responsible for this behaviour.

9.6.2 Effects of mechanosynthesis in the NiO–Al reagent system The effect of 20 hours mechanochemical synthesis in the NiO–Al reagent system was to produce a NiAl phase with Al2O3, when the ratio of the masses of the reagents was used to provide NiAl stoichiometry (Fig. 9.12a). A small amount of Ni3Al was still present. The composite powder Ni3Al/Al2O3 is formed explosively in the NiO–Al reagent system (Fig. 9.12b). The SEM microphotograph of the product is shown in Fig. 9.13.

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9.12 X-ray diffraction patterns (Philips X’Pert diffractometer, CuKα) of the NiO–Al system after 20 hours mechanochemical treatment (Pulverisette 6 Fritsch GmbH mill) with reagent proportions calculated for intermetallic phase formation of (a) NiAl and (b) Ni3Al (WieczorekCiurowa and Oleszak, 2008).

9.7

Other possible variants of the synthesis of metal matrix–ceramic composites in Cu–Al–O and Ni–Al–O elemental systems using mechanical treatment ex situ and in situ

It has already been reported that copper and nickel as well as other transition metals and intermetallics with these metals are important matrices in composite materials with fine dispersion of Al2O3 particles because of their applications. For instance, cermetals such as Ni–Al/Al2O3 are candidates for military and civil aero engines as exhaust nozzle materials. This type of composite is used in dentistry as a substance for fillings, in prosthetics, as

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9.13 SEM microphotograph illustrating the product structure of a combustion aluminothermic reaction during mechanical treatment of a NiO–Al mixture (SEM Hitachi S-4700) (Wieczorek-Ciurowa and Oleszak, 2008).

vacuum tube coatings for solar hot water systems, transistors, capacitors, and so on (see, e.g. Ying and Zhang, 2000a,b; 2003; Grahle and Arzt, 1997; Morsi, 2001; Portnoy et al., 2002). To produce these materials combustive reactions need to be suppressed during mechanochemical synthesis. Ying and Zhang (2000a, b; 2003) show that this can be achieved utilizing a technique in which the Al powder is first diluted by Cu and then the resultant Cu–Al alloy powder is reacted with CuO. The dilution was done through mechanical alloying of Al and Cu into a Cu(Al) solid solution or Cu–Al intermetallics compounds (depending on the amount of Al needed in the mixture). In the next step, powders of Cu–Al alloy and CuO were milled to form a composite structured powder. Finally, this composite powder was heated (<450°C) causing the reaction between CuO and Cu(Al) or Cu9Al4. Subsequent heating at a temperature below 800°C led to formation of Cu and Al2O3. Mechanical treatment enhances the kinetics of both reactions by refining the composite particles structure. Another technique, also a combination of oxidation and mechanical alloying, has been used to produce Cu–Al2O3, metal matrix–ceramic composite powders. With this method, Cu powder is first partially oxidized and then the powder obtained is mechanically alloyed with Al powder to facili-

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tate the aluminothermic reaction of Al2O3 particles formation (Bobrova and Besterci, 1994). Particulate reinforced Cu/Al2O3 metal matrix–ceramics composites can be mechanochemically synthesized by using several methods which include mixing of Cu melt and Al2O3 powder followed by calcination and internal oxidation of Cu–Al alloy powders (Shi and Wang, 1998). The calcination process is limited (Liang et al., 2004) because the Al2O3 particle size has to be large enough to allow effective milling, while internal oxidation can only produce composites with a low volume fraction of Al2O3 particles, such as oxide dispersion strengthened alloys. Oxidation was used in conjunction with mechanical alloying (Ogbuji, 2004). In this process, Cu–Al alloy powder was milled under an oxidizing atmosphere to produce the composite powder. The advantage of this process is that very small oxide particles can be achieved. Another approach that avoids too high activities of the reactants has been made by Venugopal et al. (2005) who introduced toluene as the milling medium (PCA) in copper–alumina composite synthesis by reactive milling of CuO/Cu2O and Al. In fact, the transformation of a combustible reaction to a progressive one yields nanocomposite particles of Cu and Al2O3 with both components with a crystallite size in the range of about 20 nm. Hwang and Lee (2005) demonstrated that Cu/Al2O3 nanocomposite powders with various vol% of alumina as a reinforcement phase have been successfully produced by mechanochemical synthesis in a high-energy attritor mill using mixtures of Cu, Cu2O and Al powder components. It is important to note that in this case the heat generated during milling was removed by a cooling tube coil attached to the outside of the milling chamber. Moreover, in every stoichiometric reaction, excess Cu was added to the system. The role of excess Cu powder was not only as a diluent to control the concentration of Al2O3 but also as a thermal conductor to remove heat from the system during milling.

9.8 •

•

Conclusions

Metallic–ceramic nano-structured composite powders can be produced in situ simply and successfully at room temperature through a mechanochemical route using high-energy ball milling. The subsequent consolidation produces nanocomposites with the desired microstructure and mechanical characteristics. Despite a number of studies having been carried out in different laboratories to synthesize nanocomposite powders, direct comparison and prediction of the course of the process and/or product properties is still very difficult (even when the precursors have a similar character). Using different kinds of mills and procedures for milling influences the input

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of mechanical energy to the treated material. Moreover, until now, the mechanisms and kinetics of mechanochemical syntheses in solid–solid systems are not fully understood. Therefore it seems to be reasonable to carry out experimental studies to determine the progress of reactions and quality of milling products at various milling stages in relation to the milling conditions and to collect the data in order to formulate a theory of mechanochemical processes. In the case of composite mechanosyntheses with a metal matrix it is especially important to inhibit self-propagating explosive reactions and to transform to the progressive ones by wider use of soft mechanochemistry (acidic, basic, hydrated salts) or, for example, different PCA.

Based on the results of two practical examples of syntheses, the transformation of combustion reactions to progressive reactions is shown by applying copper and nickel hydroxocarbonates as composite precursors (even if different in nature) instead of metal oxides, CuO and NiO.

9.9

Acknowledgements

The author would like to thank Professor Yu. G. Shirokov of the Ivanovo State University of Chemistry and Technology, Russian Federation, who in 1995 inspired the study of mechanochemical processes. The Polish Ministry of Science and Higher Education (Projects No. PB 4 T09B 07024 and PB 1 T09B 023 30) and the State Committee for Scientific Research (Grants CUT/C-1/DS) provided financial support (2002–2008).

9.10

References

agrawal p and sun c t (2004), ‘Fracture in metal-ceramic composites’, Comp Sci Technol, 64, 1167–78. avvakumov e g, senna m and kosova n (2001), Soft Mechanochemical Synthesis, A Basis for New Chemical Technologies, Kluwer Academic Publishers, Boston. balázˇ p (2008), Mechanochemistry in Nanoscience and Minerals Engineering, Springer, Berlin. barin i, knacke o and kubaschewski o (1977), Thermochemical Properties of Inorganic Substances, Supplement, Springer, Berlin. boldyrev v v (2006), ‘Mechanochemistry and mechanical activation of solids’, Russian Chem Rev, 75 (3), 177–89. boldyrev v v, pavlov s v and goldberg e l (1996), ‘Interrelation between fine grinding and mechanical activation’, Int J Mineral Process, 44–5, 181–5. bobrova e and besterci m (1994), ‘Processing of Cu–Al2O3 metal matrix nanocomposite materials by high-energy ball milling’, Powder Metall Sci Tech, 6, 7. botta p m, aglietti e f and porto lopez j m (2000), ‘Thermal and phase evolution of mechanochemical reactions in the Al–Fe3O4 system’, Thermochim Acta, 363, 143–7.

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butyagin p yu (2000), ‘Mechanochemical synthesis: mechanical and chemical factors’, J Mater Synth Process, 8 (3/4), 205–11. butyagin p yu (2003), ‘Diffusion and deformation models of mechanochemical synthesis’, Colloid J, 65 (5), 648–51. chang s t, tuan w h, you h c and lin i c (1999), ‘Effect of surface grinding on the strength of NiAl and Al2O3/NiAl composites’, Mater Chem Phys, 59, 220–4. chen y, hwang t, marsh m and williams j s (1997), ‘Study on mechanism of mechanical activation’, Mater Sci Eng, A226–8, 95–8. courtney t h (1994), ‘Modelling of mechanical milling and mechanical alloying’, Rev in Particulate Mater, 2, 63–116. fernández-bertran j f (1999), ‘Mechanochemistry: an overview’, Pure Appl Chem, 71 (4), 581–6. grahle p and arzt e (1997), ‘Microstructural development in dispersion strengthened NiAl produced by mechanical alloying and secondary recrystallization’, Acta Mater, 45, 201–11. hansen h (1958), Constitution of Binary Alloys, McGraw-Hill, New York. heegn h, birkeneder f and kmptner a (2003), ‘Mechanical activation of precursors for nanocrystalline materials’, Cryst Res Technol, 38 (1), 7–20. http://www.fritsch.com Laboratory planetary mill Pulverisette 6 (Fritsch GmbH, Germany). hwang s j and lee j-h (2005), ‘Mechanochemical synthesis of Cu–Al2O3 nanocomposites’, Mater Sci Eng, A405, 140–6. ivanov e and suryanarayana c (2000), ‘Materials and process design through mechanochemical routes’, J Mater Synth Process, 8 (3/4), 235–44. kaczmar j w (2000), ‘Production and application of metal matrix composite materials’, J Mater Process Technol, 106, 58–67. kelsall r w, hamley j w and geoghegan m (eds) (2005), Nanoscale Science and Technology, J Wiley & Sons, New York. koch c c (2001), Nanostructured Materials Processing, Properties and Potential Applications, Wiliam Andrew, Norwich, New York. kubashewski o, alcock c b and spencer p j (1993), Materials Thermochemistry, 6th edition, Pergamon Press, Boston. kwon y-s, gerasimov k b and yoon s-k (2002), ‘Ball temperatures during mechanical alloying in planetary mills’, J Alloys Comp, 346, 276–81. liang s, fan z, xu l and fang l (2004), ‘Kinetic analysis on Al2O3/Cu composite prepared by mechanical activation and internal oxidation’, Composites, A35, 1441–6. lyakhov n z, vityaz p a, grigoryeva t f, talako t l, barinova a p, vorsina i a, letzko a i and cherepanova s v (2008), ‘Nanocomposites intermetallics/oxides, produced by MA SHS’, Rev Adv Mater Sci, 18, 326–8. massalski t b (ed.) (1992), Binary Alloy Phase Diagrams, 1, 2nd edition, ASM Int., Ohio, USA. maurice d r and courtney t h (1990), ‘The physics of mechanical alloying: A first report’, Metallurgical Trans A, 21A, 289. mccormick p g, wharton v n and schaffer g b (1989), Physical Chemistry of Powder Metals Production and Processing, TMS, Warrendale, PA. morsi k (2001), ‘Review: reaction synthesis processing of Ni-Al intermetallic materials’, Mater Sci Eng, A299, 1–15.

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murty b s and ranganathan s (1998), ‘Novel materials synthesis by mechanical alloying/milling’, Int Mater Rev, 43, 101–41. ogbuji l u (2004), ‘The oxidation behavior of an ODS copper alloy Cu-Al2O3’. Oxidation of metals, 62 (3/4), 141–51. pearson w b (1958), A Handbook of Lattice Spacings and Structures of Metals and Alloys, Pergamon Press, London. portnoy v k, blinov a m, tomilin i a, kuznetsov v n and kulik t (2002), ‘Formation of nickel aluminides by mechanical alloying and thermodynamics of interaction’, J Alloys Comp, 336, 196–201. schaffer g b and mccormick p g (1992), ‘The direct synthesis of metals and alloys by mechanical alloying’, Mater Sci Forum, 88–90, 779–86. schwarz r b (1996), ‘Introduction to the viewpoint set on: mechanical alloying’, Scripta Mater, 34 (1), 1–4. sˇ epelak v (2002), ‘Nanocrystalline materials prepared by homogeneous and heterogeneous mechanochemical reactions’, Ann Chim Sci Mater, 27 (6), 61–76. shen t d, wang k y and quan m x (1992), ‘Solid state displacement reaction of Fe and CuO induced by mechanical alloying’, Mater Sci Eng, A151, 189–95. shi z and wang d (1998), ‘Processing of Cu-Al2O3 metal matrix nanocomposite materials by using high energy ball milling’, J Mater Sci Lett, 17, 477–81. smolyakov v k, lapshin o v and boldyrev v v (2007), ‘Macroscopic theory of mechanochemical synthesis in heterogeneous systems’. Int J Self-Propagating HighTemp Synth, 16 (1), 1–11. suryanarayana c (2001), ‘Mechanical alloying and milling’, Progress Mat Sci, 46, 1–184. suryanarayana c, ivanov e and boldyrev v v (2001), ‘The science and technology of mechanical alloying’, Mater Sci Eng, A304–6, 151–8. takacs l (2002), ‘Self-sustaining reactions induced by ball milling’, Progr Mater Sci, 47, 355–414 (and literature cited herein). takacs l and mandal s k (2001), ‘Preparation of some metal phosphides by ball milling’, Mater Sci Eng, A304–6, 429–33. takacs l, sojka v and balázˇ p (2001), ‘The effect of mechanical activation on highly exothermic powder mixtures’, Solid State Ionics, 141–2, 641–7. tjong s c and chen h (2004), ‘Nanocrystalline materials and coating’, Mater Sci Eng, R45, 1–88. urakaev f kh and boldyrev v v (2000), ‘Mechanism and kinetics of mechanochemical processes in comminuting devices: 2. Applications of the theory. Experiment’, Powder Technol, 107, 197–206. venugopal t, prasad rao k and murty b s (2005), ‘Synthesis of copper-alumina nanocomposite by reactive milling’, Mater Sci Eng, A393, 382–6. wieczorek-ciurowa k and gamrat k (2005), ‘NiAl/Ni3Al-Al2O3 composite formation by reactive ball milling’, J Therm Anal Calorim, 82 (3), 719–24. wieczorek-ciurowa k and gamrat k (2006), ‘Mechanochemical synthesis of nanocrystalline cermets’, Pol J Chem Technol, 8 (3), 51–3. wieczorek-ciurowa k and gamrat k (2007a), ‘Some aspects of mechanochemical reactions’, Mater Sci Pol, 25 (1), 219–32. wieczorek-ciurowa k and gamrat k (2007b), ‘Mechanochemical syntheses as an example of green processes’, J Therm Anal Calorim, 88 (1), 213–17. wieczorek-ciurowa k and oleszak d (2008), ‘Mechanochemical synthesis of Ni– Al/Al2O3 and Cu–Al/Al2O3 nanocomposite powder’, 15th International Sympo-

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sium on Metastable and NanoMaterials, ISMANAM2008, Buenos Aires, 26–30 June 2008, Argentina, Book of Abstracts. wieczorek-ciurowa k, shirokov ju g and paryło m (2000), ‘Use of thermogravimetry to assess the effect of mechanical activation of selected inorganic salts’, J Therm Anal Calorim, 60, 59–65. wieczorek-ciurowa k, paryło m, shirokov ju g, gamrat k (2001), ‘The characteristics of mechanically activated mixtures of copper hydroxocarbonate with aluminium’, J Therm Anal Calorim, 65, 359–66. wieczorek-ciurowa k, gamrat k, paryło m and shirokov ju g (2002a), ‘Alloy formation in mechanically activated mixtures hydroxocarbonate with Al0’, J Therm Anal Calorim, 70, 165–72. wieczorek-ciurowa k, gamrat k, paryło m and shirokov ju g (2002b), ‘The influence of aluminium and aluminium oxide on the effects of mechanical activation of nickel hydroxocarbonate’, J Therm Anal Calorim, 69, 237–43. wieczorek-ciurowa k, gamrat k and shirokov ju g (2003a), ‘Chemical reactions during high energy ball milling of the Cu2(OH)2CO3–Al system’, Solid State Ionics, 164, 193–8. wieczorek-ciurowa k, gamrat k and shirokov ju g (2003b),’Mechanism of mechanochemical reactions in malachite–active metal systems’, Thermochim Acta, 400, 221–5. wieczorek-ciurowa k, gamrat k, fela k and sawłowicz z (2003c), ‘Structural and morphological changes induced by mechanical activation in copper hydroxocarbonate–aluminium system’, Pol J Chem Technol, 5, 72–4. wieczorek-ciurowa k, gamrat k and sawłowicz z (2005), ‘Characteristics of CuAl2-Cu9Al4/Al2O3 nanocomposites synthesized by mechanical treatment’, J Therm Anal Calorim, 80 (3), 619–23. wieczorek-ciurowa k, oleszak d and gamrat k (2007a), ‘Mechanochemical synthesis of Cu–Al/Al2O3 composite in CuO–Al system under different conditions’, Chem for Sustainable Development, 15, 255–8. wieczorek-ciurowa k, oleszak d and gamrat k (2007b), ‘Mechanosynthesis and process characterization of some nanostructured intermetallics–ceramics composites’, J Alloys Comp, 434–5, 501–4. wieczorek-ciurowa k, oleszak d and gamrat k (2008), ‘Cu-Al/Al2O3 cermet synthesized by reactive ball milling of CuO-Al system’, Rev Adv Mater Sci, 18, 248–52. yakowitz h (1975), ‘Methods of quantitative X-ray analysis used in electron probe microanalysis and scanning electron microscopy’, In: Practical Scanning Electron Microscopy, Goldstein J I and Yakowitz H (eds), Plenum Press, New York. yang h and mccormick p g (1993), ‘Combustion reaction of zinc oxide with magnesium during milling’, J Solid State Chem, 107, 258–63. ying d y and zhang d l (2000a), ‘Processing of Cu-Al2O3 metal matrix nanocomposite materials by using high energy ball milling’, Mater Sci Eng, A286, 152–6. ying d y and zhang d l (2000b), ‘Solid-state reactions between Cu and Al during mechanical alloying and heat treatment’. J Alloys Comp, 311, 275–82. ying d y and zhang d l (2003), ‘Solid state reactions between CuO and Cu(Al) or Cu9Al4 in mechanically milled composite powders’, Mater Sci Eng, A361, 321–30. zhang f, kaczmarek w a, lu l and lai m o (2000), ‘Formation of Al–TiN metal matrix composite via mechanochemical route’, Scripta Mater, 43, 1097–102.

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11 Mechanochemical plating and surface modification using ultrasonic vibration S. V. K O M A R O V, Nippon Light Metal Co. Ltd, Japan; S. E. R O M A N K O V, N. H AYA S H I and E. K A S A I, Tohoku University, Japan

Abstract: This chapter introduces the concept of a novel ultrasonicassisted mechanochemical technique for coating metallic surfaces with other metal or ceramic materials and presents recent experimental results on its implementation in a number of material systems. In the first part the capability of the process to deposit micro-sized Ti particles on Al substrates is examined. Based on the results obtained, the process was modified to improve the degree of surface coverage. Then, it was applied to produce La-monazite coatings on stainless steel substrates. The last part of the chapter presents data on TiN nanocoatings deposited on Al, Ti and stainless steel substrates. On the whole, the results suggest that the proposed technique gives much better results when dealing with submicro and nanoparticles. Key words: mechanical coating, ultrasonic vibration, ball treatment, acoustic streaming, nanocoatings.

11.1

Introduction

Recent years have seen considerable development in coating technologies. Many of the more recent developments have been driven not only by the need to increase performance and functionality of the coatings, but also by the need to reduce the process cost and to meet new environmental regulations. Another noteworthy trend in the recent coating technologies has been the growing interest in nanocomposite and nanostructured coatings. The published results have strongly suggested that nanocoatings have the potential to be used in a variety of industrial applications. The vast number of processes involved in the coating technology makes their classification rather difficult. Roughly, the processes can be subdivided into those which proceed under (i) relatively high vacuum and temperature, (ii) normal atmosphere and higher temperatures, (iii) high vacuum and near room temperatures and (iv) normal atmosphere and near room temperature. Detailed information about these and other processes can be found in a variety of reviews, for example (Choy, 2003; Neuville and Matthews, 2007; Pogrebnjak and Tyurin, 2005) and books, for example (Cavaliero et al., 2006). 251 © Woodhead Publishing Limited, 2010

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Each process has its particular advantages and shortcomings. For instance, among the shortcomings of the processes proceeding under vacuum and at higher temperatures is that they require sophisticated equipment which is still rather expensive. Therefore, these processes are faced with the constant challenge to reduce the product costs. In addition, the vacuum chamber restricts the area of specimens to be coated and coatings produced often have a higher level of porosity. The coating processes which can be implemented under a normal atmosphere and elevated temperature have a much higher productivity and yield. However, the coatings quite often suffer from a low coating-to-substrate bond strength because of unavoidable oxidation of the coating material and/or substrate by atmospheric air. Furthermore, production of nanocoatings in these processes is very difficult because of rapid grain growth in the coating occurring at elevated temperatures. The above-mentioned circumstances require a careful choice of materials for the coatings and substrates, accurate and laborious preparation of the surface to be coated and careful control of the coating conditions. As a result, there are many materials the use of which could provide significant improvements in the coating performance, but deposition of such coatings remains a great challenge owing to these problems. In order to overcome these problems, it has been proposed that the principle of ball milling (BM) and mechanical alloying (MA) can be applied to the coating production. It is well known that MA offers greater flexibility in choosing materials systems because formation of new compounds occurs through direct synthesis achieved during fine grinding of the powdered materials in ball mills. One of the first attempts to apply this principle to coating deposition was made by Torosyan et al. (2002). In the method proposed, components to be coated were put together with milling balls and coating powdered materials in a ball mill jar, which was set in vibration motion. It was demonstrated that well-adhered work-hardened layers of Ti, Cr and Cu coatings on steel and Al substrates can be obtained after ball milling and impact treatment. In other studies (Romankov et al., 2006a), experimental results showed that dense thick coatings with good adhesion can be produced by ball treatment of Ti components with Al powder with consequent annealing for intermetallic phase formation. The results of these and some other experiments have revealed that the ball treatment is potentially applicable for producing coatings. However, the MA process in its conventional implementation, that is using treatment inside vials, has a number of disadvantages that restrict its practical application for coating production. First, the size of components for coating is limited by the vial dimension. Second, the time taken for coating deposition is rather long. Recently, we have proposed a novel method of coating and surface modification, which is also based on the principle of BM and MA pro-

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cesses, but allows us to treat wider surface areas for a shorter time (Komarov et al., 2007a, b). Because the method uses the energy of ultrasonic oscillations, we called it UMCA (ultrasonic mechanical coating and armouring). The proposed method possesses a number of attractive features that may make it useful for many practical applications. First, the method has no limitation on the area of surfaces that can be processed. Second, processing can be performed under room temperature and atmospheric pressure by using relatively cheap ultrasonic-based equipment. Third, the potential of the method includes, but is not limited to, the coating treatment. Depending on the operation conditions, the method allows armouring or reinforcing metal surfaces with particles of other metal or ceramic materials which are expected to be capable of significantly improving the performance characteristics of the metal surfaces. Fourth, in addition to coating and armouring, the method can provide rapid cold-hardening and formation of an ultrafine-grained structure in the coated surface layers in the same manner as does the ultrasonic shot peening treatment. This chapter introduces the proposed method and presents recent experimental results on its implementation in a number of material systems. In the first experiments, we investigated the capability of UMCA to deposit micro-sized Ti particles on the surface of Al. Based on the results obtained, the treatment procedure was modified to improve the surface coverage degree. Then, we applied the treatment to produce La-monazite coatings on stainless steel substrates. Most recent investigations have dealt with hard TiN nanocoatings deposited on Al, Ti and stainless steel substrates.

11.2

Key concepts

Generally, the concept of the UMCA method is based on two phenomena: (i) the ability of ultrasonically vibrating solid surfaces strongly to accelerate small balls and particles and (ii) the ability of high frequency ultrasound to generate a steady flow known as acoustic streaming. Depending on the process modification, either the first or both the possibilities can be employed. Below is more detailed description of the process concept. The process concept is schematically represented in Fig. 11.1. Balls made of a hard material (stainless steel, ZrO2 or other) are put into a bowl-shaped resonant chamber which is fixed beneath the specimen surface to be coated or armoured with powdered materials. The resonant chamber is designed in such a way that its bottom and walls can perform vertical, radial and torsional vibrations at a resonant frequency. In the first modification of the process, the powdered material was also put into the chamber. In the second, the powder layer was loosely deposited on the surface of the specimen

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Load

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11.1 Representation of the UMCA process.

before fixing it above the resonant chamber. The specimen chamber is rigidly fixed at the top of a vertically vibrating rod of a water-cooled ultrasonic transducer (not shown in the figure). An important feature of the UMCA process is that ultrasonic oscillations occur with very small amplitudes but with a huge acceleration. The small oscillation amplitudes provide an important advantage for the UMCA process over conventional MA methods used for coating. Small oscillation amplitudes make it possible to fix the resonant chamber beneath a specimen so that the gap between the specimen surface and the chamber top edge can be as small as several tenths of a micrometre. This enables the chamber to be moved over a wider specimen or, vice versa, the specimen can shift relative to the chamber, which might be necessary to treat a wider surface area, at the same time avoiding balls jumping out of the chamber during the treatment. However, at the present stage of our study, we investigate a simplified version of the process in which both the chamber and surface are fixed relative to each other. High-intensity oscillations result in vigorous acceleration and chaotic motion of the balls inside the resonance chamber leading to their colliding with each other and with the specimen surface. Moreover, when the powder is put inside the chamber, its vibrations force the particles to jump up from the chamber bottom where they lie initially. Large-scale acoustic streaming then entrains the particles into an eddy motion inside the chambers which causes some of the particles to transfer to the specimen surface and adhere on to it. Two closed loops in Fig. 11.1 show one of the possible patterns of large-scale acoustic streaming inside the chamber. Ball collisions with these © Woodhead Publishing Limited, 2010

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adhered particles (first modification) or with loosely deposited particles (second modification) cause formation of a coating or armoured layer on the specimen surface. The behaviour of balls and particles in the resonant chamber was considered in detail in our earlier study (Komarov et al., 2007a). In the present investigation, two chambers, one small and one large, were used. The inner diameters and heights of the chambers were 30 mm and 60 mm for the small one, and 85 mm and 90 mm for large one, respectively. The chambers were made of Ti alloy. They were excited by using two different generators, a 1-kW ultrasonic generator for the small chamber and a 2-kW generator for the large one. The generators operated at the frequency range of 19–22 kHz. Both the generators had a built-in unit for automatic adjustment of the resonance frequency during the treatment.

11.3

Surface modification and coating treatment of as-obtained substrates

11.3.1 Experimental The ability of the UMCA process to treat as-obtained substrates, that is when particles and balls are put together inside the chamber, was examined with Al plate and Ti powder as substrate and coating material, respectively. All experiments were carried out using the small resonant chamber. Samples that were 50 mm square and 2–3 mm thick were fixed at the top of the resonant chamber for treatment. Two kinds of Ti powder (99.9%Ti) used had median diameters, dp = 28 μm and 95 μm. Figure 11.2 is a schematic drawing of the experimental setup used in the present study. The figure was adapted from our previous paper (Komarov Pressure plate Thermocouple Sample Supporting table

Supporting bars Resonant chamber

Ar gas inlet

Gas outlet Transducer

Acrylic resin box

To ultrasonic generator Water

11.2 Experimental setup.

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et al., 2007a). A sample for coating was secured stationary above the resonance chamber with the aid of a pressure plate and two supporting plates. To measure the sample temperature during ultrasonic treatment, a K-type thermocouple was placed in contact with the plate. The stainless steel balls used in the experiments had a diameter of 3 mm. All experiments were performed inside a sealed box made of acrylic resin. To avoid oxidizing the sample surface and coating materials, Ar gas was blown into the box before and during coating. The processing time ranged from 5–60 min.

11.3.2 Degree of coverage The degree of coating was found to be mainly influenced by the number of balls and quantity of powder, the size of the powder particles and the treatment time. Figure 11.3, which was partly adapted from (Komarov et al., 2007a), shows scanning electron microscope (SEM) photographs of Al plates after 10 min treatment with Ti particles of median diameters dp = 28 μm and 95 μm. In the experiments, approximately 0.5 g of the particles was put into the resonant chamber filled with stainless steel balls. Their total weight was about 10 g, respectively. The photographs were taken using back-scattered electrons. A comparison of Fig. 11.3 (a) and (b) clearly shows that when coarser Ti particles were used for treatment, the coating efficiency was much worse. It is assumed that the particle size dependence is associated with acoustic streaming. Estimates showed that under the present experimental conditions the acoustic streaming is capable of involving relatively small particles in an eddy motion while the relatively large particles are too heavy. It is also noteworthy that a greater proportion of particles on the surface had significantly smaller sizes compared to those of original particles. This suggests that Ti particles were subjected to milling during coating operations. Figure 11.4 presents typical data on the particles size distribution after 10 min treatment. To plot these data, SEM images of the particles were processed by Scion Image software. In order to express the degree of coverage, α, quantitatively, the backscattered electron images were processed by the same software. According to the results, α increased from 12.5% to 56.6% in the range of treatment time from 5 to 10 min. However, a further increase in the treatment time up to 60 min resulted in an increase of α from 56.6% to 71%. This finding suggests that the coating rate is probably not controlled by the collision frequency but instead by the transfer of particles to the substrate surface. The particles are most likely to adhere to the surface during the first few minutes of the surface treatment, whereupon the adhesion rate decreases. This occurs supposedly because the particles remaining in the resonant chamber become flattened. This was confirmed by SEM observations. The

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(a)

400 μm

(b)

400 μm

11.3 Back-scattered electron images of Al surface (black) armoured with Ti particles (white) with median diameters of 28 μm (a) and 95 μm (b).

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Particles on AI surface Particles remaining in chamber Original particles

Frequency (%)

10 8 6 4 2 0 0

40

80

120 160 Particle diameter (μm)

200

240

11.4 Distribution of particle sizes.

flattened particles readily form agglomerates that make their entrainment by acoustic streaming difficult.

11.3.3 Surface morphology and coating microstructure after treatment Two kinds of particles can be readily distinguished on the surface after the coating treatment, namely relatively large and fine particles. The large particles had sizes close to those of original particles. Figure 11.5 shows a typical SEM view of a large Ti particle on the surface of the Al plate. This figure was adapted from our earlier work (Komarov et al., 2007a). The SEM photographs make it clear that the particles were flattened out and hammered into the surface layers of Al sample matrix in such a way that the particle top surface became flush with the matrix surface. The hammering effect is better viewed as a cross-sectional photograph. One is shown in Fig. 11.6 which was adapted from our previous study (Komarov et al., 2007a). The dash-dot line on the figure indicates the approximate location of the Al plate surface. It is readily seen that the Ti particles are deeply incorporated into the Al matrix. In some cases, admittedly as a result of successive hammering of two and more particles into the same place; the first particles penetrated deeply into the Al plate through its surface as exemplified in Fig. 11.6 by the dashed circle.

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Al

Ti

10 μm

11.5 A single Ti particle hammered in the Al matrix.

Al matrix

Ti particles

Sample surface 20 μm

11.6 Back-scattered electron image of a cross-section of Al sample with hammered Ti particles.

Because the particle hammering happens very rapidly, it should occur with a considerable normal and shearing stress at the particle–matrix interfaces. This, in turn, can be followed by cracking and peeling of thin oxide films, inevitably present on the surfaces of such easily oxidized metals as Al and Ti. Moreover, the interface friction during the particle penetration should result in heat release and rise at the near interface temperature. Both these phenomena can provide better contact between particles and matrix.

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Thus, the treatment of as-obtained Al surfaces with Ti particles made it possible to obtain only partly particle covered surfaces. Such surfaces, containing particles embedded in the surface layer, can be termed particlereinforced or armoured surfaces. It is expected that a proper choice of materials for the particles and substrates can give significant improvements in the performance characteristics of these surfaces, especially wear resistance and hardness. Furthermore, the armoured surfaces can be useful as enhanced substrates in production of new functional coatings by other coating technologies. Characterization of the armoured surfaces will be topic of our further investigations. Measurement of the microhardness of Ti particle-armoured Al surfaces showed that the microhardness of the as-obtained surface was 40 HV. After treatment for 10 min, the value of HV increased, on the average, up to 59 HV for areas between the large Ti particles and up to 150 HV for the large Ti particles themselves. It should be noted that the surface hardening in the areas between the large Ti particles might occur not only owing to the presence of micrometre or submicrometre-sized particles at the surface layers but also because the ball treatment increases the surface dislocation density and imparts a compressing residual stress to the surface layers.

11.4

Surface treatment of precoated substrates

11.4.1 Material systems and treatment conditions In order to raise the degree of coverage, we tried to apply the UMCA method to treatment of precoated surfaces. This section presents data on Ti and LaPO4 coatings produced on Al and stainless steel substrates, respectively. The Al/Ti system was chosen because it has been used for treatment of as-obtained surfaces. Choice of monazite-type lanthanum phosphate, LaPO4 for the coating was made for the following reasons. In recent years, lanthanum phosphate has attracted great interest owing to its unique high temperature stability and chemical inertness. Potential applications of lanthanum phosphate include areas such as thermal barrier coatings for materials used in extreme environments (Keller et al., 2003). However, its thermal and chemical stability makes it difficult or even impossible to produce lanthanum phosphate coatings by conventional techniques. In the present experiments we attempted to deposit the LaPO4 coatings on stainless steel samples. These coatings are of considerable industrial interest for application to components working under very high temperatures and a chemically aggressive environment. Experiments were carried out at room temperature under an Ar atmosphere using a small ultrasonic installation with a resonant chamber 30 mm in diameter. Before treatment with impact balls, powder particles were preliminarily deposited on the specimen surface using a particle-containing © Woodhead Publishing Limited, 2010

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ethanol suspension. This operation is termed ‘precoating’. Details of the precoating procedure can be found in our earlier paper (Komarov et al., 2008). After evaporating ethanol, the substrate was fixed at the top of resonance chamber for coating treatment. The treatment time was 3, 5 or 10 min per coating operation. The number of coating operations, Nc ranged from 1 to 3. In the case where Nc > 1, precoating the sample surface with the particle-containing ethanol suspension was done after completing each coating operation. In the Ti/Al system, stainless steel balls were used for impact treatment. Their number and size as well as treatment conditions were the same as described above. In the LaPO4/steel system, the specimens were treated with either stainless steel or ZrO2 balls. Most of the experiments were performed at room temperature in an Ar atmosphere. Some samples were precoated with the LaPO4-containing ethanol suspension after preliminary treatment of the as-obtained substrate surface with balls. The purpose of these experiments was to elucidate whether there is an effect of preliminary mechanical activation of the substrate surface on the coating efficiency. In addition, in part of the experiment, in order to activate LaPO4 deposition, stainless steel specimens were preheated after precoating but before ball treatment. The temperature varied from 150°C to 400°C. Further details of experiments with LaPO4 powder can be found in our previous studies (Komarov et al., 2008; Romankov et al., 2008).

11.4.2 Ti coatings on Al plates The experiments showed that ball treatment of precoated surfaces makes it possible to produce complete Ti coatings. The best results were obtained when the ball treatment process was carried out in two or three stages, each 10-min long. Each stage was preceded by a precoating operation as described above. A SEM view of such a fully coated surface is presented in Fig. 11.7. The coating thickness was estimated to be less than about 30 μm, as can be seen from a cross-section view of the coating in Fig. 11.7 (b). This figure also shows that the Ti particles, embedded in the Al matrix, are strongly deformed and their surfaces are closely adjacent each other. Additional experiments revealed the formation of a Al3Ti layer at the Al–Ti interface after annealing at 640°C for 120 min under an Ar atmosphere. According to EPMA (electron probe microanalysis), the thickness of Al3Ti layer was about 6–7 μm.

11.4.3 LaPO4 coatings on stainless steel plates Degree of coverage and coating microstructure Experimental results revealed that the degree of coverage and LaPO4 coating uniformity are affected by the number of coating operations, Nc, the © Woodhead Publishing Limited, 2010

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(b)

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Al 30 μm

11.7 SEM images of an Al surface entirely coated with Ti particles (a) and cross-section of the Ti coating (b).

ball material and the ball pretreatment. The degree of coverage and uniformity were estimated by EPMA (JEOL, JXA-8200). Oxygen was used as an element for the estimates. After one coating operation, the coating was rather non-uniform. However, the coverage degree significantly improved after repeating the coating operation three times, each for 10 min. Figure 11.8, which was adapted from our earlier work (Komarov et al., 2008), shows the coatings produced with ZrO2 (a) and stainless steel (b) balls in three coating operations. The light areas in the images correspond to the surface regions covered by LaPO4. It is readily seen that the stainless steel balls yield much lower degrees of coverage compared to the ZrO2 balls. An explanation for this has been given in our previous study (Komarov et al., 2008). Therefore, further experiments were conducted with ZrO2 balls. The coating thickness was determined from SEM cross-section micrographs. To prepare the cross-section samples, a cross-section polisher (JEOL,

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11.8 EPMA mapping of samples coated with LaPO4 in three coating operations: (a) ZrO2 balls, (b) STS balls.

Coating

Substrate

2 μm

11.9 A cross-section SEM view of a LaPO4 coating on a stainless steel surface.

SM-09010) was used. Figure 11.9 illustrates typical SEM views of crosssection samples obtained after a three-step coating procedure at room temperature. This figure was adapted from our previous work (Komarov et al., 2008). The solid line shows the location of the boundary between the substrate and coating. As can be readily seen, the coating thickness is about 2 μm. Interestingly, the coating consists of three layers relating to the

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number of coating operations. The layer thickness varies from several hundred nanometres to approximately one micrometre. The coating itself and the coating-to-substrate interfaces are fully dense; no pores or voids are observable. Further investigations showed that the coating characteristics can be further improved when the coating treatment is conducted at higher temperatures. Figure 11.10, which was adapted from our previous work (Romankov et al., 2008), depicts cross-section SEM views of coatings obtained at different temperatures, 200°C, 300°C and 400°C. At 200°C, the coating consisted of 100–350 nm particles and included many pores and planar microcracks. The larger particles were probably formed as a result of consolidation of smaller ones. Coatings with the maximum thickness were formed at 300°C. Fine nanostructured particles of about the same size as those at 200°C are clearly seen in the figure. However, more careful X-ray diffraction (XRD) investigations revealed that the crystallite size of the LaPO4 coatings varies with the treatment temperature. When the treatment temperature rose to 400°C, the granular structure was not observed by SEM (see Fig. 11.10). The coating structure became less porous and cracked than that produced at lower temperatures. Characterization of LaPO4 coatings The coatings produced were tested for microhardness and adhesion. Figure 11.11 is a plot of (a) microhardness and (b) peeling force versus treatment temperature. This figure was adapted from our earlier work (Romankov et al., 2008). Each point was obtained by averaging at least three measurements. For comparison, the microhardness of the uncoated stainless steel substrate after ball treatment is also shown in the figure by open circles. The room temperature measurements relate to LaPO4 coatings obtained after one coating operation. The data reveal that both the microhardness and peeling force increase with the treatment temperature. Specifically, a considerable rise in the microhardness of the coatings can be seen at temperatures starting from 250°C. It is worthy of note that the microhardness of the uncoated substrates increased from about 200 HV0.3 in as-obtained state to 367 HV0.3 after ball treatment at room temperature. This result is attributed to the effect of cold hardening produced by ball-to-surface collisions in much the same manner as in the ultrasonic shot peening treatment (Abramov et al., 1998). It is specially noteworthy and interesting that the microhardness of the uncoated substrates also increased with the treatment temperature. Additional XRD investigations suggested that ball treatment of preheated stainless steel substrates promotes an α-Fe → γ-Fe transformation in the surface layer of the substrates (Abramov et al., 1998). It is assumed

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300°C

400°C

1 μm

11.10 Cross-section views of LaPO4 coatings deposited at different temperatures.

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500 450

LaPO4-coated samples

400 Ball-treated non-coated samples

350 300

0

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Temperature (°C) (b)

Peeling force, p (kN m–1)

3.8 3.3 2.8 2.3 1.8

0

100

200

300

400

Temperature (°C)

11.11 (a) Microhardness and (b) peeling force as a function of treatment temperature. Coated samples have filled symbols.

that this is the main reason why heating the substrate leads to a rise in the microhardness. Moreover, this provides an explanation of the above results. The hardness of LaPO4 is about three times higher than that of the asobtained stainless steel substrate but more than half that of ZrO2. Therefore, during the first moments of the ball treatment, the harder LaPO4 particles, mainly with larger sizes, can be hammered or pressed into the surface layers of steel substrates similarly to the Ti particles discussed above. Then, as the ball treatment proceeds, the hardness of the LaPO4 particle-coated substrate increases. In this case, the hardening effect rises as the temperature increases. This should promote a rapid densification of the particles in the next stages of the coating treatment. Furthermore, the higher the substrate temperature, the faster the rate of diffusion phenomena occurring in the coating layer and at the coating–substrate interface.

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Obviously, this should result in an increase in the adhesion force between coating and substrate. All this taken together gives an explanation why higher ball treatment temperatures are favourable for LaPO4 coating deposition. In addition, the as-coated stainless steel specimens were subjected to hot corrosion tests in a Na2SO4–NaCl melt at 950°C for 50 hours. The test results showed a significant improvement in the hot corrosion resistance. Details can be found in the previous study (Komarov et al., 2008).

11.5

Production of nanostructured coatings

As mentioned in the chapter introduction, nanocoatings are promising in various areas of technology development. However, the deposition of nanocoatings presents a very difficult problem for many coating techniques because of the inevitable growth of nanograins during the coating deposition. In the UMCA method, the use of nanoparticles for nanocoating deposition seems to be beneficial for the following reasons. First, because the ball treatment is performed at room temperature and under relatively mild impact conditions, it is possible to avoid nanograin growth. Second, nanoparticles possess a huge surface energy and thus are thermodynamically metastable. This forces them to attach to each other to reduce the surface energy. This factor can be of great importance in the first stage of ball treatment when there is a need to hold the nanoparticles in contact with each other and with the substrate surface. In this section, we introduce an application of the UMCA process to deposit TiN nanocoatings on the surfaces of Al, Ti and stainless steel substrates. The choice of TiN is associated with the fact that TiN coatings and their production routes have been investigated intensively in the past. Therefore, a considerable body of data has been accumulated concerning these coatings which facilitates comparison of the coating characteristics.

11.5.1 Treatment conditions Experiments on TiN powders as coatings were carried out at room temperature under ambient pressure in air without any preparation of the substrate surface. A large 2 kW ultrasonic installation was used. About 125 g of ZrO2 balls, 3 mm in diameter, were put into the resonant chamber. The TiN particles used had median sizes of 50 nm. A particles-containing ethanol suspension was applied on to AA1050, pure Ti or SUS304 plates of 100 × 100 mm in size and 2–3 mm in thickness in the manner described above. There were three coating operations, with a ball treatment time per coating operation of 2 min. After the treatment, all samples were rinsed with alcohol and cleaned ultrasonically in a water bath.

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11.5.2 Surface morphology and coating nanostructure after treatment Figure 11.12 presents typical cross-section SEM views of the TiN coatings on the Al and SUS304 substrates. As can be seen, the coatings are dense, crack-free and well-adhered to the substrates. The majority of the TiN particles are in their original form with sizes of around 50 nm, indistinguishable by SEM. However, a significant part of the particles has much bigger sizes in the submicrometre range. These particles are evenly distributed throughout the coating layers and can be readily observed in the SEM views. Although preliminary transmission electron micrograph (TEM) observations revealed the presence of submicrometre particles in the original state (see Fig. 11.13), their sizes did not exceed 300–400 nm. Moreover, these

(a)

Al

(b)

SUS

1 μm

11.12 Cross-sectional SEM views of TiN coatings on (a) Al and (b) steel substrates.

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200 nm

11.13 TEM image of original TiN particles.

particles had triangular, rectangular or hexagonal cross-sections. In contrast, many submicrometre and micrometre particles in the TiN coatings were rounded, especially those which were deposited on the stainless steel substrate. Therefore, it is assumed that they formed as a result of consolidation of the nanoparticles during the ball impact treatment. Very coarse particles can also be observed in both SEM images. In the Al/TiN system, a particularly large particle appears as a single dense structure. The SEM view reveals that the large particle was deeply embedded in the Al matrix presumably owing to repeated ball impacts. This resulted in the formation of cavities which then filled up with TiN nanoparticles. In the SUS/TiN systems, a coarse particle appears as a cluster consisting of several finer particles (see Fig. 11.12b). Their arrangement suggests that they were most likely to be formed through fragmentation of a single large particle as a result of the ball impacts. The cluster is seen to be also incorporated into the metal matrix, although not as deep as in the Al matrix. Careful observation of the interface between the coatings and substrates suggests that, at the interface, many submicrometre-size TiN particles are also incorporated in the metal matrix. The incorporation occurred to a greater extent in the Al substrate and to a lesser extent in the stainless steel substrate, presumably owing to a large difference between the hardness of Al and steel. Particle incorporation is assumed to play an important role in the coating formation mechanism through which coating particles adhere to substrate. This mechanism is referred to as anchoring. The anchoring provides very high values of adhesion strength caused by direct contact between the anchored particles and the substrate. In the case of nanopar-

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TiN Al

20 nm

11.14 TEM image of the interface between TiN coating and Al substrate.

ticles, the adhesion can be greatly enhanced also due to a very high surface energy of nanoparticles. A TEM image (see Fig. 11.14) provides more details of the TiN nanocoating features. This figure presents a TEM image of a TiN nanoparticle layer and the adjacent area of Al substrate. By convention, the interface between them is depicted by the solid line. As can be seen, the vast majority of TiN nanoparticles have sizes within 20 nm, which is significantly smaller than those in original state. This again suggests that the original nanoparticles were subject to grinding during repeated ball-to-surface collisions. Two arrows indicate locations where some nanoparticles appear to be anchored onto the substrate surface.

11.5.3 Characterization of nanocoatings The TiN nanocoatings obtained were examined for thickness, roughness, microhardness and adhesion. The coating thickness was determined from SEM observations of the specimen cross-sections after polishing with a JEOL Cross Section Polisher SM-09010. The coating thickness was measured in different regions and then the average value was determined. The surface roughness Ra was measured by using a Keyence Color 3D Laser Confocal Microscope VK 9700. Measurements of microhardness were performed using an automated load Akashi Hardness Testing Machine at a

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Table 11.1 Characterization of TiN nanocoatings

Thickness (μm) Roughness Ra (μm) Hardness (HV0.05) Critical load (N) (Nanotec Corp.)

Al

Ti

SUS304

4.9 1.87 143 7.9 8.9

5.0 1.45 327 16.4 26.2

5.1 1.12 483 33.5 41.0

load of 50 g and indentation time of 10 s. A scratch tester (CSEM Revest) was used to evaluate the coating adhesion. During the tests, the load, applied on the diamond tip (tip radius 200 μm), was continuously increased while the tip was moving with a constant speed. The loading rate and speed were 100 N min−1 and 10 mm min−1, respectively. The critical loads, at which the coating was fully delaminated, were defined by optical microscopy and acoustic emission. The characterization results are shown in Table 11.1. The averaged thickness of coatings ranged from 4.9 μm to 5.1 μm and the thickness was almost independent of the substrate materials. The coating roughness of Al substrate was larger than that of steel. The roughness on the Ti substrate was intermediate between them. Conversely, microhardness significantly increased when going from Al to Ti and steel. It is worthy of note that the microhardness of the substrates in the as-received condition was 60, 185 and 255 HV, respectively. Thus, a comparison of these values with those in Table 11.1 reveals that the largest increase in microhardness after coating treatment was obtained for the Al substrate. The critical loads of the coatings increased in the same order as did microhardness: Al < Ti < steel, as can be seen in Table 11.1. The measured values were compared with those of industrial TiN coatings shown in the lower line of the table. They were produced by an ion-plating technique (Nanotec Corp., Japan). The coating thickness was around 3 μm. For the Al substrates, the critical loads were very close. Similar to the Ti and steel substrates, the ion-plated TiN coatings were stronger than those produced by the UMCA technique. These results could be explained by the fact that the ion-plated coatings are approximately twice as thin as the mechanically plated coatings. It is well known that coating adhesion deteriorates when the coating becomes thicker because of increasing internal stress in thick coatings. Among the reasons for this is that the ion-plating technique requires careful preparation of the substrate surface before plating which should result in an improvement in the coating adhesion. As noted earlier, the UMCA plating technique can use substrates without or with minimum preparation. Furthermore, in contrast to the ion-plating method, which

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requires a carefully controlled atmosphere, the UMCA process can take place under ambient pressure and temperature.

11.6

Conclusions and future trends

In this chapter, we have introduced the concept of a novel dry process for armouring and/or coating metal surfaces with other metal or ceramic powdered materials. The principle of the process is close to that of ball milling and mechanical alloying; because it uses the energy of resonant ultrasonic vibrations, we called it UMCA (ultrasonic mechanical coating and armouring). The process is capable of producing armoured layers and various coatings for short times by using relatively inexpensive ultrasonic equipment. The most attractive feature of the UMCA process is that it can be implemented under ambient pressure and temperature without or with minimal preparation of the substrate surfaces. Another advantage of the process is its flexibility in the choice of substrates and coating materials. It is the authors’ belief that the method can be effectively used by itself, for example to produce composite metal–ceramics coatings, and in combination with other surface processing techniques. The capabilities of the process were examined experimentally using Al, Ti and stainless steel substrates and a number of powdered coating materials. In the first modification of the process, the powdered material was put into the resonant chamber together with hard balls. Under these conditions, ball impact treatment of Al substrates with Ti particles resulted in rapid formation of a Ti-particle-reinforced layer on the Al substrate. The appeal of this modification is in its simplicity. The particles can be continuously uploaded into the chamber, which can be shifted over the surface of a wider specimen or, vice versa, the specimen can shift relative to the chamber, which may be necessary to treat a wider surface area. The main shortcoming of this treatment, however, is that it needs careful selection of the particle size. Too coarse or too fine particles have a lower capability to adhere on the specimen surface. Among possible improvements that we are working on now is to apply a magnetic field to attract the particles flying inside the chamber to the specimen surface. In a second modification, the substrates were precoated with a particlecontaining ethanol suspension and were then subjected to coating by ball impact treatment. It was shown that three precoating/coating operations provide a degree of surface coverage of 100% or close to it. However, the coating treatment time and coating thickness depended on the particle characteristics. An approximately 30 μm thick coating of Ti particles (median size 28 μm) was deposited on Al substrates during three coating operations, each 10 min long. The same time was needed to produce Lamonazite coatings of 2–5 μm thick on stainless steel substrates. Notice that

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these coatings cannot be produced by other methods because of the high temperature stability and chemical inertness of La-monazite. When nanoparticles were used for coatings, dense and complete coatings could be produced in much shorter times. For example, dense TiN nanocoatings 5 μm thick were produced on Al, Ti and stainless steel substrates in three ball processing operations, each 2 min long, without any preparation of the substrates. On the whole, the results obtained suggest that the UMCA process gives much better results when dealing with submicro and nanoparticles. We intend to investigate the deposition of nanocoatings further in future. In closing it should be noted that although in the present experiments, the UMCA technique was examined for flat substrates, it can be applied to produce various coatings and armoured layers on components with more complex shapes, for example the outer surfaces of pipes and rods. For more detail, the reader is referred to the patent issued to Komarov et al. (2007b).

11.7

References

abramov v o, abramov o v, sommer f, gradov o m and smirnov o m (1998), ‘Surface hardening of metals by ultrasonically accelerated small metal balls’ Ultrasonics, 36, 1013–19. cavaliero a and de hosson j th m (2006), Nanostructured Coatings, Springer, New York. choy k l (2003), ‘Chemical vapour deposition of coatings’, Progr Mater Sci, 48, 57–170. keller k a, mah t, parthasarathy t a, boakye e e, mogilevsky p and cinibulk m k (2003), ‘Effectiveness of monazite coatings in oxide/oxide composites in extending life after long-term exposure at high-temperature’, J Am Ceram Soc, 86, 325–32. komarov s v, son sang han, hayashi n, kaloshkin s d, abramov o v and kasai e (2007a), ‘Development of a new dry mechanochemical method for coating treatment of metallic surfaces’, Surface Coating Technol, 201, 6999–7006. komarov s v, kasai e, saito f, kanou j and hayashi n (2007b), Equipment and Method for Production of Coatings, Japanese Patent Application 2007-197767 (2007-08-09). komarov s v, romankov s e, son sang han, hayashi n, kaloshkin s d, ueno s and kasai e (2008), ‘Production of LaPO4 coatings using a novel ultrasonically assisted plating technique’, Surface Coating Technol, 202, 5180–4. neuville s and matthews a (2007), ‘A perspective on the optimization of hard carbon and related coatings for engineering applications’, Thin Solid Films, 515, 6619–53. pogrebnjak a d and tyurin yu n (2005), ‘Modification of material properties and coating deposition using plasma jets’, Physics-Uspekhi, 48, 487–514. romankov s, sha w, kaloshkin s d and kaevitser e v (2006a), ‘Fabrication of Ti–Al coatings by mechanical alloying method’, Surface Coating Technol, 201, 3235–45.

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romankov s e, kaloshkin s d and pustov l yu (2006b), ‘Synthesis of Ti–Al coatings on the Ti and Al surface by mechanical alloying method’, Fizika Metallov I Metallovedenie, 101, 65–73 (in Russian). romankov s, komarov s v, hayashi n, ueno s, kaloshkin s and kasai e (2008), ‘Effect of temperature on the deposition of the LaPO4 coatings produced by ultrasonicbased coating process on the steel substrate’, Surface Coating Technol, 202, 4285–90. torosyan a r, tuck j r, korsunsky a m and bagdasaryan s a (2002), ‘A new mechanochemical method for metal coating’, Mater Sci Forum, 386–8, 251–6.

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12 Mechanochemically activated powders as precursors for spark plasma sintering (SPS) processes R O B E RT O O R R Ù, R O B E RTA L I C H E R I, A N T O N I O M A R I O L O C C I and G I AC O M O C AO, Università degli Studi di Cagliari, Italy

Abstract: The influence of mechanochemical activation of reactants on the simultaneous synthesis and consolidation of different materials by spark plasma sintering (SPS) is examined. The state of the art in this field is first described in the introductory section, where several examples of bulk products obtained following this approach are given. The main results relating to the fabrication of dense nanostructured MoSi2, TiC–TiB2 and NbAl3 are then reported and discussed in detail. It is found that milling significantly affects SPS process dynamics, the kinetics of the synthesis process, product relative density and crystallite sizes, which can be reduced down to nanometric levels. Possible future developments of this approach with the purpose of obtaining dense nanostructured materials are finally proposed. Key words: mechanochemical activation, ball milling, reactive sintering, spark plasma sintering.

12.1

Introduction

It is well recognized that cold-welding, microstructure refinement and interface formation processes, which favour the occurrence of structural and chemical transformations towards the synthesis of equilibrium and farfrom-equilibrium phases, represent the main effects induced by mechanical treatment of powders by ball milling (BM) (Suryanarayana, 2001; Beyer and Clausen-Schaumann, 2005). Moreover, ball milling is found to significantly affect powder reactivity by so-called mechanical activation (MA) (Charlot et al., 1999; Takacs et al., 2001; Khina and Formanek, 2006). Chemical and structural transformations that occur during BM clearly depend on the corresponding intensity of mechanical treatment, which is typically tuned through the appropriate selection of milling time and charge ratio, that is the ball to powder mass ratio (Delogu et al., 2003). Starting from powders processed by BM, several densification techniques, for instance hot pressing (HP) (Il Moon et al., 2001; Krasnowski and Kulik, 275 © Woodhead Publishing Limited, 2010

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2003; Zheng et al., 2003; Krasnowski et al., 2007; Tavoosi et al., 2009) and shock consolidation (Korth and Williamson, 1995; Yamasaki et al., 2003) have been utilized with the aim of fabricating bulk advanced materials and particularly nanostructured products. Along with these sintering methods, spark plasma sintering (SPS), which belongs to the more general class of electric current activated sintering (ECAS) techniques, has been also recently proposed (cf. Orrù et al., 2009). As schematically represented in Fig. 12.1, during SPS, powders to be sintered and/or the die containing them are crossed by an electric pulsating current while a mechanical load is simultaneously applied. Other than providing rapid Joule heating, mass transport enhancement through electromigration and the controversial presence of a plasma within the voids surrounding the powder particles are also postulated (Orrù et al., 2009). It is generally found that SPS allows sintering to be performed at relatively lower temperatures, shorter times and with a considerable energy saving with respect to conventional HP (Musa et al., 2009). The milder sintering conditions encountered during SPS are relevant when the final goal is the fabrication of difficult-to-sinter materials or to obtain nanostructured products, since grain growth is significantly limited.

Fixed stainless steel electrode

Graphite block

I Powders

Graphite spacer DC pulse ~ current generator

Graphite die Graphite plunger

Moving stainless steel electrode Mechanical load

12.1 Schematic representation of a spark plasma sintering (SPS) system.

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As reported in Table 12.1, a wide variety of bulk materials have been fabricated to date by combining BM treatment with similar ECAS techniques, mainly SPS. Bulk materials include pure metals, different alloys and intermetallics, carbides, oxides, silicides, metal–ceramic and ceramic–ceramic composites, as well as other systems suitable in several industrial fields like aeronautical, microelectronics, biomedical, and so on. Two different approaches are generally followed when coupling mechanical treatment methods with the appropriate consolidation processes. Namely, milling can be adapted for the mechanochemical activation of a starting mixture that will subsequently be reacted and consolidated (reactive sintering) by ECAS in a single step (Lee et al., 2001; Sannia et al., 2003; Paris et al., 2004). In contrast, the synthesis and densification steps can be conducted separately. In this case, reactants are first completely converted by reactive milling into the desired products, whose microstructure is also simultaneously refined, and the resulting powders are then sintered (Tavoosi et al., 2009). Alternatively, products previously prepared by suitable processing routes other than milling are subjected to comminution and/or crystal size refinement through specific mechanical treatments. The powders obtained are subsequently consolidated (Licheri et al., 2007). In this work, the effect of mechanical activation of the starting reactants on the electric current activated simultaneous synthesis and densification of MoSi2, TiC–TiB2 and NbAl3 by means of SPS is examined and discussed. All systems taken into account in this work are very promising materials in several application sectors owing to their known properties. Specifically, MoSi2 is an attractive structural material for industrial and military applications, owing to its resistance in oxidizing environments at elevated temperatures (Petrovic and Vasudevan, 1999). In addition, TiC–TiB2 composites exhibit superior wear resistance and fracture toughness with respect to their constituent ceramic components (Bhaumik et al., 2000). Finally, niobium aluminides are considered to be optimal candidates in structural applications for use beyond the operating temperatures of conventional nickelbased superalloys (Hanada, 1997). The results obtained will be analysed in view of possible utilization of the adopted processing route as a tool for the consolidation of difficult-to-sinter materials.

12.2

Experimental materials and methods

12.2.1 Mechanochemical activation of starting reactants The raw powders used in the present investigation are commercially available products whose characteristics are reported in Table 12.2. Powders were milled in stoichiometric ratios according to the following reactions:

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Table 12.1 Selected systems investigated in the literature which combine ball milling and electric current activated sintering techniques System

Reference

Al-based alloys • Pure Al • Al–17Si–5Fe–3.5Cu–1.1Mg–0.6Zr • 5083 Al alloy • Al–5 at% Fe • Al–Cu–Fe • Al–10La–4Ni–4Fe • (Al+12.5 at%M)3Zr (M = Cu, Mn) Co-based alloys • Co65Ti20W15 Fe-based alloys • Fe–C • Fe–Co • Fe–Pt • Fe–Cr–Ni–N, Fe–Cr–Mo–N • Fe–18Cr–11Mn • Fe–17Cr–13Ni–2.2Mo–0.8Si–0.02C Ni-based alloys • Ni–43 at% Mo Nb-based alloys • Nb–Al, Nb–Al–W, Nb–Al–Mo, Nb–Al–N Ti-based alloys • Ti–Si–Fe • Ti–Zr–Ni • Ti–10Ta, Ti–10Nb • Ti/Ti silicides Niobium aluminides • NbAl3 Iron aluminides • FeAl-based intermetallics • Fe3Al-based intermetallics • Fe–Al–C Titanium aluminides • Ti50Al50 • TiAl–X, Al3Ti–X (X = Cr, Mn, Fe) • TiAl–Cr • Ti(Al+xMn)3 • Ti–Al (FGM) Ni–Ti Sm–Co Sm–Co–Fe Ti–Si SiC WC B4C Al3BC

Kubota, 2007 Sastry et al., 2004 Ye et al., 2006 Sasaki et al., 2007 Nicula et al., 2007a Choi et al., 2007 Lee et al., 2006a El-Eskandarany et al., 2005 Zhang et al., 2005a Nicula et al., 2007b Gebert et al., 2007 Miura and Ogawa, 2001 Cisneros et al., 2005 Oleszak et al., 2007 De La Torre et al., 2000 Murakami et al., 1999

Matsumoto et al., 2002 Matsumoto et al., 2007 Maziarz et al., 2006 Handtrack et al., 2008 Locci et al., 2007 Paris et al., 2004 He et al., 2006 Minamino et al., 2005 Liu and Liu, 2007 Calderon et al., 2002 Maziarz et al., 2004 Jang et al., 2004 Nagae et al., 2000 Ye et al., 1998 Yan et al., 2008 Sreenivasulu et al., 2008 Handtrack et al., 2006 Yamamoto et al., 2004 Shinoda et al., 2008 Heian et al., 2004 Kubota and Cizek, 2008

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Table 12.1 Continued System

Reference

Ti3SiC2 Ti3AlC2 (Fe0.95Mn0.05)75C25 β–Si3N4 MoSi2 CoSb3 SiAlON BaTiO3 Ba1−xSrxTiO3 WC–Co WC–Co–VC Cu–TiB2 SiCp/Al TiAl–SiC SmCo5–Fe Al2O3–Cu AlN–Cu Al2O3–Ni–Co Cr–Al2O3, Nb–Al2O3 Ni3Al–Al2O3 Mo–ZrC Mo–TaC Fe–Mo–SiO2 TiC–Ni FeAl–Y2O3 Mg2Ni/LaNi5 TiB2–Fe–Al Al2O3–TiC Al2O3–Ti3SiC2 Al2O3–BaTiO3 Bi2Te3–SiC Mo(Si0.75Al0.25)2–SiC Si3N4–C SiC–AlN SiC/BN SiC– RBC (rice bran carbon) TiC–TiB2 TiN–TiB2 ZnS–SiO2 ZrB2–ZrC ZrO2–MgAl2O4 HA/Y–TZP HA–Ag HA–TiO2 (Bi2Te3)0.9–(Bi2−xAgxSe3) CNTs/Fe/Al2O3 LaxCo4Sb12 Fe35Pt35P30, Fe50Pt50

Feng et al., 1999 Yang et al., 2009 Terashima et al., 2006 Xu et al., 2005a Orrù et al., 2001; Sannia et al., 2003 Zhang et al., 2004 Xu et al., 2005b Licheri et al., 2007 Hungria et al., 2005 Cha et al., 2003 Zhu et al., 2003 Kwon et al., 2006 Hong et al., 1992 Lee et al., 1997 Rao et al., 2007 Kim et al., 2001 Lee et al., 2007 Oh et al., 2001 Saucedo-Acun˜a et al., 2007 Meng et al., 2006 Takida et al., 2000 Takida et al., 2004 Libardi et al., 2007 Kobayashi and Ozaki, 2006a Ji et al., 2007 Okumura et al., 2002 Kobayashi and Ozaki, 2006b Zhang et al., 2005b Luo et al., 2002 Zhan et al., 2003 Zhao et al., 2008 Krakhmalev et al., 2003 Xu et al., 2007 Shirai et al., 2006 Kodera et al., 2008 Zhou and Hirao, 2005 Locci et al., 2006a Lee et al., 2001 Kim et al., 2008 Shim et al., 2002 Morita et al., 2005 Guo et al., 2003 Lee et al., 2006b Que et al., 2008 Cui et al., 2007 Yoo et al., 2006 Liu et al., 2006 Gopalan et al., 2005

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Table 12.2 Starting powders used in this investigation Reactant

Vendor

Size

Purity (%)

Mo Si Ti B 4C C (amorphous) Nb Al

Alfa Aesar Alfa Aesar Sigma-Aldrich Alfa Aesar Alfa Aesar Alfa Aesar Sigma-Aldrich

3–7 μm −325 mesh −325 mesh 1–7 μm 0.042 μm (average) −325 mesh −200 mesh

99.95 99.5 99.98 99.4 99.9+ 99.8 99

Table 12.3 Milling devices and summary of the ball milling conditions investigated Materials of vials and balls (number, diameter)

Rotation Mill Milling Charge speed, R frequency, time, ratio, CR (rpm) tM (h) ω (rad s−1)

System

Milling device

MoSi2

Planetary ball ZrO2 vials, ZrO2 0–15 mill (Fritsch balls (50, model 10 mm) Pulverisette 5) High-energy Stainless steel 0–24 shaker mill vials, stainless (SPEX mod. steel ball 8000 (1, 13 mm) CertiPrep) High-energy Stainless steel 0–35 shaker mill vials, stainless (SPEX mod. steel ball (1, 8000 13 mm) CertiPrep)

TiC–TiB2

NbAl3

7–21

250–300

–

1

–

90–130

1

–

90–130

Mo + 2Si → MoSi2

[12.1]

4Ti + B4C + C → 2TiC + 2TiB2

[12.2]

Nb + 3Al → NbAl3

[12.3]

Depending upon the system investigated, ball milling experiments were conducted using different milling devices, media and conditions as summarized in Table 12.3. In order to minimize oxidation during mechanical activation, powder handling, loading and milling were performed under an argon atmosphere. Milling runs generally included repeated cycles of one hour milling followed by one hour cooling.

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Phase identification and crystallite size evolution during the milling process were determined by X-ray diffraction (XRD) analysis. Crystallite size (L) was evaluated using the Williamson–Hall method (Williamson and Hall, 1953) for the MoSi2 and NbAl3 cases and the Halder–Wagner method (Halder and Wagner, 1966) for the TiC–TiB2 system. Thus, the separation of the two contributions by line broadening, namely the refinement of crystallite size and the internal strain, was made possible.

12.2.2 Spark plasma sintering of mechanochemically activated powders The two SPS facilities employed for simultaneously reacting and consolidating the activated reactant powders as described in Section 12.2.1. are indicated in Table 12.4 along with the corresponding experimental conditions. A measured amount of powder mixture was placed in a graphite die which was lined with graphite foil. The die was then placed inside the reaction chamber of the SPS apparatus and the system was evacuated. This step was followed by the application of a mechanical load through the plungers. A DC power supply provides both apparatuses with a pulsed current through the processing powders and the die that contains them. The pulse cycle is 12 ms on and 2 ms off for a total pulse cycle duration equal to about 46.2 ms, while the characteristic time of a single pulse is equal to about 3.3 ms. SPS experiments are initiated by the application of previously set electric current (I) or temperature (T) programmes and a mechanical pressure (P) for a selected time interval (tSPS). The temperature of the external surface of the graphite die was monitored by a pyrometer. Temperature measurement during consolidation was also made with thermocouples inserted

Table 12.4 Spark plasma sintering apparatus (Sumitomo Coal Mining Co. Ltd, Japan) and summary of the consolidation conditions investigated

System

SPS facility

MoSi2

Mod. SPS1050 Mod. SPS515S Mod. SPS515S

TiC–TiB2 NbAl3

Controlled parameter (temperature/ current)

Holding temperature, TD or current, I

Mechanical pressure, P (MPa)

SPS time, tSPS (min)

Current

850–1250 A

63.5

0–20

Current

1100 A

20

0–4

Temperature

600–1000°C

20–40

0–25

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inside a small hole in the lateral surface of the graphite die. Voltage, current, load, sample and displacement were measured in real time during SPS. Both (i) sample consolidation, which takes place gradually during classical sintering processes after heating, in addition to the application of the mechanical load, and (ii) thermal expansion of the sample, of both electrodes, of graphite blocks and of spacers and plungers, all contribute to the variation of the displacement parameter. In particular, the variation of the displacement parameter reflects the evolution of physicochemical phenomena (melting, combustion reactions, etc.) taking place during the process. For instance, when the sintering process is accompanied by a rapid synthesis reaction, as in the systems investigated in this work, the most significant and typically sudden change of sample displacement is an indication of the onset of the synthesis reaction. If the densification level achieved by the product at this point is not satisfactory, the sample is mantained at the sintering temperature for the time interval required until the desired consolidation is obtained. After the process, the sample is allowed to cool before it is removed from the die. The relative density of the product is determined by geometrical measurements and by the Archimedes method. Phase identification and crystallite size analysis of the consolidated samples are made using XRD analyses, following the procedure described in Section 12.2.1 for the characterization of milled powder.

12.3

Results and discussion

12.3.1 MoSi2 system Mo and Si powders were milled either separately or together (co-milled) in a stoichiometric ratio corresponding to MoSi2. The influence of milling time (tM), charge ratio (CR) and rotational speed (R) on the characteristics of the elemental powders milled individually was systematically investigated. As an example, Fig. 12.2 shows the variation of the XRD patterns for Mo powder with a milling time for R = 250 rpm and CR = 7. With increasing milling time, the initially sharp diffraction peaks become broader and their intensity decreases, while relatively modest changes are observed in the intensity and width of the Si diffraction peaks. This fact also holds true when changing the rotational speed in the range 250–300 rpm. However, a significant effect on the peaks of both elemental powders is produced with an increase in CR. The dependence of the calculated crystallite size of Mo, L, on milling time for three CR ratios is shown in Fig. 12.3. An analogous systematic study was conducted using co-milled (1 Mo : 2 Si) reactants. The variation of the related XRD patterns with tM (CR = 7, R = 250 rpm) is shown in Fig. 12.4. This result indicates, within the

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Mechanochemically activated powders

12 h

Intensity (a u)

8h

4h

0h 40

60

80 100 2θ (degree)

120

12.2 X-ray diffraction patterns of Mo as a function of milling time (R = 250 rpm, CR = 7).

120 CR = 7 CR = 14 CR = 21

100

L (nm)

80 60 40 20 0

0

2

4

6

8

10

12

14

Milling time, tM (h)

12.3 Dependence of Mo crystallite size on milling time at different charge ratios (R = 250 rpm).

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High-energy ball milling Mo Si α–MoSi2

Intensity (a u)

9h 7h 5h 3h 1h

20

40

60

80

100

120

0h

2θ (degree)

12.4 XRD patterns of co-milled Mo + 2Si powders as a function of milling time (R = 250 rpm, CR = 7).

detection limit of XRD analysis, that when CR = 7 a product is not formed at milling times less than 9 hours. Furthermore, with increasing milling time, the Si peaks broaden and their intensity decreases rapidly. Si reflections almost disappeared at a point corresponding to the first evidence of formation of the product phase, MoSi2. The reduction in the intensity of the Si peaks with milling is in agreement with previous observations (Schwarz et al., 1992; Ma et al., 1993). It was suggested that this feature is the consequence of Si dissolution in Mo or the formation of the silicide phase. The fact that, in our study, the disappearance of the Si peaks coincided with the appearance of the peaks of the product phase, supports the proposed explanation. Similar behaviour was found when the CR parameter was increased. Figure 12.5 summarizes the results of the conditions under which product phase formation occurs for R = 250 rpm. The milling time required for product formation decreases as CR increases. The dashed line in the figure represents the boundary separating the region of product formation from the region where only reactants are detected by XRD analysis. Moreover, using the Williamson–Hall analysis, it was found that the smallest molybdenum crystallite size (125 nm) before product formation commenced is obtained when tM = 3 h and CR = 14 (R = 250 rpm). However,

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16 Product formation Reactants only

14

Milling time, tM (h)

12 10 8 6 4 2 0

6

8

10

12

14

16

18

20

22

Charge ratio, CR

12.5 Effect of milling time and charge ratio for co-milled Mo + 2Si powders on product formation (R = 250 rpm).

since under these milling conditions the silicon peaks almost disappeared from the XRD pattern, no crystallite size analysis was possible for Si. A change of the rotational speed, R, from 250 to 300 rpm, not reported here for the sake of brevity, was investigated for the case of CR = 7 and tM = 7 hours. A slight decrease in Si peak intensity with an almost insignificant change in the Mo peaks and the beginning of α-MoSi2 formation, were the main effects evidenced. As far as the reactive sintering process is concerned, it was observed that the milling treatment of the reactants significantly affects their reactivity during SPS. In fact, when unmilled powder mixtures of Mo + 2Si were processed under an electric current equal to 925 A, no reaction occurred even after the imposition of this current for 20 min. In contrast, reactant powders milled separately (both at R = 250 rpm and CR = 21, being tM = 12 hours for Si and tM = 8 hours for Mo) reacted in SPS at I = 800 A. From the related results shown in Fig. 12.6(a) and (b), it is seen that an abrupt displacement occurred at about 130 s when the maximum temperature recorded by the pyrometer is approximately 920°C. The XRD pattern of the final product is shown in Fig. 12.6(b) along with that corresponding to the starting mixture. It is evident that the reactants have almost completely converted to α-MoSi2. Only minor amounts of Mo5Si3 are found. Furthermore, the initially broad diffraction peaks of the elemental reactants are contrasted with the narrow

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5 800

4

700

3 2

600

Intensity (a u)

900

500

Mo Si α–MoSi2 Mo5Si3

6 Displacement (mm)

Temperature, T (°C)

(a) 1000

(p)

1 0

50 100 Time, tSPS (s)

0 150

(r) 20

40

80 100 120 60 2θ (degree)

12.6 (a) Temperature and sample displacement changes and (b) XRD patterns for the starting mixture (r) and the final product (p) obtained by SPS (I = 800 A) for Mo (tM = 8 hours; CR = 21; R = 250 rpm) + 2Si (tM = 12 hours; CR = 21; R = 250 rpm) powders milled separately.

peaks of the product, indicating a significant increase in crystallite size during SPS. A more noticeable effect of mechanical treatment on powder reactivity is displayed when processing co-milled reactants. In fact, when (Mo + 2Si) milled powders (tM = 3 hours, CR = 14 and R = 250 rpm) were used, a reaction in SPS takes place only after 1.5–2.5 s from the application of the electric current. This event corresponded to the almost complete conversion of reactants into α-MoSi2. As in the case of powders milled separately, some traces of Mo5Si3 were also found in the end product and a relatively large crystallite size (1–2 μm) was obtained. Thus, under the experimental conditions examined so far, it is apparent that the reaction/consolidation stage was affected by the milling treatment in terms of powder reactivity and process dynamics while the nanometric character of the starting powders was not retained during SPS. Reactant samples co-milled until small amounts of the product are formed in the mixture were then processed by SPS. The XRD results for the spark plasma sintered product obtained for the case where I = 800 A using powders milled at R = 250 rpm, tM = 7 hours, and CR = 14, are shown in Fig. 12.7. The starting powder here contained some molybdenum disilicide, both in the α and β modifications, formed during milling. Similarly to the case of powders co-milled for tM = 3 hours, a steep sample displacement occurred after only 1–2 s from the application of the electric power. The product mainly contained the tetragonal α-MoSi2 with small amounts of the hexagonal β-MoSi2, Mo5Si3 and Mo. However, in this case, the final product

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Intensity (a u)

Mo α–MoSi2 β–MoSi2 Mo5Si3

(p)

(r) 20

40

60

80 2θ (degree)

100

120

12.7 XRD patterns of the starting mixture (r) and of the final product (p) obtained by SPS (I = 800 A) from co-milled (tM = 7 hours; CR = 14; R = 250 rpm) Mo + 2Si powders.

is in the nanometric range with the crystallite size of α-MoSi2 evaluated to be about 140 nm and the relative density is 94.8%. Based on the results reported above, it should be emphasized that, although the separately milled powders were treated under much stronger conditions (tM = 8–12 hours, CR = 21) relative to the co-milled case (tM = 3 hours, CR = 14), a considerably longer period of SPS is required for their reaction. This demonstrates the importance of interface formation during milling relative to crystallite size reduction. Furthermore, when nanometric reactants which did not contain any measurable amounts of the product phase were reacted in SPS, the product was not nanometric. On the other hand, the formation of a product phase during milling was found to have a marked effect on the reaction and on the crystallite size of the product. Thus, when MoSi2 is formed during milling in both the α and β modifications, the product after the SPS reaction was nanometric with a crystallite size of 140 nm.

12.3.2 TiC–TiB2 system The variation of the XRD patterns of the (4Ti + B4C + C) powder mixture as milling time, tM, is increased is shown in Fig. 12.8. It is seen that the

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Intensity (a u)

24 h

12 h

6h

20

30

40

50

60

70

0h 80

2θ (degree)

12.8 XRD patterns of (4Ti + B4C + C) reactants as a function of milling time.

initially sharp diffraction peaks become broader owing to powder crystallite size refinement and internal strain increase induced by BM. In addition, these XRD results indicate that, within the detection limit of this analysis, additional phases are not formed at tM up to 6 hours. First evidence of product formation (TiC) is observed when tM = 12 h. The TiC content clearly increased as tM was extended up to 24 hours. The apparent formation during milling of TiC before TiB2 is consistent with the higher diffusivity of carbon into Ti compared to boron (Lasker et al., 1990). The milled powders were then processed in the SPS apparatus under optimized operating conditions (I = 1100 A, P = 20 MPa, tSPS = 240 s) identified using unmilled powders (Locci et al., 2006b). The sample shrinkage and temperature time profiles corresponding to unmilled and milled (tM = 24 hours) powders, are shown in Fig. 12.9. It is clearly seen that the milling step does not have a marked effect on temperature evolution during the process. On the other hand, the behaviour of the system is significantly different with respect to sample shrinkage variation with time. In particular, the shrinkage profile relating to unmilled powders shows the most significant steep increase (to 2.2 mm) at approximately 70 s. The displacement then increases gradually to reach its maximum value (3.4 mm) at 240 s. Conversely, when the starting powders were mechanically treated for 24 hours,

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7 tM = 0 h tM = 24 h

6 5

1600

4 1200 3 800 2 400

Displacement (mm)

Temperature, T (°C)

2000

0

289

1

0

50

100

150 200 Time, tSPS (s)

250

300

0 350

12.9 Sample shrinkage and temperature time profiles of SPS (I = 1100 A, P = 20 MPa, tSPS = 240 s) outputs for the case of unmilled (tM = 0 hours) and milled (tM = 24 hours) (4Ti + B4C + C) reactants.

an abrupt variation of the shrinkage up to approximately 3.4 mm takes place only after about 20 s from the beginning of the SPS process. The shrinkage then gradually increases up to 5.5 mm at about 180 s. The result of the XRD analysis of the product obtained when unmilled (4Ti + B4C + C) powders are processed at I = 1100 A for 70 s, that is immediately after the occurrence of the rapid sample shrinkage, is shown in Fig. 12.10. While a significant amount of TiC is achieved, only traces of TiB2 are revealed. In addition, secondary phases such as TiB and Ti3B4, along with small amounts of reactants are present in the final sample. Therefore, as reported in previous works (cf. Locci et al., 2006b), the electric current needs to be applied for longer periods of time, up to 240 s, to obtain full conversion. The situation changed significantly when considering spark plasma sintered specimen obtained starting from 24-hour milled reactants. In fact, as shown in Fig. 12.10, the composition of the sample obtained when interrupting the application of the current after 20 s, that is just after the abrupt shrinkage variation occurs, demonstrates the complete conversion of the starting reactants to the desired TiC and TiB2 phases. It is interesting to note that the temperature of the die corresponding to the abrupt shrinkage is very low, equal to about 150°C, when powders milled for 24 hours are pro-

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Intensity (a u)

Ti B4C

TiC TiB2

24 h

0h 33 34 35 36 37 38 39 40 41 42 43 44 45 46 2θ (degree)

12.10 XRD pattern of SPS (I = 1100 A, P = 20 MPa) sample when starting from unmilled (tM = 0 hours, tSPS = 70 s) and milled (tM = 24 hours, tSPS = 30 s) (4Ti + B4C + C) reactants.

cessed, while is about 850°C if unmilled reactants are used. In addition, temperatures are in both cases much lower than the melting point of Ti (1639°C), that is the low-melting-point-component, thus the presence of molten phases can be excluded. Consequently, the rapid changes of the sample displacement manifested during the SPS process indicates that combustion reactions are likely to take place. Taking advantage of the results obtained when starting from unmilled powders (Locci et al., 2006b), some considerations related to the mechanism of formation of TiC–TiB2 by SPS when using co-milled reactants can be made. It was suggested by Locci et al. (2006b) that the formation of the composite from unmilled (4Ti + B4C + C) powders is due to the following reaction path: Ti + C → TiC

[12.4]

5Ti + B4C → 4TiB + TiC

[12.5]

16TiB + B4C → 5Ti3B4 + TiC

[12.6]

3Ti3B4 + B4C → 8TiB2 + TiC

[12.7]

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Specifically, it was postulated that the direct carburization of elemental Ti by carbon, Equation [12.4], is the first step taking place during SPS of TiC–TiB2 from unmilled reactants and corresponds to the greatest shrinkage change. Furthermore, Equations [12.5]–[12.7] are expected to be responsible for the subsequent TiB2 formation, as well as of TiB and Ti3B4, detected before the synthesis reaction was completed. Reactions [12.5]– [12.7] most probably did not evolve under the combustion mode but through a gradual solid–solid diffusion mechanism, since their occurrence is not accompanied by a steep displacement change (cf. Fig. 12.9). On the other hand, when reactants mechanically activated for 24 hours were used, examination of the sample shrinkage (cf. Fig. 12.9) in parallel with product composition (cf. Fig. 12.10b), allows us to conclude that the entire reaction path [12.4]–[12.7] evolves in this case under a combustion regime, thus giving rise to complete conversion of reactants to the desired composite. Figure 12.11 reports the effect of milling time on the density of the final product obtained by SPS. All experiments refer to the same SPS operating conditions, that is I = 1100 A, P = 20 MPa, tSPS = 4 min. It may be seen that as the milling time is increased, the density first increases, reaches a maximum value and then slightly decreases. In particular, the final product density increased from about 97% to about 100% of the theoretical density (4.687 g cm−3) when tM was 6 hours. A further increase in tM led to a slight

Relative density (%)

100

99

98

97

96

0

5

10 15 20 Milling time, tM (h)

25

30

12.11 Effect of milling time on the density of the obtained TiC–TiB2 SPS samples, (I = 1100 A, P = 20 MPa, tSPS = 4 min, theoretical density = 4.678 g cm−3).

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decrease in product density, which becomes about 99.5% when tM = 24 hours. This feature can be explained as follows. As the milling time is augmented, the interfacial area between the reactants increases and the powder particle size decreases. These phenomena induced by mechanical treatment are able to enhance mass transport by diffusion and, consequently, synthesis and sintering processes. However, as seen in Fig. 12.8, some products are also formed as a consequence of mechanical treatment. The presence of these refractory phases, that is carbides and borides, at the interfaces between reactant particles, can hinder the sintering process. From this point of view, the mechanical treatment of the starting powders could lead to SPS products characterized by relatively lower density. It is also possible that increased milling contributed to the formation of agglomerates which make sintering more difficult. The average crystallite sizes of spark plasma sintered product phases with milling time are shown in Fig. 12.12. As the milling time is increased, crystallite sizes start decreasing markedly and keep decreasing at a lower rate down to 50 nm after the starting powders have been mechanically treated for 24 hours. This behaviour is related to the corresponding particles refinement in the milled powders and the formation of interfaces between reactants. These features may induce a relative increase of nucleation sites, compared to unmilled powders, thus favouring nucleation with respect to grain growth. In addition, the presence of nanostructured reaction products formed during milling, whose amount increases as mechanical treatment 120 TiC TiB2

L (nm)

100

80

60

40 0

4

8 12 16 20 Milling time, tM (h)

24

28

12.12 Effect of milling time on the crystallite sizes of TiC and TiB2 in the spark plasma sintered samples (I = 1100 A, P = 20 MPa, tSPS = 4 min).

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proceeds, also contributes to the formation of materials by SPS with grain sizes in the nanoscale range. This is in agreement with the behaviour noticed previously during the synthesis of dense nanometric MoSi2.

12.3.3 NbAl3 system The influence of milling time on the XRD patterns of the (Nb + 3Al) co-milled powder reactants is shown in Fig. 12.13 for the case where CR = 1. This result indicates that both Nb and Al peaks broaden with increasing milling time. In addition, within the detection limit of this analysis, intermetallic phases are not formed at milling times equal to or less than 20 hours, while the first evidence of formation of NbAl3 was observed at tM = 35 h. The synthesis and consolidation of the NbAl3 intermetallic phase was then performed using the SPS apparatus under a temperature controlled

Nb Al NbAl3

Intensity (a u)

35 h

20 h

10 h

0h 20

30

40

50

60

70

80

2θ (degree)

12.13 XRD patterns of (Nb + 3Al) reactants as a function of milling time (CR = 1).

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Nb Al NbAl3

Intensity (a u)

35 h

20 h

10 h

0h 20

30

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12.14 Influence of milling time on the composition of Nb–Al samples obtained by SPS (TD = 600°C, P = 20 MPa).

mode and starting from unmilled and co-milled powders. The effect of milling time on the composition of SPS products was investigated first, when pressure and dwell temperature, TD, were maintained constant at 20 MPa and 600°C, respectively. It should be noted that the relatively low temperature and pressure levels adopted aimed to avoid possible spilling of molten aluminum (660°C melting point) from the die in the presence of a mechanical load. The XRD results of the samples obtained are reported in Fig. 12.14, which shows that, as a consequence of the milling treatment, the reactivity of the solid reactants during the SPS process is remarkably enhanced. In particular, the degree of conversion to the desired intermetallic phase reached during reactive sintering monotonically

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increases as the milling time is augmented. Moreover, no evidence of reaction was observed when starting from unmilled reactants. Nevertheless, the best result (tM = 35 hours) is still characterized by incomplete conversion and the product obtained was rather porous (about 70% of the theoretical density). Thus, with the aim of increasing both product conversion and densification, the experimental conditions (temperature and pressure) have been modified during the course of the SPS process. Specifically, while the starting powders still consist of Nb + 3Al reactants co-milled for 35 hours and CR = 1, the temperature and pressure programmes indicated in Fig. 12.15(a) have been chosen. Correspondingly, as shown in Fig. 12.15(b), a final product consisting of a single NbAl3 phase was obtained. It is worth noting that, if the SPS process was performed by adopting only a two-step thermal cycle while maintaining the mechanical pressure equal to 20 MPa, the corresponding end sample was fully converted but remains still porous (about 88% of the theoretical density). Regarding the formation of NbAl3 during SPS, the corresponding sample displacement did not display any sharp transition as time proceeded, thus indicating that a gradual synthesis mechanism takes place. The consolidation level was significantly improved by properly increasing the mechanical load as indicated in Fig. 12.15(a). Indeed, a near fully dense (97.5 ± 2.5% of the theoretical density) NbAl3 product with a crystallite size below 100 nm was obtained. It should be noted that the increase in temperature and mechanical load is possible only when the presence of Al is significantly reduced, so that its potential spill from the die is avoided.

(b) 100

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(a)

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12.15 The two-step temperature and mechanical load temporal profiles (a) adapted during the SPS of NbAl3 and the corresponding XRD pattern of the obtained product (b) from co-milled (tM = 35 h) Nb + 3Al reactants.

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Conclusions and future trends

The effect of the mechanochemical activation of reactants processed by an electric current activated sintering process, namely the SPS, is examined in this chapter. In particular, after reviewing the most relevant systems investigated following this route so far, the syntheses and simultaneous consolidation of MoSi2, TiC–TiB2 and NbAl3 are discussed in detail. The general conclusion is that the milling treatment strongly affects SPS process dynamics and the kinetics of the synthesis process, as well as the product relative density and crystallite sizes, which are reduced to nanometric levels. These features may be associated with physicochemical modifications of solid reactants taking place during BM, which in turn affect the nature and rate of chemical and structural transformations involved in the SPS process. Specifically, during milling, the interfacial area between the reactants increases and the powder particle size decreases, thus enhancing mass transport by diffusion and, consequently, synthesis and sintering processes as well as product microstructure refinement. Furthermore, it was experimentally evidenced that the presence of nanostructured reaction products formed during milling also contributes to the formation of products with grain sizes in the nanoscale range. However, these benefits can partially or totally vanish when the final aim is the fabrication of nanostructured materials in bulk form. In fact, as observed in the Introduction, conventional consolidation methods require high sintering temperatures and prolonged processing times, which significantly favour grain growth. In contrast, grain growth can be considerably reduced since solid–solid reactions and consolidation are strongly accelerated using the SPS process, as powders to be sintered and/or the die containing them are directly and rapidly crossed by an electric current. These features are confirmed by the results obtained when investigating the MoSi2, TiC–TiB2 and NbAl3 systems described in this chapter. Joule heating and mass transport enhancement through electromigration (Huntington, 1975) are the most likely phenomena which explain the observed acceleration of reactive sintering that occurs during SPS, although other controversial mechanisms possibly involved in the process, like the plasma generation within the voids surrounding powder particles (cf. Orrù et al., 2009), are also postulated. Although the combination of BM and SPS techniques has been already applied to hundreds of systems (cf. Table 12.1), these methods can be additionally coupled for the fabrication of other nanostructured materials of practical importance. Moreover, in contrast to the cases investigated in detail in this chapter, milling treatment can be also further exploited to improve the refining process of nanostructured powders to be densified by SPS. Following these ideas, potential benefits are expected by taking

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advantage of the cryomilling technique, where the mechanical treatment is performed in presence of liquid medium, generally nitrogen (Lavernia et al., 2008). In fact, the extremely low temperature encountered during milling (about 77 K for the case of liquid N2) restrains recovery and recrystallization which leads to more rapid grain refinement thus giving finer grain structures. In summary, the combination of this technology with SPS may represent a powerful strategy for the development of bulk nanostructured products in the near future.

12.5

Acknowledgements

The authors acknowlege the collaboration with Prof. Zuhair A. Munir, University of California, Davis (USA). IM (Innovative Materials) S.r.l., Italy, is also gratefully acknowledged for granting the use of the SPS-515S apparatus.

12.6

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yamasaki t, zheng y j, ogino y, terasawa m, mitamura t and fukami t (2003), ‘Formation of metal-TiN/TiC nanocomposite powders by mechanical alloying and their consolidation’, Mat Sci Eng A, 350 (1–2), 168–72. yan x, song x, lu n, li e d and zhang j (2008), ‘A novel route for preparing binary Sm-Co bulk amorphous alloys’, Mater Lett, 62 (17–18), 2862–4. yang c, jin s z, liang b y and jia s (2009), ‘Low-temperature synthesis of high-purity Ti3AlC2 by MA-SPS technique’, J Eur Ceram Soc, 29 (1), 181–5. ye l l, liu z g, raviprasad k, quan m x, umemoto m and hu z q (1998), ‘Consolidation of MA amorphous NiTi powders by spark plasma sintering’, Mat Sci Eng A, 241, 290–3. ye j, ajdelszajn l and scoenhung j m (2006), ‘Bulk nanocrystalline aluminum 5083 alloy fabricated by a novel technique: crysomilling and Spark Plasma Sintering’, Metall Mater Trans A, 37, 2569–79. yoo s h, wang h j, oh s t, kang s g and choa y h (2006), ‘Synthesis and densification of CNTs/Fe/Al2O3 nanocomposite powders by chemical vapor deposition’, Key Eng Mater, 317–18, 665–8. zhan g d, kuntz j, wan j, garay j and mukherjee k (2003), ‘Spark-plasma-sintered BaTiO3/Al2O3 nanocomposites’, Mat Sci Eng A, 356, 443–6. zhang j x, lu q m, liu k g, zhang l and zhou m l (2004), ‘Synthesis and thermoelectric properties of CoSb3 compounds by spark plasma sintering’, Mater Lett, 58, 1981–4. zhang h w, gopalan r, mukai t and hono k (2005a), ‘Fabrication of bulk nanocrystalline Fe–C alloy by spark plasma sintering of mechanically milled powder’, Scripta Mater, 53, 863–8. zhang y f, wang l j, jiang w and chen l d (2005b), ‘Microstructure and properties of Al2O3–TiC composites fabricated by combination of high-energy ball milling and spark plasma sintering (SPS)’, J Inorg Mater, 20, 1445–9. zhao l d, zhang b p, li j f, zhou m, liu w s and liu j (2008), ‘Thermoelectric and mechanical properties of nano-SiC-dispersed Bi2Te3 fabricated by mechanical alloying and spark plasma sintering’, J Alloy Compd, 455, 259–64. zheng y, yamasaki t, mitamura t, terasawa m and fukami t (2003), ‘Formation of ultra-fine grained TiC-dispersed SUS316L by ball-milling and their consolidation by hot isostatic pressing’, J Jpn Inst Met, 67, 74–8. zhou y and hirao h (2005), ‘Preparation and tribological properties of SiC/rice bran carbon composite ceramics’, J Mater Res, 20, 3439–48. zhu l h, huang q w and zhao h f (2003), ‘Preparation of nanocrystalline WC–10Co0.8VC by spark plasma sintering’, J Mater Sci Lett, 22, 1631–3.

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13 Synthesis of titanium dioxide-based, visible-light induced photocatalysts by mechanochemical doping S H U Y I N, Q I W U Z H A N G, F U M I O S A I T O and T S U G I O S AT O, Tohoku University, Japan

Abstract: An environmentally friendly soft chemical process for the synthesis of visible-light induced photocatalysts is introduced in this chapter. Titania-based anion doped photocatalysts can be effectively prepared by the ‘low-temperature mechanochemical doping treatment’. Compared with conventional preparation processes, mechanochemically synthesized titania-based photocatalysts show excellent ultraviolet and visible-light induced photocatalytic activities, because of their fine particles, high crystallinity and effectivity of doping treatment. Key words: mechanochemical doping, visible-light induced photocatalyst, anion-doped titanium dioxide, environmental purification.

13.1

Introduction

Modern society has been sustained by the remarkable development of advanced materials and by a huge consumption of energy and resources, but the waste from these materials, chemicals, energy and heat will increase markedly in the future causing environmental problems for the Earth. Since the Earth is an almost closed system, with limited energy resources, energy and heat wasted in preparing materials should be minimized. It is necessary to look for materials which are less hazardous to humans and find more environmentally friendly processing methods to fabricate such materials. Many researchers in the world are focusing their attention on the photocatalytic reaction system, which utilizes clean and inexhaustible solar energy. The photocatalytic system is a chemical process with a low environmental impact and is expected to be widely used for the purification of air and polluted water. Titanium dioxide and its derivatives, such as titanate, are well-known photocatalytic materials and are widely used in the purification of air and water, solar-powered systems and so on. It is well known that photocatalytic 304

Photocatalyst synthesis by mechanochemical doping

305

activity is closely related to the crystallinity, phase composition, specific surface area, microstructure and morphology of the prepared powders. Visible-light induced photocatalysts are expected to be developed in order to utilize solar energy effectively. Usually, highly active photocatalytic materials are synthesized by expensive processes, such as a hightemperature solid state reaction, ion injection or a hydrothermal process, for which a high temperature or high performance load is necessary. In the following sections, a general introduction to TiO2 photocatalysts, their structure, application and characterization is given and recent developments in visible-light reactive photocatalysts are reviewed. Environmentally friendly mechanochemical doping processes for the synthesis of photocatalyst materials are introduced. Compared with conventional preparation processes, materials prepared by the mechanochemical doping process, which has a low environmental impact, show excellent photocatalytic activity because of their fine particles, high crystallinity and especially their high doping efficiency. Both ultraviolet (UV) and visible-light irradiation activity can be achieved by optimizing the reaction conditions.

13.2

Titanium dioxide and its photocatalytic applications

Titania is the world’s most important white pigment and has wide applications in cosmetics, catalysts and catalyst carriers because of its excellent physical and chemical properties.1 Fujishima and Honda2 discovered a process for the photocatalytic splitting of water on TiO2 electrodes. Since then, an understanding of the fundamental processes and of ways to enhance the photocatalytic efficiency of titania have come from extensive research performed by chemists, physicists and chemical engineers. In recent years, these studies have been related to energy conversion such as the electrochemical photolysis of water and use of high-efficiency solar cell systems.3–7 Titania has also received much attention with respect to the degradation of various environmental pollutants in water and the atmosphere. A titania photocatalyst can also be used in many other applications such as self-cleaning, sterilization and deodorization.8 Although various inorganic substances could be used as photocatalysts, titania is the most effective because of its great efficiency and excellent chemical stability.9

13.2.1 Structure of titanium dioxide There are seven known polymorphs of TiO2, six of which have a distinct structure (Table 13.1). Three of these polymorphs, rutile, anatase and brook-

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Table 13.1 TiO2 polymorphs with a distinct structure. Adapted from Banfield et al.10 Structure

Space group

Density (g cm−3)

Unit cell data (nm)

Anatase Rutile Brookite

I41 /amd P42 /mnm Pbca

3.79 4.13 3.99

TiO2 (B), B-phase

C2/m

3.64

TiO2 II, highpressure phase TiO2 (H), Hollandite

Pbcn

4.33

I4m

3.46

a = 0.379, c = 0.951 a = 0.459, c = 0.296 a = 0.917, b = 0.546, c = 0.514 a = 1.216, b = 0.374, c = 0.651, β = 107.29 a = 0.452, b = 0.550, c = 0.494 a = 1.018, c = 0.297

ite, occur naturally.10 Anatase is known to consist of TiO6 octahedra, which share corners and four edges with neighbouring octahedra; rutile TiO6 octahedra share corners and only two edges with neighbouring octahedra, while brookite TiO6 octahedra share corners and three edges with neighbouring octahedra. Anatase and brookite are thermodynamically metastable phases, while rutile is a thermodynamically stable phase. These differences in lattice structure lead to different mass densities and electronic band structures for the three forms of TiO2.11 Recently, monoclinic TiO2(B), which is composed of edge- and corner-linked pairs of edgesharing octahedra structures, has been discovered.10 Its structure is less compact than that of the anatase and rutile forms of titanium dioxide.10,12 Rutile and anatase are commonly used in photocatalysis, with anatase showing higher photocatalytic activity.13 However, almost no research has been carried out on the brookite and monoclinic phases. The photocatalytic enhancement of anatase has been ascribed to differences in the Fermi level, surface area and the extent of surface hydroxylation of the solid.14–16 In addition, rutile usually shows harder agglomeration and a larger particle size than anatase, since rutile is normally prepared by high-temperature calcination of anatase.

13.2.2 Photocatalytic reaction mechanism It is accepted that TiO2 particles absorb UV light with energy greater than the band gap of c. 3.2 eV to generate electron–hole pairs (Equation [13.1]). The electrons are photoinduced to the conduction band (e−BC), and the holes in the valence band (h+VB) are subsequently trapped by H2O to yield H+ and ⋅OH radicals (Equation [13.2]). hv − + TiO2 ←⎯ → TiO2 ( eBC + h VB )

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Potential (V) vs normal hydrogen electrode

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307

–2 •O2– –1 O2 0

H2O

(a)

(b) ECB

H2/H2O

e–

H2

1 O2/H2O

(hν>Eg)

Eg

2 EVB 3

•OH h+

H2 O

13.1 Schematic illustration of the energy correlation and the redox mechanism on a titania catalyst surface (a) with and (b) without oxygen. Adapted from Yin and Sato.20

Figure 13.1 illustrates the energy diagram and photoinduced charge transfer process in a TiO2 particle. Although the band gap of titania is larger than the potential energy of the electrochemical decomposition of water (1.23 eV), pure H2O cannot be photodecomposed on a clean titania surface because of the rapidly occurring reverse reaction of H2 and O2. In the absence of oxygen and in the presence of sacrificial species, such as methanol, the holes generated by the light oxidize methanol to HCHO, while electrons in the conduction band of the crystal lattice simultaneously reduce water molecules and/or protons in solution to form gaseous H2, as shown by Equations [13.2]–[13.5].17–21 These reactions proceed competitively with the recombination of the photoinduced electrons and holes. h+VB + H2O → ⋅OH + H+

[13.2]

CH3OH + ⋅OH → ⋅CH2OH + H2O

[13.3]

⋅CH2OH → HCHO + H+ + e−CB

[13.4]

2H2O + 2e−CB → H2↑ + 2OH−

[13.5]

On the other hand, in the presence of oxygen, the electrons in the conduction band can be rapidly trapped by the molecular oxygen and adsorbed on the TiO2 particle surface, to form ⋅O−2 which can then generate highly

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active ⋅OOH radicals18 (Equations [13.6] and [13.7]). An environmental pollutant, such as phenol, in the solution reacts with these reactive oxygen radicals, and/or molecular oxygen, to produce intermediate or complete decomposition products, CO2 and H2O21,22 (Equation [13.8]). In the case of the gas phase reaction, an environmental pollutant such as nitrogen monoxide, NO, reacts with the reactive oxygen radicals to produce nitrous acid (HNO2)/nitric acid(HNO3) or nitrogen and oxygen directly23 (Equations [13.9] and [13.10]). The acid produced is thought to be neutralized by basic ash that exists in the air to achieve environmental purification. e−CB + O2 → ⋅O−2 − 2

+

[13.6]

⋅O + H → ⋅OOH

[13.7]

C6H5OH +{⋅OH, ⋅OOH, and/or O2} → intermediates → CO2 + H2O

[13.8]

NO +{⋅OH, ⋅OOH, and/or O2} → HNO2 and/or HNO3 (80%)

[13.9]

→ N2 + O2 (20%)

[13.10]

13.2.3 Highly active visible-light induced titanium dioxide and its applications Utilization of a photocatalytic reaction system might effectively achieve purification of the atmosphere. Figure 13.2 illustrates an environment purification system based on a photocatalytic reaction. Building materials coated with a photocatalyst might be used for noise suppression board fences along highways, and also as deNOx functional materials under solar light irradiation. The photocatalytic production of nitrous acid (HNO2)/ nitric acid (HNO3) might be neutralized by basic ash existing in the air or washed out by rainwater. After collection and neutralization treatment, the rainwater could be released into rivers. One problem with using anatase as a photocatalyst is the large band gap energy, that is anatase only shows photocatalytic activity under UV light irradiation with a wavelength <387 nm, which corresponds to its band gap value of 3.2 eV. It is known that solar energy contains only about 5% UV light (wavelength <400 nm) and most of the rest is visible light (wavelength 400–700 nm). In order to utilize solar energy efficiently, it is necessary to develop a photocatalyst that is reactive to visible light. Asahi, Morikawa and co-workers (2001)24,25 reported that nitrogen-doped titanium oxide with high visible-light photocatalytic activity for the decomposition of methylene blue and acetaldehyde could be prepared by sputtering the TiO2 target in an N2 (40%)/Ar gas mixture followed by annealing in N2 gas at 550°C for

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Building materials hν TiO2-xNy photocatalyst slurry

TiO2-xNy photocatalyst coated noise suppression board fence along the roadside

NOx NO3–

Spray Road NOx

Neutralization

Washed by rainwater

NOx

Drain ditch

River

Drain ditch Settling tanks

13.2 Illustration of an environment purification system by photocatalytic reaction.

4 hours, or by treating anatase TiO2 powder in an NH3(67%)/Ar atmosphere at 600°C for 3 hours. It was observed that nitrogen-doped titania was yellow. Optical absorption spectra show that the TiO2−xNx noticeably absorbs light at less than 500 nm. It was also predicted that other anions, such as C, S, and F, would have a similar effect to nitrogen. Asahi et al.24 theoretically calculated the band structure of anion-doped titania by first principles calculations and pointed out that N-doping led to visible light activity and a narrowing of the band gap by mixing the N 2p and O 2p state (see Fig. 13.3). Until now, there have been various methods for the preparation of photocatalysts that respond to visible light. However, most of them are hightemperature processes or they use expensive preparation instruments.26–36 It is well known that photocatalytic activity is strongly related to physical properties such as crystal phase, particle size, specific surface area, crystallinity and morphology. It is thought that high-temperature treatment results in the formation of rutile, crystal growth, reduction of specific surface area, decrease in the amount of anion doping and consequently a reduction in the visible-light responsive photocatalytic activity. In order to avoid these disadvantages, it is necessary to develop low-temperature soft chemical processes. In addition, it is hoped to avoid using irritant gases such as NH3, H2S, expensive precursors such as TiN, TiS2 and to reduce the formation of

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High-energy ball milling TiO2 F-doped N-doped

DOS (a.u.)

C-doped S-doped P-doped Ni-doped Ni+s-doped

–10

–5

0 E (eV)

5

10

13.3 Total densities of states (DOS) of doped TiO2. From Asahi et al. (2001), Science, 293, 269–71.24 Reprinted with permission from AAAS.

poisonous by-products such as SOx and NOx during the photocatalyst preparation process. This review introduces the preparation of photocatalysts that are highly responsive to visible light by a low-temperature mechanochemical process.

13.3

Synthesis of anion-doped, titanium-dioxide, visible-light induced photocatalysts by mechanochemical doping

Mechanochemical technology has recently become a promising method for material synthesis, surface modification, waste processing, recycling, and so on. In many cases, mechanical stressing induces the formation of a fresh oxygen-rich surface, resulting in electron transfer from the O2− ion on the oxide surface to the other organic substance and, as a result, leads to the destruction of the weak bonding in the organic substance and the formation of a new bonding between the oxide and the non-metallic element.37,38 We were the first to discover that mechanochemical technology might be an effective method for doping non-metallic elements such as nitrogen and fluorine into photocatalysts at around room temperature.39–42 Since then, many researchers have concentrated on this field of research. It has been

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found that mechanochemical treatment conditions have a great effect on phase transformation, adsorption and other physicochemical properties.

13.3.1 Effect of mechanochemical treatment on the structure, sorption and physicochemical properties of titanium dioxide The phase transformation of titania by ball milling has been reported by some researchers. Begin-Colin reported that anatase titania can be transformed by high-energy ball milling into rutile via the high-pressure modification of TiO2-II,43 while Pan et al.44 reported that high-energy ball milling induced the transformation from anatase to srilankite and rutile at room temperature and ambient pressure. They also pointed out that milling atmospheres with different oxygen partial pressures had an influence on the transformation kinetics of anatase. The transformation rate, in turn, increased with a decrease in oxygen partial pressure of the milling atmosphere, owing to the reducing concentration of oxygen vacancies in the titania lattice of particles. During ball milling, some of the Ti-O bonds on the titania surface layer may be broken by mechanical activation and oxygen is released from the titania lattice.44 In the synthesis of nitrogen-doped titania, some nitrogen sources were used during the mechanochemical process. It was found that not only the milling time, but also the rotation rate and the quantity of nitrogen source additives had a great effect on the phase transformation of titania.42,45 Figure 13.4 shows the rutile fraction in the ball-milled products prepared from a mixture of a commercial titania (Deggusa P25, which consists of about 77% anatase and 23% rutile phases) and a nitrogen source, hexamethylenetetramine (HMT). It is clear that the fraction increased with the rotation rate and milling time and decreased with the amount of nitrogen source. The HMT powder that was added both eased mechanical stress and led to a slight reduction in phase transformation. The temperature during the mechanochemical treatment was less than 50°C even after milling for a long time. Generally, the transformation temperature of anatase to rutile is much higher at 700°C. These results strongly suggested that high mechanical energy accelerated the phase transformation of anatase to thermodynamically stable rutile, although some intermediate phase might also be formed. Figure 13.5 shows the particle size distribution of P25 titania powders before and after planetary ball milling at 70 rpm for 60 min. The original P25 powder showed well-defined particles of diameter 20–40 nm. Some small particles were formed on the surface after milling and agglomeration of the particles continued. It is probable that the fresh surfaces exposed by the milling lowers their high surface energy by agglomeration. Particle size

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High-energy ball milling 100.0 R/(A+R) (%)

80.0

Zirconia vessel / balls

60.0 40.0 Agate vessel / balls (a)

20.0 0.0

0

60 90 120 Milling time (min)

30

150

180

100.0 R/(A+R) (%)

80.0 60.0 40.0

Zirconia vessel / balls

20.0 0.0

(b) 0

100

200

300 400 500 600 Milling speed (rpm)

700

800

R/(A+R) (%)

100.0 80.0 60.0

Zirconia vessel / balls

40.0 20.0 0.0

(c) 0

5

10 HMT additive (%)

15

20

13.4 Effect of (a) milling time with 10 wt% HMT at 700 rpm, (b) rotation speed with 10 wt% HMT for 180 min and (c) amount of HMT at 700 rpm for 120 min on the mole fraction of rutile phase in titania powders prepared by planetary ball milling of P25 titania with HMT using agate or zirconia vessels. A, anatasre; R, rutile; R/(A+R), molar ratio of rutile phase. Adapted from Yin.42

distribution analysis indicated that P25 initially had a medium diameter (d50) of around 0.023 μm. After milling without any additive, the peak of d50 shifted to around 0.074 μm, indicating that particle agglomeration was increased by the mechanochemical treatment. Similarly, in the case of milling with a 10 wt% nitrogen reagent mixture, the d50 values changed to about 0.056–0.075 μm. Uzunova-Bujnova et al.46 reported that the sorption

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Particle size (10–6 m)

0

20 30 Particle amount (%)

40

0.023 μm

P25 as received

10

0.074 μm

P25 without additive

d50

50

>5.116 4.154 3.373 2.739 2.224 1.806 1.466 1.190 0.967 0.785 0.637 0.517 0.420 0.341 0.277 0.225 0.183 0.148 0.120 0.098 0.079 0.064 0.052 0.042 0.035 0.028 0.023 0.018 <0.015 0

40

0.056 μm

10 wt % (NH4)2CO3

20 30 Particle amount (%)

0.077 μm

10 wt % urea

10

0.075 μm

10 wt % HMT

d50

50

13.5 Particle size distribution and the medium diameter, d50 (particle size in case of 50% oversize), of titania powders prepared by planetary ball milling at 700 rpm for 60 min with or without 10 wt% nitrogen reagents. Adapted from Yin et al.45

>5.116 4.154 3.373 2.739 2.224 1.806 1.466 1.190 0.967 0.785 0.637 0.517 0.420 0.341 0.277 0.225 0.183 0.148 0.120 0.098 0.079 0.064 0.052 0.042 0.035 0.028 0.023 0.018 <0.015 Particle size (10–6 m)

314

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ability of the titania significantly increased as a result of mechanoactivation, due not only to the increased dispersity, but also to the changes in the surface properties, phase composition and crystal defect concentration (see Fig. 13.6).

13.3.2 Nitrogen-doped titanium dioxide and its photocatalytic activity Nitrogen-doped yellowish rutile titania with excellent visible-light photocatalytic deNOx activity was first successfully synthesized by a mechanochemical doping process in 2002–03.39 It was a very simple process: white titania powder was mixed with 5 wt% urea or 10 wt% HMT before being placed in a reaction pot. Seven zirconia balls 15 mm in diameter and 4 g of mixed powder were then placed in zirconia vessels with 45 cm3 inner volume. Grinding was performed using planetary ball milling apparatus at a rotation speed of 100–700 rpm for the desired time. Diffuse reflectance spectroscopy (DRS) is a technique that collects and analyses scattered light energy. It is used to measure fine particles and powders, as well as rough surfaces. Sampling is fast and easy because little or no sample preparation is required. Diffuse reflectance spectra are used

1.0

6 Ct/Cd

0.8 5

In (Cd /Ct)

4

0.6 0.4 0.2 0.0

3

0

50 100 150 200 250 Time (min)

2 1 0 0

50

100 150 Time (min)

200

250

13.6 Sorption and degradation of malachite green oxalate under the action of TiO2 (䊐) and TiO2 mechanoactivated in an agate vessel for 1 hour as a suspension in methanol (䊊). In the insert, the experimental data (symbols) and the theoretical fit of the data (lines) are shown. Cd, concentration at the end of the ‘dark period’; Ct, concentration in the moment t; t, illumination time. From Uzunova-Bujnova et al. (2008) Mater. Chem. Phys., 110, 291–8.46 Reprinted with permission from Elsevier.

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in the characterization of the absorbance and band structure of inorganic semiconductors. By using diffuse reflectance spectra measurements of nitrogen-doped titania, the band gap of the original titania structure and nitrogen doping status could be determined, respectively. Figure 13.7 shows the diffuse reflectance spectra of P25 titania and samples prepared by planetary ball milling a P25–10 wt% HMT mixture for various milling times. The P25 titania showed an absorption edge at 408 nm corresponding to a band gap of 3.04 eV. The powders prepared by planetary ball milling the P25 titania-HMT mixture showed two absorption edges at 400–408 nm (3.04–3.10 eV) and 530–550 nm (2.25–2.34 eV). It is thought that the first and second edges are related to the band structure of the original titania and the newly formed N 2p band, which is located above the O 2p valence band,24,25 respectively. It is obvious that the band gap could be greatly narrowed by mechanochemically doping nitrogen in titania. The nitrogen-doped amount increased with milling time and rotation speed, but slightly decreased with nitrogen source reagent content. The nitrogen-doped amount using agate balls and a vessel was less than 0.1 wt%, whereas the amounts attained using zirconia balls and a vessel were around 0.20 wt%. Figure 13.8 shows the relationship between the wavelength of the light source and the photocatalytic ability of the powders prepared by the mechanochemical process. Under irradiation of visible light at wavelength >510 nm, nearly 27% of the nitrogen monoxide could be continuously removed by nitrogen-doped titania prepared by planetary milling the P25

100 (a)

Reflectance (%)

80

(f) 60 (b) 40

(d)

(c)

20 0 200

(e)

300

400 500 Wavelength (nm)

600

700

13.7 Diffuse reflectance spectra of P25 titania (a) and the samples prepared by planetary ball milling of a P25 titania–10 wt% HMT mixture using zirconia vessel and balls at 700 rpm for (b) 15 min, (c) 60 min, (d) 120 min, (e) 150 min and (f) 180 min. All samples were calcined in air at 400°C for 60 min. Adapted from Yin et al.42

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> 400 nm

Residual NO concentration (%)

100.0

> 290 nm

Light off

(b) (c)

80.0

(d)

(a)

(e)

60.0 40.0 Vis-light

UV-light

20.0 0.0

0

2

4

6

8

0

2

4

6

8 0 2 4 Time (min)

6

8

0

2

4

6

8 10

13.8 Relationship between light wavelength and the photocatalytic ability for oxidation of nitrogen monoxide in the prepared samples. (a) 䉬, P25 titania; (b) 䊉, planetary milled P25–10 wt% urea mixture at 700 rpm for 1 hour; (c) 䊊, calcined (b) in air at 400°C for 1 hour; (d) 䉱, planetary milled P25–5 wt% HMT mixture at 700 rpm for 1 hour; (e) 䉭, calcined (d) in air at 400°C for 1 hour. Adapted from Yin et al.39

titania–5% HMT mixture followed by calcination in air at 400°C. The powder prepared by planetary milling the P25 titania–10% urea mixture showed lower activity than that with 5% HMT. On the other hand, the photocatalytic activity under visible-light irradiation of the as-prepared powders without calcination was negligible. This is due to the depression of NO adsorption by the remaining reaction products such as CO2, NH3, carbon, organic molecules, and so on. In addition, as expected, no photocatalytic activity was observed for P25 because of its large band gap energy of 3.1 eV. In the case of irradiation by light of wavelength >400 nm, similar results were observed. The activities of the powders after calcination ((c) and (e)) were at almost the same level as that of P25 TiO2 but higher than those before calcination ((b) and (d)). When the light was turned off, the nitrogen concentration at the outlet of the reactor returned to the initial concentration of 1 ppm within 10 min, indicating that light energy is essential for the oxidation of nitrogen monoxide. It is thought that the chemical reaction between titania and HMT or urea occurs together with the phase transformation of titania. Usually, HMT/ urea hydrolyses to form ammonia and formaldehyde/CO2; however, it was found that many carbon/related compounds, such as biuret and cyanuric acid, were generated by the mechanochemical reaction of HMT or urea, as shown by Equations [13.12] and [13.13], owing to the lack of H2O.

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Photocatalyst synthesis by mechanochemical doping C6H12N4 → 6C + 4NH3

(mechanochemical reaction)

317 [13.11]

2(NH2)2CO → NH3 + NH2CONHCONH2 (biuret) (mechanochemical reaction)

[13.12]

3(NH2)2CO → 3NH3 + C3H3N3O3 (cyanuric acid) (mechanochemical reaction)

[13.13]

This is the reason why the as-prepared powders showed low photocatalytic activity. The intermediate products or residual nitrogen sources depressed the adsorption of irradiated light on the surface of the photocatalyst. As shown in Fig. 13.8, the photocatalytic activity could be promoted by calcination at 400°C. Recently, it was also found that nitrogen-doped titania loaded with metallic ions (such as Fe or Pt) on the surface led to a great improvement in the photocatalytic activity and quantum yield. It was confirmed that NOx could be continuously removed successfully, not only under UV light irradiation but also under solar light and even some long-wavelength monochrome LED light (red LED: λ > 627 nm; green LED: λ > 530 nm; blue LED: λ > 445 nm, with light intensity of 2 mW).47

13.3.3 Carbon-doped titanium dioxide and its photocatalytic activity In addition to nitrogen-doped titania, carbon-doped titania has also attracted the attention of researchers. Asahi et al.24 predicted that other anions, such as C, S, and F, would possess a similar effect to nitrogen. In 2002, Khan et al.26 synthesized chemically modified TiO2−xCx by flame pyrolysis of a Ti metal sheet at 850°C in a natural gas flame with a controlled amount of oxygen. Irie et al.27 also prepared C-doped TiO2 by oxidative annealing of a TiC precursor in air/O2 at 600°C. Janus et al.48 pointed out that the photocatalytic ability of C-doped TiO2 could be improved by decreasing the recombination rate in a photogenerated electron–hole pair using electron scavenger carbon doped into TiO2. The nature of the carbon-induced modifications of the TiO2 electronic band structure remained unclear and the enhancement of photocatalytic activity was attributed either to the band gap narrowing or to the formation of a localized mid-gap state in the TiO2 band gap. In principle, any element could be doped into a discretionary matrix material by the aforementioned mechanochemical doping process. Greycolored carbon-doped titania was successfully and easily prepared by grinding TiO2 with ethanol and applying heat treatment.49 Anatase titania powder supplied by Wako Pure Chemical Industry (Osaka, Japan) was mixed with ethanol in 5% by mass and introduced into a planetary ball

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50 40 30

(c) (d)

(c) (d) (b)

(b)

(c) (a)

20 10

(a)

(a) Raw (b) As prepared (c) 200°C (d) 400°C

(b)

(d)

(a)

0 > 510 nm

> 400 nm

> 290 nm

13.9 DeNOx activity of the carbon doped titania treated by postmechanochemical heating; (a) raw-TiO2, (b) without heating (as prepared product), heated at (c) 200°C and (d) 400°C for 1 hour in air. Adapted from Kang et al.49

mill with a partially stabilized zirconia pot and grinding was then carried out at 700 rpm for various periods of time, followed by heat treatment in air to remove impurities from the ground products. Figure 13.9 illustrates the NOx gas decomposition activity of carbon-doped titania. This result also confirmed that surface impurities on the as-prepared samples prevented optical irradiation of the surface of the TiO2 in a redox reaction. The best photocatalytic activity of the C-doped TiO2 in the visible irradiation region was achieved by grinding for 2 hours followed by heating in air at 200°C. It was confirmed that grinding TiO2 with ethanol caused the formation of Ti–C and C–O bonds in the ground product. Both the Ti–C and C–O bonds remained after heating at 200°C. Further heating the product at 400°C caused the dissociation of Ti–C bonds in the product, whilst the C–O bonds still remained. This result induced a decrease in photocatalytic activity of the product when heated at 400°C. Using adamantine (ADM) instead of ethanol as the carbon source also led to carbon doping in the titania lattice. Details will be given in Section 13.3.5.

13.3.4 Sulfur-doped titanium dioxide and its photocatalytic activity In general, it is particularly difficult to dope sulfur, which has a large ionic size, into the lattice of titania. It has been reported that sulfur-doped titania could be synthesized by heating TiO2 at a high temperature under a gas

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flow of NH3 and H2S. However, calcination of TiO2 under H2S gas flow does not give satisfactory results.50 Heating TiS2 in air has been proposed for preparing sulfur-doped TiO2; however, use of the expensive TiS2 precursor could not be avoided.28,29 Ohno et al.30 also prepared yellowish S-doped TiO2 by mixing titanium isopropoxide with thiourea and ethanol, followed by evaporation and calcination in air at 400–700°C. Zhang et al.51 provided the basic information for easy-to-perform sulfur doping of TiO2 by mechanochemical treatment, where an inexpensive sulfur sample could be used as the doping agent. Commercial titania powder supplied by Wako Pure Chemical Industry was mixed with 10 wt% sulfur and 2 g of the mixture were put in a zirconia pot (45 cm3 inner volume) with seven zirconia balls (15 mm in diameter) and subjected to grinding using a planetary ball mill at a rotation speed of 700 rpm in air for different periods of time. Because the sulfur element was easy to oxidize in high-temperature air, the ground samples were calcined at 673 K for 60 min in an argon flow. Figure 13.10 shows the S 2p X-ray photoelectron spectroscopy (XPS) spectra of the mixtures ground for 20 min and 120 min. The S 2p state has a broad peak because of the limited amounts of sulfur doped, as well as an overlap of the split sublevels, the 2p3/2 and 2p1/2 states, with separation of 1.2 eV by spin-orbit coupling.29,52,53 A distinct difference in the peak position was observed for spectra of two samples ground for different periods of time. The peak of the sample ground for 20 min was located at 165–163 eV, whereas that of the one

Absorbed SO2 120 min

20 min

172 170 168 166 164 162 160 158 Binding energy (eV)

13.10 XPS spectra of S 2p of the samples ground for 20 min and 120 min. Adapted from Zhang et al.51

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ground for 120 min was positioned at 163–161 eV. The peak positioned at 169 eV was related to the absorbed SO2 molecules on the TiO2 surface,16,17 while the peak at 162 eV resulted from bonding between sulfur and titanium atoms, and sulfur itself usually exhibits a peak position at 164 eV. The S–S bond remained as the main component in the product ground for a short time (20 min). Prolonged mechanochemical treatment is crucial to ensure the formation of S–Ti bonds in the titania lattice.

13.3.5 Co-doping effect on titanium dioxide photocatalysis Instead of hexamethylenetetramine, (NH4)2CO3 (denoted as NHC) could also be used as a nitrogen source to prepare nitrogen-doped titania; while adamantine (denoted as ADM), which has the same basic structure as that of HMT, could be used as a carbon source to prepare carbon-doped titania samples, respectively.54 Figure 13.11 shows the N 1s and C 1s XPS spectra of P25 and the TiO2−xAy (A = N, C) products prepared by mechanochemical doping. No peaks for N–Ti and C–Ti bonding were observed in the commercial P25 powders, indicating that no nitrogen or carbon was doped in the P25 powders. The binding energy around 396 eV, which relates to the existence of N–Ti, was

(a)

(b) C 1s

C

N 1s

C–Ti

With HMT

N–N N–Ti N–C N–O

Intensity (a.u.)

Intensity (a.u.)

With HMT With ADM With NHC

P25

298

With ADM With NHC

P25

294 290 286 282 Binding energy (eV)

278

408 406 404 402 400 398 396 394 392 Binding energy (eV)

13.11 (a) C 1s, and (b) N 1s XPS spectra of P25, the commercial powder, and the powders prepared by the planetary ball milling of P25 with 10 wt% NHC, ADM and HMT, followed by calcination in air at 400°C. The measurement was carried out after Ar+ ion sputtering for 3 min. Adapted from Yin et al.54

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confirmed in the TiO2−xNy samples prepared by using NHC and HMT, indicating that Ti–N bonds were actually formed in the lattice of the titania crystal during the mechanochemical treatment. It is known that the peak around 400 eV is related to the N–N, N–O or N–C bonds and the one around 402 eV is related to the N–H bonds.24,25 According to the integral intensity of the peaks around 396 and 400 eV, it was estimated that about 63% and 58% of the nitrogen atoms in the samples were actually incorporated into the TiO2 lattices (peak at 396 eV) by using HMT and (NH4)2CO3. It was also found that the binding energy around 282 eV, relating to the existence of C–Ti, was confirmed in the titania samples prepared using ADM and HMT. It is obvious that both nitrogen and carbon were actually co-doped in the lattice of the titania crystal during the mechanochemical treatment using HMT as a reaction reagent. Figure 13.12 shows the relationship between the wavelength of light irradiated and the photocatalytic ability for the oxidative destruction of nitrogen monoxide. The photocatalytic deNOx ability was calculated by Equation [13.14]: deNOx ability (%) = [(C − C0)/C0] × 100%

[13.14]

where C0 is the initial NOx concentration and C is the NOx concentration after photocatalytic reaction.

20 10 0

With ADM

Blank

30

Blank

With HMT

40

With HMT

(b) P25 P25-milled With NHC

(a)

P25 P25-milled With NHC With ADM

NO degradation ability (%)

50

>510 nm

>400 nm Wavelength (nm)

13.12 DeNOx ability of the titania powders prepared by a mechanochemical doping process using NHC, ADM and HMT as reaction reagents. Open dotted bars: as-prepared by mechanochemical doping; solid bars: calcined at 400°C for 1 hour. Diagonal lines: commercial titania P25. The photocatalytic activity is carried out under light irradiation of λ > 510 nm (a) and 400 nm (b). For comparison, the results of P25 before and after ball milling are also plotted. Adapted from Yin et al.54

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The photocatalytic activity of P25 decreased by c. 10–20% after milling without any additive at 700 rpm. This may have been due to the introduction of crystal defects or the phase transformation from anatase to rutile. It is obvious that nitrogen or carbon-doped titania prepared with NHC and ADM possessed relatively higher photocatalytic activity than non-doped titania under visible light >510 nm. The powder prepared with HMT showed photocatalytic activity about twice as high as that prepared with NHC and ADM, that is nearly 37% of nitrogen monoxide could be continuously removed by the nitrogen/carbon co-doped titania prepared with a 10 wt% HMT mixture. The activity was c. 6–7 times higher than that of P25, estimated by taking the blank result into consideration. This result indicated that nitrogen and carbon co-doping had multiple effects on the improvement of visible light photocatalytic activity. Similar results were observed for a wavelength >400 nm, although the differences between the samples were not so clear. The nitrogen-doped amount is controlled by optimizing the reaction conditions.

13.3.6 Titanium dioxide-based binary composite photocatalyst It has been found that the photocatalytic activity of classical TiO2 in the purification of polluted effluents containing volatile organic compounds can be greatly improved using coupling semiconductors and oxides, such as WO3/SiC–TiO2, TiO2–V2O5,55 SnO/TiO2,56 WO3–TiO2, ZnO–TiO2,57–59 SiO2/TiO2.60 The results indicate that this combination can result in an enhancement of photocatalytic activity. The efficiency of the binary composite system strongly depends on its composition. The increase and synergy in photo-oxidation activity are related to retarding the electron– hole recombination process by an electron-trapping effect and to the modification of surface acidity compared to that of pure TiO2. In titania-based photocatalytic research, it is well known that commercial photocatalyst Degussa P25 shows excellent photocatalytic activity because of its mixed structure of anatase and rutile at the nano level. The rutile phase acts as a hole acceptor and the anatase phase operates as the electron acceptor, resulting in a long lifespan for the photoinduced electron and hole and leading to effective separation of the photoinduced electron–hole.61,62 In principle, similar to Degussa P25 titania, the visible-light-induced photocatalytic activity of titania could be improved by the addition of nitrogendoped titiania. Figure 13.13 shows the mechanism of the electron–hole separation in TiO2−xNy/TiO2 composition during visible light-irradiated photocatalysis. In the case of the commercial Degussa P25 titania sample, rutile TiO2 has 3.0 eV band gap energy, so its photoactivity under visible light is less sensi-

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B– Light e–

323

e–

CB

CB

TiO2

Light e–

TiO2–xNx h+

h+ VB

A+

VB h+ A

13.13 Mechanism of electron-hole separation in TiO2−xNy /TiO2 composition during photocatalysis. Adapted from Kang et al.63

tive. Therefore, if a photocatalyst that is more sensitive to visible light replaces rutile, a photocatalyst that is more reactive to visible light may be obtained. Based on these results, it is reasonable to use nitrogen-doped TiO2 instead of the pure rutile phase. N-doped TiO2/anatase–TiO2 (TiO2−xNy/ TiO2) composites were prepared in ethanol with various weight ratios of TiO2−xNy/TiO2 (from 10 : 0 to 0 : 10) under relatively mild mechanochemical conditions. A 2.0 g sample was mixed with 30 ml ethanol and introduced into the mill pot with 40 g balls 5 mm in diameter, then the mechanochemical treatment was carried out at 200 rpm for 15 min. Finally, the residual solvent was removed by heating at around 80°C with magnetic stirring, followed by drying in an oven at 80°C for 12 hours.63,64 Figure 13.14 shows the acetaldehyde photodegradation ability of the TiO2−xNy/TiO2 composite samples under irradiation by black light at wavelength 352 nm. The CO2 concentration increased linearly with the light irradiation time. The commercial titania nanoparticles showed an obviously higher acetaldehyde photodestruction ability than those of TiO2−xNy under UV light irradiation. Furthermore, the acetaldehyde degradation activity of the composite samples was also higher than that expected by the additive property (dotted lines), that is the activity increased up to a mixture ratio of TiO2−xNy :TiO2 of 1 : 9–2 : 8. When the TiO2−xNy :TiO2 ratio = 2 : 8, the acetaldehyde decomposition activity of the composite was about 1.4 times higher than that of titania alone. In addition, the CO2 generation activity also increased up to TiO2−xNy :TiO2 = 1 : 9–2 : 8. These results also supported the idea that the photocatalytic activity of TiO2 was improved by mixing it with TiO2−xNy owing to the depression of electron-hole recombination by the heterogeneous transfer of photoinduced holes from TiO2 to TiO2−xNy.

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0.2

0.4

0.6

0.8

TiO2–xNy content (wt%)

1.0

CO2 generation rate (ppm min–1)

CH3CHO decomposition rate (ppm min–1)

(a)

6 5 4 3 2 1 0 0.0

0.2

0.4

0.6

0.8

1.0

TiO2–xNy content (wt%)

13.14 Acetaldehyde photo degradations ability of TiO2−xNy/TiO2 composites under the irradiation by black light (light wavelength 352 nm). The influence of TiO2−xNy/TiO2 mix ratio on the acetaldehyde degradation rate (a), and CO2 generation rate (b) is shown. Adapted from Liu et al.64

13.3.7 Fluorine doping and multiple-element co-doped photocatalysts Besides titania powders, SrTiO3 is also an important titania-based photocatalyst matrix and widely studied by many researchers because it has a similar band structure to that of titania. It was found that nitrogen could be doped into the lattice of SrTiO3 by the same mechanochemical method as for nitrogen-doped titania.65 It was also found that not only nitrogen, but also other elements such as fluorine could be doped by the mechanochemical treatment method. Saito66 synthesized lanthanum oxyfluoride (LaOF) by using polytetrafluoroethylene (PTFE, [-CF2-]n), as the fluoride source. Benziada-Taibi et al.67 reported that SrF2 and LiF could be used as fluorine co-doping sources in SrTiO3 ceramics containing fluorine by a solid-state reaction process. Wang et al.40,41 successfully synthesized fluorine-doped SrTiO3 using polytetrafluoroethylene, SrF2 and LiF as doping fluorine sources during the mechanochemical treatment process. From the XPS spectra of the prepared fluorine-doped SrTiO3 sample and that of commercially available SrF2 (Fig. 13.15), it is clear that the peak at 684.1 eV in SrF2-doped SrTiO3 can be attributed to the doping state of fluorine being shifted to 0.8 eV lower energy compared to that of the starting material SrF2 (684.9 eV), indicating that chemical bonds between the titanium and fluorine in the lattice of SrTiO3 were formed. The fluorine-doped SrTiO3 showed higher photocatalytic activity compared with that before doping, although the doped SrTiO3 samples had lower activity than that of the titania samples.

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(a) F-SrTiO3

(b) SrF2 675

680

685

690

695

Binding energy (eV)

13.15 XPS spectra of F1s of SrF2-doped SrTiO3 sample and commercially available SrF2 powder: (a) SrF2-doped SrTiO3, (b) pure SrF2. Adapted from Wang et al.40

As mentioned above, because of the narrowing of the band gap, nitrogendoped titania possessed excellent photocatalytic activity. However, replacing O2− with N3− would result in the formation of anion defects in the charge compensation and the anion defects would act as electron–hole recombination centres, which would limit improvement in the photocatalytic activity. It is expected that the charge compensation would be satisfied if O2− was replaced with N3− together with F−, without producing a large lattice strain. In a similar mechanochemical manner, SrTiO3 with nitrogen/fluorine codoping was successfully synthesized and showed higher photocatalytic activity than that synthesized by the heat treatment method.68

13.4

Future trends

Photocatalysis appears to be a promising technology which has various applications in environmental systems such as air and water purification. It can be demonstrated that an environmentally friendly, low-temperature mechanochemical process is effective for the synthesis of well-crystallized nanosize photocatalyst materials, which possess high UV and visible light activity. Besides nitrogen-doped titania, other anion-doped photocatalysts such as carbon, sulfur and fluorine were successfully prepared. Developments in the synthesis of highly active photocatalysts continue to be made. For example, a reformative complete low-temperature mechanochemical treatment method has been developed. In this process, a post-mechanochemical reaction washing process was applied instead of calcination treatment. Where soluble nitrogen reagents such as urea and/or (NH4)2CO3 were used, the effect was obvious. Compared with the samples treated with a post-mechanochemical reaction calcination process, the samples prepared by water-washing showed similar or higher deNOx activity (Fig. 13.16).

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DeNOx ability (%)

60 50 40 30

(c) (d)

20 10

(f) (g)

(b) (a)

(e)

(h)

0 510 Wavelength (nm)

13.16 Effect of washing process on the photocatalytic activity of TiO2−xNy prepared by mechanochemical doping. (a) P25, (b) as-prepared by using urea, (c) calcined at 400°C for 60 min, (d) washed with distilled water followed by drying at 80°C, (e) as-prepared by using (NH4)2CO3, (f) calcined (e) at 400°C for 60 min, (g) washed with distilled water followed by drying at 80°C, (h) blank (without catalyst).

These results also indicated that nitrogen doping was carried out during the mechanochemical treatment but not caused by the post-mechanochemical reaction calcination. In principle, cation doping, anion/cation co-doping in titania, in other oxides and in other composites could be achieved by mechanochemical treatment. Some successful examples such as nitrogen-doped ZnO, nitrogen and lanthanum co-doped SrTiO3 have already been reported.69–71 The mechanochemical doping process combined with another low-temperature process, such as O2 plasma treatment, is also an effective way to remove intermediate or unreacted organic compounds and to produce the photocatalyst at room temperature.72 In addition to photocatalytic applications, other advanced functional materials are expected to be synthesized by the mechanochemical doping process.

13.5

Acknowledgement

The authors would like to acknowledge AAAS and Elsevier for permission to use material from their published sources.

13.6

References

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60. liu z, quan x, fu h, li x and yang k (2004), ‘Effect of embedded-silica on microstructure and photocatalytic activity of titania prepared by ultrasoundassisted hydrolysis’, Appl. Catal. B: Environ., 52, 33–40. 61. hurum d, agrios a and gray k (2003), ‘Explaining the enhanced photocatalytic activity of degussa P25 mixed-phase TiO2 using EPR’, J. Phys. Chem. B, 107, 4545–9. 62. ohno t, tokieda k, higashida s and matsumura m (2003), ‘Synergism between rutile and anatase TiO2 particles in photocatalytic oxidation of naphthalene’, Appl. Catal. A, 244, 383–91. 63. kang i, zhang, q, yin s, sato t and saito f (2008), ‘Improvement in photocatalytic activity of TiO2 under visible irradiation through addition of N–TiO2’, Environ. Sci. Technol., 42, 3622–6. 64. liu b, yin s, li r, wang y and sato t (2007), ‘Preparation of visible light-activated titania photocatalyst by mechanochemical method’, J. Ceram. Soc. Jpn., 115, 692–6. 65. wang j, yin s, zhang q, saito f and sato t (2004), ‘Photo-oxidation properties of nitrogen doped SrTiO3 made by mechanical activation’, Appl. Catal. B: Environ., 52, 11–21. 66. lee j, zhang q and saito f (2001), ‘Effect of mechanical activation on the preparation of SrTiO3 and Sr2TiO4 ceramics from the solid state system SrCO3– TiO2’, J. Alloys Compd., 329, 230–8. 67. benziada-taibi l and kermoun h (1999), ‘Structural and nonlinear dielectric properties in fluoride containing SrTiO3 or BaTiO3 ceramics’, J. Fluorine Chem., 96, 25–9. 68. wang j, yin s, zhnag q, saito f and sato t (2004), ‘Photo-oxidation property of nitrogen and fluorine co-doped SrTiO3 made by mechanochemical method’, Transactions Mater. Res. Soc. Jpn., 29, 2693–6. 69. wang j, lu j, zhang q, yin s, sato t and saito f (2007), ‘Mechanochemical doping of a non-metal element into zinc oxide’, Chemistry for Sustainable Development, 15, 249–53. 70. lu j, zhang q, saito f, yin s and sato t (2004), ‘Mechanochemical doping of non-metallic elements on oxides and their characterization’, Sozaiken Iho, 60, 7–11. 71. wang j, yin s, zhnag q, saito f and sato t (2005), ‘Lanthanum and nitrogen co-doped SrTiO3 powders as visible light sensitive photocatalyst’, J. Euro. Ceram. Ceram. Soc., 25, 3207–12. 72. yin s, ihara k, komatsu m, zhang q, saito f, kyotani t and sato t (2006), ‘Low temperature synthesis of TiO2KxNy powders and films with visible light responsive photocatalytic activity’, Solid State Commun., 137, 132–7.

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14 Soft mechanochemical synthesis of materials for lithium-ion batteries: principles and applications N. V. K O S O VA, Institute of Solid State Chemistry and Mechanochemical SB RAS, Russia

Abstract: A soft mechanochemical synthesis approach was applied to prepare some fine particulate lithium-containing compounds with a spinel (LiMn2−xMxO4, Li4Ti5O12), layered (LiNi1−x−yCoxMnyO2, LiV3O8) and framework structure (LiFePO4, Li2FeSiO4, LiTi2(PO4)3). It consists of the fast formation of amorphous or low-crystalline precursors during mechanical activation of LiOH (or Li2CO3) with corresponding anhydrous oxides of d-metals, solid hydroxides, crystal hydrates and acidic salts in highly energetic planetary mills and transformation of these precursors into nanostructured final products under ‘soft’ heating conditions (lower temperatures and shorter time of treatment) after removal of water molecules, hydroxide groups and other volatile compounds. Most of as-prepared materials show good electrochemical properties as cathodes, anodes or solid electrolytes for lithium-ion batteries. Key words: soft mechanochemical synthesis, materials for lithium-ion batteries.

14.1

Introduction: principles of soft mechanochemical synthesis of solid inorganic compounds

Recently, nanostructured materials have attracted a great deal of attention. They have noticeable advantages over micrometre-sized materials. For instance, the advantages of nanostructured active electrode materials used in lithium-ion batteries can be summarized as follows: •

a larger electrode/electrolyte contact area leads to higher charge/ discharge rates; • short path length for both electronic and Li-ion transport permits operation of materials even with low electronic or low Li-ion conductivity; • significantly reduced specific current density owing to very large surface area effectively stabilizes an electrode and preserves its capacity at high charge/discharge rates; • stability to large volume expansion/contraction conserves the integrity of the electrode and leads to stable cycle performance. 331 © Woodhead Publishing Limited, 2010

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As a rule, manufacturing nanomaterials requires more complex synthesis process, generally based on solutions, including metallorganics, and is accompanied by large amounts of waste products. Therefore, the development of new facile synthesis methods for large-scale production of such materials is highlighted. The idea of performing the reactions directly between solids, excluding the dissolution stage, has been always attractive. Mechanical activation (MA) is a promising solid state method for preparing highly dispersed materials. It is a time-saving, energy efficient, economically convenient, ecologically clean and capable to scale-up method providing reduction of material cost. MA of mixtures involves the dispersion of solids and their plastic deformation. These processes generate defects in solids, accelerate the migration of defects in the bulk, increase the number of contacts between particles and renew the contacts, providing chemical interaction between solids. MA taking place in highly energetic planetary mills allows final products to be obtained either directly at close to room temperature or after a short-time one-step thermal treatment of activated mixtures at moderate temperatures, thus increasing phase homogeneity and decreasing particle size. It was revealed that hydroxides are usually more reactive in solid state synthesis compared with anhydrous oxides. They yield only water as a byproduct, giving a much ‘cleaner’ decomposition in contrast to metallorganic precursors, such as acetates or citrates (Pechini process) which produce large amounts of volatile organics on pyrolysis. According to Avvakumov et al.,1 hydroxides along with other solids containing oxygen and hydrogen groups (solid acids, acidic and basic salts, crystal hydrates, etc.) appear to be more reactive in mechanochemical reactions owing to the high reactivity of their surface oxide/hydroxide groups in proton and electron transfer and bond formation. This type of acid–base reaction has been called ‘soft mechanochemical synthesis’. The acid–base reaction is realized when more than two types of M–O(H) bond with differing acid–base properties are brought into contact. When using hydrated compounds with a hardness three to four times lower than that of anhydrous oxides, the level of mechanical loading and contamination decreases. The method consists of fast formation of amorphous or low-crystalline precursors during MA and their transformation into nanostructured final products under ‘soft’ heating conditions (lower temperatures and shorter time of treatment) after removal of water molecules, hydroxide groups and other volatile compounds. The present short review contains recent experimental results on mechanochemical interactions of LiOH (or Li2CO3) with anhydrous oxides, hydroxides, crystal hydrates of d-metals and acidic salts revealing the formation of precursors after a short time MA (1–10 min). The final products of these reactions – complex oxides with layered, spinel and framework structured compounds with polyanions – were prepared after a short-time

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heat treatment (1–4 hours) at moderate temperature. The experiments were performed using AGO-2 planetary mills with stainless steel, titanium or corundum milling bodies (830 rpm). The results might be interesting for the scientists who have been developing new methods of preparation of cathode, anode and electrolyte materials for rechargeable lithium batteries.

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Reactions of LiOH with anhydrous oxides

14.2.1 Synthesis of LiMn2O4 LiMn2O4 is an inexpensive and environmentally benign cathode material used for lithium-ion batteries with a practical capacity of about 120 mA h g−1 in the 4 V range. An important problem that prevents LiMn2O4 from wider use as a cathode is its unstable rechargeability, that is the amount of reversibly inserted lithium decreases gradually during cycling. At room temperature, LiMn2O4 has a cubic spinel-type structure, space group Fd3¯ m. The structure can be described as ideally consisting of closepacked arrangements of oxygen ions in 32e sites. Li+ ions occupy tetrahedral 8a sites, while Mn3+ and Mn4+ ions occupy octahedral 16d sites. The intercalation properties of LiMn2O4 have been widely studied.2,3 The extraction/ insertion of two lithium ions from/into the Li[Mn2]O4 spinel framework proceeds in two distinct steps. Lithium extraction/insertion from/into 8a tetrahedral sites occurs at around 4 V with maintenance of the initial cubic spinel symmetry, while that from/into the 16c octahedral sites occurs at around 3 V by a two-phase mechanism involving the cubic spinel Li[Mn2] O4 and the tetragonal lithiated spinel Li2[Mn2]O4. The cubic to tetragonal transition is due to Jahn–Teller distortion associated with Mn3+ (t2g3eg1) ions and is the origin of poor cycleability in the 3 V range. LiMn2O4 tends to exhibit capacity fade in the 4 V region as well, particularly at elevated temperatures (50°C). Several factors have been suggested to be the source of capacity fade, such as Jahn–Teller distortion occurring on the surface of the particles, manganese dissolution into the electrolyte, formation of two cubic phases, loss of crystallinity and the development of micro-strain during cycling. Electrochemical properties strongly depend on the method and conditions of synthesis, crystal and electronic structure, particle size, and so on. Soft mechanochemical synthesis of submicrometre LiMn2O4 was performed starting from LiOH and manganese oxides with different oxidation states of manganese ions (MnO2, Mn2O3, MnO).4,5 The interaction of lithium hydroxide with manganese oxides can be considered as acid–base reactions. According to Lewis classification, lithium hydroxide is a base whereas manganese oxides are acids with different strengths. Acid–base properties of

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14.1 X-ray patterns of activated mixtures of LiOH with different manganese oxides: (1) MnO, (2) Mn2O3, (3) MnO2. 䊉, LiMn2O4; 䉬, Mn2O3; 䉲, MnO; +, Mn3O4.

oxides significantly depend on the oxidation degree of the cation. The acidity increases along the series: MnO (2+) < Mn2O3 (3+) < MnO2 (4+). Reflections of LiOH are present in the background on X-ray diffraction (XRD) patterns of activated mixtures. No observable interaction occurs in the mixtures of LiOH with Mn2O3 or MnO (Fig. 14.1). The reflections of Mn2O3 remain practically unchanged after MA while a significant transformation of MnO to Mn3O4, an oxidation process, occurs. In contrast, MnO2 reacts almost completely forming cubic lithium–manganese spinel (probably, with a defective structure) after 10 min of MA without heating. The pattern can be indexed in the Fd3¯ m space group. The broadening of reflections is associated with reduced particle size and residual strain. On differential analysis (DTA) curves of the LiOH+MnO2 non-activated mixture, endothermic peaks corresponding to phase transitions of the reagents (TLiOH = 440°C, TMnO2 = 540°C) are present (Fig. 14.2). These effects practically disappear when the mixture was activated for 1 min. Moreover, on the DTA curve of the mixture activated for 10 min, a sharp exothermic peak at 485°C appears associated with crystallization of the final product, LiMn2O4. Thermogravimetric (TG) curves of activated mixtures have a

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sloping form. Mass loss, induced by elimination of gaseous products according to reaction [14.1] 4LiOH + 8MnO2 = 4LiMn2O4 + 2H2O + O2

[14.1]

becomes significantly lower than the theoretical one and is complete at a lower temperature (∼550°C) compared with a non-activated mixture. Solid state 7Li NMR (nuclear magnetic resonance) was used to study interaction in the LiOH+MnO2 mixtures activated by corundum milling bodies for 1–10 min. Strong shifts of 7Li lines to weaker magnetic fields appear in the NMR spectra, although no reflection of LiMn2O4 was observed in the X-ray patterns of these mixtures (Fig. 14.3). The value of the shift increases with the time of activation and does not coincide with that of the final paramagnetic crystalline product, LiMn2O4 (470 ppm).6 Such a shift was attributed to the appearance of paramagnetic Mn3+ and Mn4+ ions in the second coordination sphere of Li+ ions indicating the occurrence of intimate mixing of reagents (at the atomic level) and the beginning of spinel formation. According to photoelectron spectroscopy (XPS),6 the Li/Mn ratio on the surface of particles of the MnO2+LiOH mixture activated for 1 and 10 min is equal to 3.2 and 1.6, respectively, compared to 0.5 in the

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14.3 Li NMR spectra of the LiOH+MnO2 activated mixture. (—) 1 min MA, (—) 10 min MA, (---) spectrum of initial LiOH.

final spinel LiMn2O4, indicating that the surface layer of particles obtained is enriched with amorphous LiOH. As-prepared precursor accelerates the formation of crystalline LiMn2O4 with subsequent heat treatment. Spinels prepared at moderate temperatures (400–600°C) are characterized by high dispersion (secondary particles of about 50–100 nm) and a low cell parameter owing to increased numbers of Mn4+ ions in their structure. This results in high capacity and good structural stability of spinel when cycling in the 3 V range. As the temperature of heating increases to 750–800°C, the cell parameter of spinel increases to 8.24Å, indicating the average oxidation state of Mn ions to be +3.5, while particle size does not noticeably change. The electrochemical performance of these spinels at 4 V is similar to that of the ceramically prepared spinel.

14.2.2 Synthesis of LiNi0.5Mn1.5O4 Recently it has been shown that substitution of Mn ions for the other d-ions in LiMn2O4 results in the appearance of a 5 V plateau on the charge/ discharge curves. Among others, LiNi0.5Mn1.5O4 is of the most interest, since all nickel ions are in the 2+ oxidation state, while all manganese ions are in the 4+ state. Thus, a 5 V plateau is associated with the two-electron oxidation process Ni2+ → Ni4+, whereas Mn4+ ions remain electrochemically inactive. According to XRD, LiNi0.5Mn1.5O4 has a spinel-type structure (space group Fd3¯m) similar to LiMn2O4.7,8 However, its IR spectrum has a complex

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form indicating a more ordered structure described by the P4332 space group.9 Preparation of LiNi0.5Mn1.5O4 with this crystal and electronic structure, high discharge capacity and cycling stability will require the development of new synthesis methods. Earlier, using solution (sol–gel and combustion) methods, it was shown that fine particles are preferable for better LiNi0.5Mn1.5O4 electrochemistry. In this study, high-dispersed LiNi0.5Mn1.5O4 was prepared using soft mechanochemical synthesis. Figure 14.4 shows DTA and TG curves of non-activated and mechanically activated mixtures of LiOH, NiO and MnO2. Removal of adsorbed water occurs below 400°C for both mixtures. The corresponding mass loss for an activated mixture in this region exceeds that for a non-activated one. However, it stops as the temperature increases; moreover, one can observe a small mass gain, probably associated with oxidation of some Mn3+ ions, which appears during MA. For a non-activated mixture, in contrast, mass loss continues up to 565°C. New reflections appear on the X-ray pattern of the mixture activated for 10 min, which were assigned to lithium-rich manganese spinel (e.g. Li4Mn5O12), similar to results described above for the LiOH+MnO2 activated mixture (Fig. 14.5). However, the reflections of NiO remain, become less intensive and slightly wider. The formation of lithiated NiO with a cell parameter close to NiO is not excluded. The final product, LiNi0.5Mn1.5O4, with a submicrometre particle size is formed after heat treatment of the precursor at 900°C via interaction of lithium–manganese spinel with NiO. It has been shown that the structural and electrochemical properties of LiNi0.5Mn1.5O4 used as 5 V cathode material strongly depend on the regime of subsequent heat treatment: cooling rate, gas atmosphere, and so on.

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14.2.3 Synthesis of Li4Ti5O12 Recently, the quest to find a substitute for the graphite anode has received much attention. Lithium–titanium oxide Li4Ti5O12 or Li1.33Ti1.67O4 is one of the best candidates. It is a partly inverse spinel with a Fd3¯m space group, in which Li+ ions occupy both tetrahedral 8a and octahedral 16d sites: LiTd[Ti5/3Li1/3]OhO4.10–12 Li4Ti5O12 has a working voltage of 1.5 V and shows no volume changes during intercalation/deintercalation of lithium ions (‘zero strain’ material) providing a significant structural stability at cycling. It has been shown that the theoretical capacity can be achieved using fine particulate material. Below, the results of soft mechanochemical synthesis of Li4Ti5O12 are presented.

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Two endothermic peaks at 100 and 400°C are observed on DTA curves of non-activated mixture of LiOH with TiO2 (anatase). They are induced by removal of absorbed water and melting of LiOH, respectively (Fig. 14.6). In the case of a mixture activated for 10 min, the first peak becomes less intensive whereas the second one disappears evidencing the consumption of all LiOH below 400°C. Instead of a stepped mass loss (at 20–100°C, 100–350°C and 350–650°C) for the non-activated mixture, the TG curve of the activated mixture has a smooth form. Mass loss is already complete at ∼350°C. Its value considerably decreases, evidencing a significant degree of interaction, associated with water elimination during MA. According to X-ray data, MA of the LiOH and TiO2 mixture is accompanied by full and partial amorphization of LiOH and TiO2, respectively (Fig. 14.7). New reflections were assigned to monoclinic Li2TiO3 (space group C2/c). The final single-phase product, Li4Ti5O12, is formed after a short-time heating of as-prepared precursor at 800°C. Li4Ti5O12 anode material with enhanced properties was prepared after additional short-time grinding, especially in the presence of carbon as an electron-conductive additive, thus forming Li4Ti5O12/C composite material.13

14.2.4 Synthesis of Li1+xV3O8 Li1+xV3O8 has been widely studied as a promising 3 V cathode material because of its high discharge capacity (up to 250 mA h g−1), facile preparation and stability in air. Li1+xV3O8 is a bronze-type compound in which some of the V5+ ions are reduced to V4+. Li1+xV3O8 crystallizes in a monoclinic

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system (space group P21/m). The structure of Li1+xV3O8 consists of double and single zigzag octahedral strings arranged to provide interlayer sites (tetrahedral and octahedral) for Li ions. It is achieved by expanding the framework of vanadium polyhedrons.14–17 The synthesis conditions were shown to induce important differences in the reversible capacity and cycling behaviour of Li1+xV3O8. Enhanced electrochemical characteristics were observed for low-crystalline materials with increased specific surface area and expanded interlayer spacing which causes an increase in Li-ion mobility. In this study, LiOH and V2O5 were used as initial reagents to prepare Li1.07V3O8. Mechanical activation was carried out for 30 s to 10 min with corundum jars and balls in air.18,19 The thermal effects of phase transfor-

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14.8 DTA and TG curves of the LiOH + V2O5 mixture activated for 10 min.

mations of reagents (TLiOH = 445°C, TV2O5 = 690°C) are absent on DTA curves of the LiOH + V2O5 mixture activated for 10 min (Fig. 14.8). Instead, an endothermic peak appears at 20–250°C accompanied by release of weakly bonded H2O and CO2. It transforms into a strong exothermic peak, associated with crystallization of the final product. The next endothermic effect at 590°C is due to melting of Li1+xV3O8. It is accompanied by mass gain (∼0.5 mass %) as a result of oxidation process of some of the V4+ ions. Figure 14.9 shows X-ray patterns of precursors formed after activation of the LiOH + V2O5 mixture for 30 s to 10 min and heat treated at 400°C. It is seen that the intensity of V2O5 reflections decreases if the milling time increases. Reflections of the initial LiOH are present in the background owing to its amorphization and weak scattering ability of Li atoms. One can see the appearance of new weak reflections assigned to lithium–vanadium bronzes after MA for only 1 min. After 10 min, the reflections of the final product, Li1+xV3O8, are present in the pattern, indicating that chemical interaction and crystallization occur at the early stage of MA. In infrared (IR) spectra of the LiOH + V2O5 activated mixtures, the intensity of absorption band at 1020 cm−1, corresponding to stretching vibrations of double vanadyl bonds ν (V苷O) in V2O5, decreases versus time of MA (Fig. 14.10). New bands characteristic of Li1+xV3O8 appear in the 900– 1000 cm−1 range. After heating the activated mixtures at 400°C, all reflections on the X-ray patterns belong to Li1+xV3O8. Note that only partial interaction occurs when

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14.9 X-ray patterns of the LiOH + V2O5 mixtures activated for 30 s (1), 1 min (2), 5 min (3), 10 min (4), 10 min followed by annealing to 400°C (5), 5 min followed by ageing for 6 months (6). 䊉, V2O5; 䉬, LiV3O8; 䉲, Li–V bronzes.

a non-activated mixture was heated to temperatures below the melting point of V2O5. Thus, MA of mild and basic LiOH with hard and acidic V2O5 results in fast deformation mixing of reagents at atomic level and chemical interaction caused by mechanical and chemical forces. LiOH probably also acts as a surfactant towards V2O5, accelerating its comminution. Surprisingly, X-ray patterns and IR spectra of the sample, activated for 5 min, significantly change after ageing the sample for 6 months at room temperature (without heating). X-ray reflections of V2O5 sharply decrease, while those of Li1+xV3O8 increase, evidencing the occurrence of chemical reactions

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14.10 IR spectra of the LiOH + V2O5 mixtures activated for 30 s (1), 1 min (2), 5 min (3), 10 min (4), 10 min followed by annealing at 400°C (5), 5 min followed by ageing for 6 months (6).

by the relaxation of accumulated mechanical energy and diffusion processes in the solid state at room temperature (Fig. 14.9). Simultaneously, the intensity of the IR bands, corresponding to Li1+xV3O8, increases; a complex line in the 900–1050 cm−1 range become narrow; a clear maximum at 950 cm−1 with a shoulder at 970 cm−1 appears. The reduction process V5+ → V4+ during MA and ageing of activated mixture was confirmed by EPR spectroscopy.18 The final product Li1+xV3O8 was prepared by heat treatment of the precursor at 400–600°C. Its electrochemical capacity was about 300 mA h g−1 when discharging to 1.5 V; discharge curves have no distinct plateaus like ceramically prepared samples.

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Reactions of LiOH with solid hydroxides

14.3.1 Synthesis of LiCoO2 LiCoO2 is the most suitable lithium intercalation material for today’s high density and long life lithium-ion batteries. However, it is also the most expensive. High-temperature fabrication processes result in significant costs in terms of energy usage. This has led many researchers to seek lower temperature fabrication routes, resulting in fine particulate material. LiCoO2 can be prepared in two modifications: a high-temperature (HT) and a low-temperature (LT) one.20–23 The structure of HT-LiCoO2 is related to the trigonal space group R-3m with oxygen anions occupying 6c sites and Li and Co ions occupying 3a and 3b octahedral sites, respectively. Li layers alternate with Co ones between layers of oxygen. LT-LiCoO2 is believed to adopt a spinel-like structure, consisting of a cubic close packed oxygen network with alternating cation layers with compositions 0.75 Co, 0.25 Li and 0.75 Li, 0.25 Co perpendicular to each of the cubic (111) directions. The space group is Fd3¯m; Li and Co ions are in 16c and 16d sites, respectively and O ions are in 32e sites. Unfortunately, only 50% of the theoretical capacity of Li1−xCoO2 could be practically utilized. This corresponds to a reversible extraction/insertion of 0.5 lithium ions per cobalt and a practical capacity of 120–140 mA h g−1. Capacity fade occurs at (1 − x) < 0.5. The limitation of practical capacity has been attributed to an ordering of lithium ions and consequent structural distortion around x = 0.5. LiCoO2 was prepared using preliminary MA of LiOH with Co(OH)2.24,25 The time of MA varied from 1 to 10 min. In reaction with LiOH, cobalt hydroxide exhibits its amphotheric (acidic) properties. Moreover, the synthesis of layered LiCoO2, using Co(OH)2, is more efficient compared with Co3O4 owing to the structural similarity of the initial hydroxide and final LiCoO2 and the large specific surface area of the hydroxide. Interaction of LiOH with Co(OH)2 to prepare LiCoO2 should be accompanied by oxidation of Co2+ to Co3+ ions. Therefore, additional oxygen is needed to provide reaction [14.2]: 4LiOH + 4Co(OH)2 + O2 = 4LiCoO2 + 6H2O.

[14.2]

TG–DTA data highlight an appreciable difference in the thermal behaviour of the activated and non-activated mixtures. Heating the non-activated LiOH+Co(OH)2 mixture is accompanied by three endothermic peaks: two peaks in the low-temperature region (below 300°C) and one peak at 400– 500°C. They were assigned: (i) to the removal of adsorbed water from LiOH in the course of the sample preparation; (ii) to the dehydration of the initial cobalt hydroxide; and (iii) to LiOH melting, respectively. In

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contrast, the DTA curve of the activated mixture is less busy presenting a single endothermic peak at 50–200°C, associated with the removal of gaseous products. Thermal effects associated with the formation and crystallization of LiCoO2 were absent in both cases. Only broadened reflections of Co(OH)2 are observed on the X-ray patterns of activated mixtures, while the reflections of LiOH are absent, probably owing to its amorphization and weak scattering ability of Li ions (Fig. 14.11). Partial oxidation of Co(OH)2 with the formation of CoOOH occurs. In the IR spectrum of the LiOH + Co(OH)2 activated mixture, the intensity of the bands at 307 and 490 cm−1, corresponding to stretching vibrations of Co2+—O bonds in the Co2+O6 octahedra of the Co(OH)2 structure, as well as those of hydroxide groups of the initial reagents at 3630 and 3680 cm−1, sharply decreases (Fig. 14.12). New bands in the 500–700 cm−1 range, corresponding to stretching vibrations of Co3+—O bonds in the Co3+O6 octahedra appear, evidencing the beginning of LiCoO2 formation. Correspondingly, in electronic diffuse reflectance spectra (EDRS) of the LiOH + Co(OH)2 activated mixture, the overall absorption above 15000 cm−1 increases as a result of strong distortion of the Co(OH)2 structure (Fig. 14.13). A new band at 25000 cm−1, corresponding to low-spin [Co3+]Oh ions, appears, indicating partial oxidation of Co2+ to Co3+ followed by transformation of Co(OH)2 to CoOOH (but not to Co3O4). As-prepared precursors were subsequently annealed at 400°C, 600°C and 800°C for 1–4 hours in air. The annealing results in fast chemical interaction with the formation of homogeneous product, LiCoO2. The width of X-ray reflections decreases and their intensity increases with respect to the temp-

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14.11 X-ray patterns of LiOH + Co(OH)2 mixtures activated for 1 min (2) and 10 min (3) compared with initial Co(OH)2 (1). 䊉, Co(OH)2; 䉬, CoOOH; 䊏, LiCoO2.

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14.12 IR spectra of the LiOH + Co(OH)2 mixture activated for 1 min (2) and 10 min (3) compared with initial Co(OH)2 (1).

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14.13 Diffuse reflectance spectra of LiOH + Co(OH)2 mixtures activated for 1 min (2) and 10 min (3) compared with initial Co(OH)2 (1). F(R) is the Kubelka–Munk function.

erature of heat treatment as a result of the processes of crystallization and particle growth. Splitting of 006 and 012 and 018 and 110 reflections, as well as the value of c/a ratio equal to 4.99, are evidence of the formation of HT-LiCoO2 even at 400°C. Note that annealing the non-activated mixture at 400°C leads to formation of the mixture of HT- and LT-LiCoO2. The electrochemical performance of LiCoO2 cathode material prepared using

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MA and heat treatment at high temperatures is similar to materials prepared by the ceramic method, although its synthetic process is more facile, time- and energy-conserving.

14.2.2 Synthesis of LiNi1−x−yCoxMnyO2 Cathode materials based on solid solutions of LiNi1−yCoyO2 are less expensive than LiCoO2 and show a larger discharge capacity at 4 V. Their discharge curves have a sloping form. It has been found that Ni3+ ions are oxidized first on charging LixNi1−yCoyO2 (in the 0.85 < x < 1 range), while Co3+ ions are oxidized at a higher voltage. In the LiCoO2–LiNiO2 system, a continuous row of solid solutions is formed. Depending on the temperature of synthesis, they can adopt either a HT- or a LT-structure, similar to LiCoO2.26–28 Introduction of Mn in LiNi1−yCoyO2 was proposed to improve further the electrochemical and thermal stability. LiNi1−x−y CoxMnyO2 materials are viewed as candidates for large-scale batteries for electric vehicles. LiNi1−x−yCoxMnyO2 materials are usually considered to be solid solutions in the LiCoO2–LiNi0.5Mn1.5O2 system with Ni2+, Co3+ and Mn4+ ions.30,31 The synthesis conditions are important in determining the metal ordering in the LiNi1−yCoyO2 and LiNi1−x−yCoxMnyO2 layered structure and in the subsequent electrochemistry. Thus, these compounds are best formed by mixing Ni, Co and Mn salts in aqueous solution, precipitation by LiOH and corresponding firing. In this study MA was applied as an alternative to the solution method. Solid solutions of LiNi1−yCoyO2 (y = 0.2, 0.4, 0.6 and 0.8) were prepared by MA of mixtures of LiOH with double Ni1−yCoy-hydroxides.29 Figure 14.14 shows DTA and TG curves of activated and non-activated mixtures as well as of the double Ni0.8Co0.2-hydroxide alone. This hydroxide loses water in two steps accompanied by endothermic effects at 100°C and at 220–290°C (with minimum at 275°C) owing to removal of absorbed and structural water, respectively. When heated in a mixture with LiOH, the temperature of the latter effect decreases by ∼30°C in the initial mixture and by ∼80°C in the activated mixture. The TG curve of the activated mixture has a smoother form; mass loss is completed at lower temperature. The endothermic peak associated with melting of non-reacted LiOH (∼400°C) significantly decreases. A similar effect was observed for the mixtures of LiOH with Ni1−x−y CoxMny hydroxides. Triple hydroxides lose water in the 100–400°C temperature range accompanied by two endothermic effects at 100°C and 290°C and a stepped mass loss.32 In both non activated and activated mixtures, these effects shift by ∼20° to the lower temperature region, while the effect associated with LiOH melting disappears.

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14.14 DTA and TG curves for initial (---) and activated (—) mixtures of LiOH with Ni0.8Co0.2(OH)2 compared with the double hydroxide (......).

Since the average oxidation state of d-ions in the final products is higher than in initial hydroxides, the process of LiNi1−yCoyO2 and LiNi1−x−yCox MnyO2 synthesis should be accompanied by oxidation of d-ions. However, we did not observe exothermic effects on the DTA curves associated with oxidation or with the crystallization of final products. It should be noted that the total mass loss for both mixtures is practically identical. The water eliminated in the course of dehydration of d-metal hydroxides, probably, enters intermediate products. According to X-ray analysis, formation of the final products is observed after heating these precursors to 400°C, however their crystallinity is poor (Fig. 14.15). Final crystallization occurs at 750°C for LiNi1−yCoyO2 and at 850°C for LiNi1−2yCoyMnyO2 independently of y.33 All as-prepared materials are defined as single-phase HT-modifications. They are characterized by a specific discharge capacity of 170 mA h g−1 for LiNi0.8Co0.2O2 and LiNi0.6Co0.2Mn0.2O2 and good rechargeability.

14.4

Reactions of LiOH and Li2CO3 with crystal hydrates and acidic salts

14.4.1 Synthesis of LiFePO4 Lithium metal phosphates are among the most promising cathode materials for high capacity lithium-ion batteries. Their promise is due to the ‘inductive effect’ of the PO4 polyanion, which elevates the M2+/M3+ redox couple by about 1.5–2 V. A redox potential of 3.45 V vs. Li/Li+ results in the case of LiFePO4, making it a particularly appealing material for hybrid energy

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018 110 113

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Li2Ni8O10 [ASTM 23-362]

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14.15 X-ray patterns of LiNi1−yCoyO2, prepared by annealing activated mixtures at 750°C: y = 0.2 (1), 0.4 (2), 0.6 (3) and 0.8 (4). Trace (5) is the LiOH + Ni0.8Co0.2(OH)2 mixture activated for 5 min and annealed at 400°C.

systems where cost and safety are of major concern. LiFePO4 adopts an olivine-type structure, built on an oxygen hexagonal packing in which Li+ and Fe2+ ions occupy half of the octahedral sites and P5+ ions occupy 1/8 of the tetrahedral sites (space group Pmnb). The peculiar distribution of Li+ and Fe2+ within the octahedral sites generates MO6 layers.34–36 Owing to the inherently low electronic conductivity of LiFePO4, it is essential to optimize its properties: to minimize the defect concentration and crystalline size (down to submicron level), control morphology, and so on. MA is one of the most convenient methods of preparing LiFePO4 with enhanced electrochemical performance. Four endothermic peaks at 108°C, 188°C, 223°C and 402°C are present on a DTA curve of non-activated mixture of LiOH with FeC2O4⋅2H2O and NH4H2PO4 heated in Ar atmosphere, associated with dehydration and decomposition of FeC2O4⋅2H2O and NH4H2PO4 and melting of LiOH, respectively (Fig. 14.16). The same peaks are observed for the mixture, activated in Ar for 5 min, however, the peak at 188°C is less intensive and the peak at 402°C practically disappears. A new low-intensive endothermic peak is present at ∼255°C. The TG curve adopts a smooth form. Whereas the crystallization of the final product in the non-activated mixture occurs at about 500°C followed by a low intensive exothermic peak, a distinct peak at 443°C is observed for the activated mixture. A decrease in the crystallization temperature of LiFePO4 is an important factor in preparing cathode material with a smaller grain size.37

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The reflections of FeC2O4⋅2H2O and NH4H2PO4 on the X-ray pattern of the activated mixture become broader and less intensive, while the reflections of LiOH completely disappear (Fig. 14.17). New weak reflections assigned to Li3PO4 as an intermediate product are present. This indicates that the initial interaction occurs during MA according to Reaction [14.3]: NH4H2PO4 + 3LiOH → Li3PO4 + NH3 + 3H2O

[14.3]

Figure 14.18 shows the Mössbauer spectrum of the activated mixture. The experimental curve can be decomposed into two components: FeC2O4⋅2H2O and Fe3O4. This correlates with the scheme for thermal decomposition of FeC2O4⋅2H2O proposed by Hermanek:38 FeC2O4⋅2H2O → FeC2O4 + 2H2O (170–230°C)

[14.4]

3FeC2O4 → Fe3O4 + 4CO + 2CO2 (230–390°C)

[14.5]

3FeC2O4 + 2CO → Fe3C + 7CO2 (360–410°C)

[14.6]

Fe3C → 3Fe + C (410–530°C)

[14.7]

443

Accordingly, LiFePO4 is formed under subsequent heating via interaction of Li3PO4 with Fe3O4. Crystallization of LiFePO4 occurs after heat treatment of the activated mixture at 450°C. The crystallinity of the final product does not noticeably change up to 600°C, while its crystallinity increases sharply after heating to 700°C. According to Mössbauer spectroscopy, as the temperature increases, the amount of Fe3+ impurities in LiFePO4 decreases. The specific discharge capacity of as-prepared LiFePO4

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14.16 DTA and TG curves of initial (---) and activated (—) mixtures of LiOH with FeC2O4⋅2H2O and NH4H2PO4.

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(112)

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(031)

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14.17 X-ray patterns of activated mixture LiOH with FeC2O4⋅2H2O and NH4H2PO4 (1) and after heat treatment at different temperatures: (2) 450°C, (3) 500°C, (4) 550°C, (5) 600°C, (6) 700°C.

is about 125–130 mA h g−1, while that of mechanochemically prepared composite cathode of LiFePO4 with carbon as an electronic conductive agent is more than 150 mA h g−1.

14.4.2 Synthesis of Li2FeSiO4 In the on-going search for alternative cathode materials, the Li–Fe–Si–O system holds great promise since one can expect the same lattice stabilization effect (‘inductive effect’) as in LiFePO4 through the presence of strong Si–O bonds. The lower electronegativity of Si versus P will result in a lowering of the Fe3+/Fe2+ redox couple. Such materials should therefore have a lower electronic band gap and therefore a higher electronic conductivity. A

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Experimental spectrum Fe2+ + Fe3+ Fe2+ Fe3+

Component

% Width Chemical Quadrupole shift splitting 1.768 FeC2O4•2H2O 58 0.299 1.187 31 0.383 0.530 0.264 Fe3O4 11 0.370 0.080 Fe

14.18 Mössbauer spectrum of the Li2CO3, FeC2O4⋅2H2O and NH4H2PO4 mixture activated for 5 min in Ar atmosphere and its decomposition.

material based on Fe and Si would also be potentially low cost in terms of raw material costs. Indeed, Fe- and Si-oxides represent >10% of the Earth’s crust. Iron orthosilicate Li2FeSiO4 has been reported to crystallize in the orthorhombic system with a Pmn21 space group, isostructural to Li3PO4. Li, Fe and Si atoms are tetrahedrally coordinated to four oxygen atoms.39–41 Reducing the particle size and optimizing its distribution will certainly improve the electrochemical performance of this cathode material. Compared with other methods, MA is very attractive in this case. Li2FeSiO4 is usually prepared starting from Li2SiO3 and FeC2O4⋅2H2O. We used a mixture of Li2CO3, SiO2⋅xH2O (silica gel) and FeC2O4⋅2H2O. Four endothermic peaks are present on DTA curves of a non-activated mixture heated in Ar atmosphere, associated with dehydration and decomposition of SiO2⋅xH2O and FeC2O4⋅2H2O (at 100°C, 195°C and 397°C), and melting of Li2CO3 (at 716°C) followed by three-step mass loss (in the RT– 200°C, 200–400°C and 400–720°C regions), while only two peaks are observed for the activated mixture (in the RT–150°C and 350–450°C regions) followed by two steps on the TG curves (in the RT–250°C and 250–500°C regions) (Fig. 14.19). The total mass loss for the activated mixture is slightly lower. According to X-ray analysis, amorphization of the initial reagents is observed after 2.5 min of MA. When as-prepared precursor was heated at 600°C for 4 hours, the synthesis was completed and crystallization of the final product occurred,

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14.19 DTA curves of initial (---) and activated (—) mixtures of Li2CO3 with FeC2O4⋅2H2O and SiO2⋅0.1H2O.

while ≥20 hours were needed when low-energy ball milling is used. Thus, preliminary MA, resulting in intimate mixing and amorphization of reagents significantly accelerates formation of Li2FeSiO4 under further heating, keeping the small particle size.

14.4.3 Synthesis of LiTi2(PO4)3 Recently, high energy density and long life all solid state lithium-ion batteries using solid lithium-ion electrolytes have been requested. NASICONtype materials (isostructural with NaZr2(PO4)3) are good ionic conductors with negligible electronic conductivity and are stable in air. Among them, LiTi2(PO4)3 is one of the most promising electrolytes. The structure of LiTi2(PO4)3 (space group R-3c) consists of a three-dimensional network made up of TiO6 octahedra sharing all their corners with PO4 tetrahedra and vice versa to form so-called ‘lantern’ units, all oriented in the same direction (along the c-axis).42,43 It has been shown that conduction at grain boundaries is the rate determining step for the lithium-ion migration at room temperature and significantly depends on the method of synthesis, as well as on the nature and the number of dielectric impurities. In order to prepare LiTi2(PO4)3, a stoichiometric mixture of LiOH, TiO2 and NH4H2PO4 was initially preheated at 400°C to decompose NH4H2PO4. As-prepared precursor was then activated for 5 min with Ti milling bodies. Two endothermic peaks are present on the DTA curve of the non-activated mixture. They were assigned to decomposition of NH4H2PO4 (200°C) and to melting of LiOH (400°C) accompanied by mass loss up to 600°C (Fig. 14.20). In contrast, these peaks disappear on the DTA curves of the

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14.20 DTA and TG curves of initial (---) and activated (—) LiOH + TiO2 + NH4H2PO4 mixture.

activated precursor; mass loss significantly decreases indicating the occurrence of the reaction. Figure 14.21 shows X-ray patterns of the preheated mixture and the products of its activation and heat treatment at 800°C. When heat treated at 400°C, the reflections of LiOH disappear, while the reflections of TiO2 become less intensive and broaden out. Besides, the reflections of a new phase – TiP2O7 – are present as a result of interaction of TiO2 with NH4H2PO4. After MA of the preheated mixture, the reflections of the final product – LiTi2(PO4)3 – appear. It crystallizes during annealing at elevated temperatures (800–1000°C). While TiP2O7 is the main impurity for ceramically prepared LiTi2(PO4)3, only tiny amounts of TiO2 were observed for that prepared using MA.44,45 Thus, as-prepared ceramic and MA samples of LiTi2(PO4)3 are different in nature and in the number of dielectric admixtures they form. It has been shown that lithium ionic conductivity of mechanochemically prepared LiTi2(PO4)3 is characterized by a significant reduction in grain boundary resistance compared with the ceramically prepared sample, while the bulk conductivity remains unchanged.

14.5

Conclusions

The peculiarities of soft mechanochemical reactions in mixtures of LiOH (or Li2CO3) with anhydrous oxides of d-metals (MnO, Mn2O3, MnO2, NiO, TiO2, V2O5), solid hydroxides (Co(OH)2, Ni1−x−yCoxMny(OH)2), crystal hydrates (FeC2O4⋅2H2O) and acidic salts (NH4H2PO4) can be summarized as follows.

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LiTi2(PO4)3 TiO2, anatase TiP2O7

3

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50

60

2θ (degree)

14.21 X-ray patterns of the LiOH + TiO2 + NH4H2PO4 mixture preheated at 400°C (1), activated (2) and heated at 800°C (3).

1. All mechanochemical reactions are characterized by fast formation of amorphous or low-crystalline nanostructured precursors. These precursors are shown to consist of intimately mixed mechano(nano)composites of LiOH with another component. In the presence of the other component fast and almost complete amorphization of composites of LiOH (or Li2CO3) occurs. The structure of solid hydroxides, crystal hydrates and acidic salts being activated in the mixtures with LiOH (or Li2CO3) undergo significant distortion followed by a decrease in the dehydration and decomposition temperature. 2. Highly reactive precursors accelerate the formation of low-crystalline final (LiMn2O4 in the LiOH + MnO2 mixture; LiV3O8 in the LiOH + V2O5 mixture) or intermediate (e.g. Li3PO4 in the LiOH + FeC2O4⋅2H2O + NH4H2PO4 mixture; Li2TiO3 in the LiOH + TiO2 mixture) products

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during MA. Moreover, in some cases (e.g. LiOH + V2O5), the accumulated mechanical energy relaxes in chemical energy with formation of the product even during aging at room temperature due to fast diffusion processes. 3. Fast chemical interactions in the above mixtures are provided by intimate mixing of reagents and high reactivity of surface functional groups, notably, OH groups. They can be considered as acid–base reactions. Although most metal oxides absorb water molecules at their surface, replacing anhydrous oxides by hydroxides or crystal hydrates results in much higher OH density per surface area. The bond energy of Me–O and O–H is a measure of the acidity of the OH-group. The acid–base reaction is realized when more than two types of Me–O(H) bonds with differing acid–base properties are brought into contact. Thus, the efficiency of chemical interaction in mixtures of LiOH with d-metal compounds depends on their acidic properties. It is evident that the highest output might be expected in mixtures exhibiting a maximum difference in the acid–base properties of the compounds. The strength of solid acids and bases can be characterized by proton affinity (PA). Paukschtis46 has developed a method to determine PA based on IR spectroscopy of adsorbed molecules. In general, the changes in the acidity of OH groups on the oxide surface follow the regularities governed by the Periodic law: the higher the position occupied by an element in the Mendeleev’ periodic table (for each group), the higher is the proton donor ability of the surface. For each period, the acidity of groups increases with the element number. Acidic properties of d-metal oxides significantly increase with the degree of oxidation of the cation. Indeed, the highest acidity of MnO2, compared to MnO and Mn2O3, results in its noticeable efficiency in reaction with LiOH during MA leading to the formation of a low-crystalline final product, lithium– manganese spinel. On the other hand, when LiOH is activated with TiO2, the Li4Ti5O12 spinel is not formed because of the lower acidity of TiO2. The acidity of V2O5 is comparable with that of MnO2; therefore MA of V2O5 with LiOH results in the formation of LiV3O8 with a monoclinic structure. Thus, acid–base properties of reagents are more significant in soft mechenochemical reactions than the structure of the final product. 4. Crystalline nanostructured and homogeneous Li-containing final products are formed via transformation of as-obtained amorphous or lowcrystalline precursors under ‘soft’ heating conditions (lower temperatures and a shorter treatment time) after removal of water molecules, hydroxide groups and other volatile compounds. These groups inhibit the growth of particles under subsequent heat treatment of precursors. Therefore, grain growth is less significant than in conventional solid state synthesis.

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5. All Li-containing compounds obtained by the reactions described above show good electrochemical properties as cathode, anode or solid electrolyte materials for lithium-ion batteries. In addition, as an energysaving fabrication route, soft mechanochemical synthesis results in a significant reduction of material cost.

14.6

Future trends

This short review reveals the principles of ‘soft mechanochemical synthesis’, generally based on the acid–base properties of the reagents. The method can be widely used for the synthesis of different functional inorganic materials. Making a choice of reagents and tuning conditions of MA one can vary the particle size and morphology, crystal and electronic structure, transport properties and, as a result, functional properties of the asprepared materials. In any case, this method has noticeable advantages over other methods of preparing high-dispersed materials, including simplicity and accessibility and should find successful industrial application in the field of materials science in the near future.

14.7

Acknowledgements

The author would like to thank E. Devyatkina, I. Asanov, S. Kozlova, S. Gabuda, V. Anufrienko, V. Kaichev, A. Slobodyuk, A. Titov, N. Uvarov, A. Stepanov and A. Buzlukov for their contribution to this study.

14.8

References

1. avvakumov e, senna m and kosova n (2001), Soft Mechanochemical Synthesis. A Basics for New Chemical Technologies, Kluwer, Dordrecht. 2. tarascon jm and guyomard d (1991), ‘Li Metal-free rechargeable batteries based on Li1+xMn2O4 cathodes (0 ≤ x ≤ 1)’, J. Electrochem. Soc., 138, 2864–8. 3. thackeray mm, de kock a, rossouw mh, liles d, bittihn r and hoge d (1992), ‘Spinel electrodes from the Li–Mn–O system for rechargeable lithium battery applications’, J. Electrochem. Soc., 139, 363–6. 4. kosova nv, uvarov nf, devyatkina et and avvakumov eg (2000), ‘Mechanochemical synthesis of LiMn2O4 cathode material for lithium batteries’, Solid State Ionics, 135, 107–14. 5. kosova nv, devyatkina et and kozlova sg (2001), ‘Mechanochemical way for preparation of disordered lithium-manganese spinel compounds’, J. Power Sources, 97–8, 406–11. 6. kosova nv, asanov ip, devyatkina et and avvakumov eg (1999), ‘State of manganese atoms during the mechanochemical synthesis of LiMn2O4’, J. Solid State Chem., 146, 184–8. 7. ohzuku t, takeda s and iwanaga m (1999), ‘Solid-state potentials for Li[Me1/2Mn3/2]O4 (Me: 3d-transition metal) having spinel-framework structures:

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26. rougier a, saadoune i, gravereau p, willmann p and delmas c (1996), ‘Effect of cobalt substitution on cationic distribution in LiNi1−yCoyO2 electrode materials’, Solid State Ionics, 90, 83–90. 27. delmas c, menetrier m, croguennec l, saadoune i, rougier a, poillerie c, prado g, grune m and fournes l (1999), ‘An overview of the Li(Ni,M)O2 systems: synthesis, structures and properties’, Electrochim. Acta., 45, 243–53. 28. zhecheva e and stoyanova r (1993), ‘Stabilization of the layered crystal structure of LiNiO2 by Co-substitution’, Solid State Ionics, 66, 143–9. 29. kosova nv, devyatkina et and slobodyuk ab (2008), ‘Structure and electrochemical properties of LiNi1−yCoyO2 solid solutions, prepared using mechanical activation’, Chem. Sustainable Development, 17, 141–9. 30. yabuuchi n and ohzuku t (2003), ‘Novel lithium insertion material of LiCo1/3Ni1/3Mn1/3O2 for advanced lithium-ion batteries’, J. Power Sources, 119–21, 171–4. 31. koyama y, makimura y, tanaka i, adachi h and ohzuku t (2004), ‘Systematic research on insertion materials based on superlattice models in a phase triangle of LiCoO2–LiNiO2–LiMnO2. I. First-principles calculation on electronic and crystal structures, phase stability and new LiNi1/2Mn1/2O2 material’, J. Electrochem. Soc., 151, A1499–A1506. 32. kosova nv, devyatkina et and kaichev vv (2007), ‘Mixed layered Ni–Co–Mn hydroxides: crystal structure, electronic state of ions, and thermal decomposition’, J. Power Sources, 174, 735–40. 33. kosova nv, devyatkina et and kaichev vv (2007), ‘Optimization of Ni2+/Ni3+ ratio in layered Li(Ni,Mn,Co)O2 cathodes for better electrochemistry’, J. Power Sources, 174, 965–9. 34. padhi k, nanjundaswamy ks and goodenough jb (1997), ‘Phospho-olivins as positive-electrode materials for rechargeable lithium batteries’, J. Electrochem. Soc., 144, 1188–94. 35. andersson as and thomas jo (2001), ‘The source of first-cycle capacity loss in LiFePO4’, J. Power Sources, 97–8, 498–502. 36. yamada a, chung sc and hinokuma k (2001), ‘Optimized LiFePO4 for lithium battery cathodes’, J. Electrochem Soc., 148, A224–A229. 37. kosova nv and devyatkina et (2004), ‘On mechanochemical preparation of materials with enhanced characteristics for lithium batteries’, Solid State Ionics, 172, 181–4. 38. hermanek m, zboril r, mashlan m, machala l and schneeweiss o (2006), ‘Thermal behaviour of iron(II) oxalate dihydrate in the atmosphere of its conversion gases’, J. Mater. Chem., 16, 1273–80. 39. nytén a, abouimrane a, armand m, gustafsson t and thomas jo (2005), ‘Electrochemical performance of Li2FeSiO4 as a new Li-battery cathode material’, Electrochem. Comm., 7, 156–60. 40. nytén a, kamali s, häggström l, gustafsson t and thomas jo (2006), ‘The lithium extraction/insertion mechanism in Li2FeSiO4’, J. Mater. Chem., 16, 2266–72. 41. zaghib k, salah aa, ravet n, mauger a, gendron f and julian c (2006), ‘Structural, magnetic and electrochemical properties of lithium iron orthosilicate’, J. Power Sources, 160, 1381–6. 42. aono h, sugimoto e, sadaoka y, imanaka n and adachi g (1989), ‘Ionic conductivity of the lithium titanium phosphate (Li1+xMxTi2−x(PO4)3, M = Al, Sc, Y, and La) systems’, J. Electrochem. Soc., 136, 590–1.

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43. arbi k, mandal s, rojo jm and sanz j (2002), ‘Dependence of ionic conductivity on composition of fast ionic conductors Li1+xTi2−xAlx(PO4)3, 0 ≤ x ≤ 0.7. A parallel NMR and electric impedance study’, Chem. Mater., 14, 1091–7. 44. kosova n, devyatkina e and osintsev d (2004), ‘Dispersed materials for rechargeable lithium batteries: Reactive and non-reactive grinding’, J. Mater. Sci., 39, 5031–6. 45. kosova nv, devyatkina et, stepanov ap and buzlukov al (2008), ‘Lithium conductivity and lithium diffusion in NASICON-type Li1+xTi2−xAlx(PO4)3 (x = 0; 0.3) prepared by mechanical activation’, Ionics, 14, 303–11. 46. paukschtis e (1992), Infrared Spectroscopy in Heterogeneous Acid–Base Catalysis, Nauka, Novosibirsk (in Russian).

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15 Materials for lithium-ion batteries by mechanochemical methods L. C. YA N G, Q. T. Q U, Y. S H I, and Y. P. W U, Fudan University, China; and T. VA N R E E, University of Venda, South Africa

Abstract: Lithium-ion batteries have many advantages over traditional rechargeable batteries and their development has been very rapid. In this chapter, preparation and electrochemical performance of their key materials including cathode materials such as LiCoO2 and LiMn2O4, anode materials such as carbon, alloys and nitrides, and electrolytes such as oxides and sulfides by mechanochemical (MC) methods are primarily summarized. Compared with conventional solid state reactions at high temperature, the MC methods appear to accelerate and simplify the synthesis process and decrease the energy expenses as well as the cost of the material. In addition, the prepared materials present good electrochemical performance. When MC methods are combined with other techniques, their advantages can be more fully displayed. In the meanwhile, MC reactions will have some unfavourable actions to some materials which should be avoided. Finally, some further aplications for MC methods in lithium-ion batteries are pointed out. Key words: mechanochemical method, lithium-ion battery, anode, cathode, electrolyte, electrochemical performance, capacity, cycling behaviour.

15.1

Introduction

After several decades of research and development, lithium-ion batteries came into use in the late 1980s and early 1990s (Wu et al., 2004). Since then their development has been very rapid owing to their evident advantages over traditional rechargeable batteries: • • • • • • • •

high output voltage (average 3.6–3.7 V or 3.2–3.3 V) and power high energy density (>180 Wh kg−1) low self discharge (<3% per month) no memory effect long cycle life (>1000 times) high rate capability (1 C) high coulomb efficiency (near 100% after the 1st cycle) easy measurement of the residual capacity 361 © Woodhead Publishing Limited, 2010

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maintenance free no environmental pollution (green battery) wide working temperature range (−25 to +45°C, extended to −40 to 70°C).

Lithium-ion batteries have now become the major power sources for mobile electronics such as cellular phones, portable computers, digital cameras/videos and portable multimedia players. As competition in the market increases, lithium-ion batteries are required to have low cost and high performance, properties which are known to be dependent on the materials in the batteries including cathodes, anodes and electrolytes. Recently, many methods for preparing materials for lithium-ion batteries have been widely explored, for example incorporation of heteroatoms (Wu et al., 2002a), composite technology (Ning et al., 2004), soft chemistry routes (Manthiram and Kim, 1999) and some non-classical methods such as template methods, pulsed laser deposition, plasma-enhanced chemical vapour deposition, radio-frequency magnetron sputtering (Ning et al., 2004) and sol–gel methods (Fu et al., 2005). Several mechanochemical (MC) methods have also been tried. A MC method is one that can achieve its goal via chemical reactions in the solid state by various mechanical methods, mainly including grinding and milling such as ball milling, colloidal milling and jet milling. During milling or grinding, the mechanical energy is transferred to the particles of the reactants, causing many changes to the particles such as deformation, friction, fracture, amorphization, quenching and so on. After grinding or milling, a product phase(s) at the interfaces of the reactant particles is (are) formed. The main advantage of the MC method is its simplicity in the synthesis process, effective mixing accompanying a breaking of chemical bonds and their recombination, and a decrease in energy expenses as well as the cost of the materials (Ning et al., 2004). As a result, MC methods show some promise in the preparation of materials with low cost and good performance in lithium-ion batteries.

15.2

Lithium-ion batteries

15.2.1 The principle of lithium-ion batteries Originally, rechargeable batteries were found to function based on pure redox reactions. Later, rechargeable batteries were based on doping and undoping performance with a gain and loss of electrons (Novak et al., 1997, Wu et al., 2002b). Lithium-ion batteries are based on intercalation and deintercalation of lithium ions accompanying gain and loss of electrons or redox behaviour with the reaction shown in Equations [15.1]–[15.3], which is schematically shown in Fig. 15.1 (Wu et al., 2004).

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Li+/e Discharge Li+/e Charge

: Li

: C

: O

: Co

: Li

15.1 Principle of lithium-ion batteries. Adapted from Wu et al. (2002b).

During the charge process, lithium ions deintercalate from the cathode which is oxidized after losing electrons. The lithium ions deintercalated from the cathode pass through the electrolyte and intercalate into the anode. In this process, the anode gains electrons and is reduced. Prior to intercalation of lithium ions into the anode, a solid electrolyte interface (SEI) should be formed via the reduction of the electrolytes so that the anode will be stable upon contact with organic electrolytes. During the discharge process, the reverse happens. Lithium ions deintercalate from the anode which is oxidized after losing electrons. The deintercalated lithium ions from the anode pass through the electrolyte and then intercalate into the cathode. In this process, the cathode gains electrons and is reduced. Since the intercalation compounds for the anode are much less reactive than lithium metal, lithium-ion batteries present stable cycling performance and satisfactory safety. The reactions of the electrodes are: Discharge Cathode: LinMyXz X Lin−xMyXz + xLi+ + xe−

[15.1]

Charge Discharge +

Anode: A + xLi + xe− X LixA

[15.2]

Charge Discharge Total reaction: A + LinMyXz X Lin−xMyXz + LixA Charge

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[15.3]

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15.2.2 Requirements for materials In lithium-ion batteries, cathode, anode and electrolyte are three key components (Wu et al., 2004). As they perform various roles they must be manufactured from different materials in order to meet different requirements. In order to obtain a high output voltage, the intercalation compounds for the cathode should meet the following requirements: • • • • • • • •

to have a high redox potential for the intercalation of Li+ in the intercalation compound LixMyXz, providing a high output voltage; to have a large value of x for the reversible intercalation and deintercalation of Li+ in the intercalation compound Lin−xMyXz so that high reversible capacity can be achieved; to show little or no change in the host structure of [MyXz] during the repeated intercalation and deintercalation of Li+ so that the best cycling behaviour can be obtained; to show little or no change in the redox potential with varying x so that charge and discharge voltages remain stable; to have good electronic (σe) and ionic (σLi+) conductivities for LinMyXz so that the overpotential is small and the rate capability at large current density remains satisfactory; to have a high chemical stability for LinMyXz during the total charge and discharge process and no reaction between LinMyXz and the electrolyte; to have a large diffusion coefficient of Li+ in LinMyXz would be favourable for rapid charge and discharge; the manufacturing process for LinMyXz should be easily realized, at low cost and presenting no unfavourable environmental effects.

In the case of the anode materials, the requirements, almost similar to those for the cathode materials, are as follows: • • • • •

a low redox potential for the intercalation of Li+ in the intercalation compound LixA, providing a high output voltage; a large x value for the reversible intercalation and deintercalation of Li+ in the intercalation compound LixA so that high reversible capacity can be achieved; little or no change in the host structure of [A] during the repeated intercalation and deintercalation of Li+ so that optimum cycling behaviour can be obtained; little or no change in the redox potential with varying x, so that charge and discharge voltages will remain stable; good electronic (σe) and ionic (σLi+) conductivities for LixA so that the overpotential is small and the rate capability at large current density is satisfactory;

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a good surface structure so that an excellent solid–electrolyte interface (SEI) film can be formed during the initial charge process; a high chemical stability for LixA during the total charge and discharge process and no reaction between LixA and electrolyte after the formation of the SEI film; a large diffusion coefficient of Li+ in the host [A] so that the LixA complex can favour rapid charge and discharge; the manufacturing process for A should be easily realized, at low cost and with no unfavourable environmental effects.

As far as the electrolyte state is concerned, liquid electrolyte, gel electrolyte and solid electrolyte are possible. For application in lithium-ion batteries, the ideal solid electrolyte should have the following properties: • • • •

a high ionic conductivity (σLi+) especially at room temperature and low electronic conductivity (σe) to avoid or reduce self discharge; a stable phase structure even during the charge and discharge processes; recrystallization should be prevented especially in the case of glassy solid electrolytes; high chemical stability during the total charge and discharge process and no redox reaction upon contact with Li metal and electrode materials; good electrochemical stability and a wide electrochemical window above 4.2 V.

15.2.3 Materials for cathode, anode and electrolyte From the above requirements, it can be seen that powdered materials like layered LiCoO2 and LiNiO2, spinel LiMn2O4 and olivine LiFePO4 can be good cathode materials. Anode materials include carbon materials, alloybased materials including silicon, tin and antimony, and nitrides. There are mainly two kinds of solid electrolyte: solid polymer electrolyte and solid inorganic electrolyte (e.g. oxides- and sulfides-based electrolyte). Usually, most of the above-mentioned materials are prepared via high temperature heat treatment, which requires high energy consumption. Using MC methods, however, they can be prepared at room temperature, which makes them promising for practical applications.

15.3

Mechanochemical preparation of cathode materials

15.3.1 Lithium cobalt oxides (LiCoO2) As early as 1987, it was reported that a layered phase LiCoO2 was achieved using MC methods by ball milling mixtures of lithium and cobalt hydroxides

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(Fernandez-Rodriguez et al., 1987). However, since the mobility of lithium is increased by mechanical activation, spinel Co3O4, which is not favourable, will appear after an extension of the milling time. During the ball milling process, Co(OH)2 decomposes at first into CoOOH and then reacts with LiOH to form LiCoO2 (Jeong and Lee, 2001). Consequently, CoOOH and LiOH can be used to prepare LiCoO2 with a cubic spinel-related structure by ball milling for a shorter time (Kosova et al., 2002). However, the LiCoO2 prepared by the mechanical activation is characterized by high dispersion, structural disorder and different electronic states of the transition metal ions. MC treatment can be used to decrease the heat treatment temperature required to achieve the desired crystal structure and to shorten the heat treatment time (Jeong and Lee, 2002). If the ratio of ball to powder weight is increased, the MC reaction will be promoted owing to the increased number of collisions and greater collision energy per particle. In order to improve the safety and electrochemical performance of LiCoO2, an inert layer is coated on the surface of LiCoO2 particles so as to decrease their contact with the electrolyte. As a result, dissolution of Co4+ will be prevented (Li et al., 2006a). The most simple, economical and convenient way to coat LiCoO2 particles is by mechanical milling. In the case of 1.0 wt% silica-coated LiCoO2, the silica species can partially diffuse into the bulk and form a solid solution LiSiyCo1−yO2+0.5y on the surface, which is evident from the diminished lattice parameter c upon coating (Fey et al., 2004). In addition, other kinds of nano oxides such as Al2O3, TiO2, Li4Ti5O12, ZrO2, or their mixed oxide ZrTiO4 can also be coated on the surface of LiCoO2 by mechanical milling (Fey et al., 2005, 2006, 2007). For example, in the case of an Al2O3 coating derived from methoxyethoxy acetate-alumoxane (MEA-alumoxane) by a MC process followed by calcination at 723 K in air for 10 hours, a thin layer forms on the surface of the core material with an average thickness of 20 nm, which can prevent direct contact of LiCoO2 with electrolytes (Fey et al., 2006). This process is summarized in Table 15.1 and it is clear that MC methods provide a good method of preparing good quality LiCoO2 material with less energy consumption.

15.3.2 Spinel lithium manganese oxides (LiMn2O4) There are several kinds of lithium manganese oxides (Wu et al., 2004). Spinel LiMn2O4 is the most widely used in the application of lithium-ion batteries. MC methods can be used to synthesize highly dispersed LiMn2O4 spinel starting from different manganese oxides (MnO2, Mn2O3, MnO) and lithium compounds (Li2O, LiOH, LiOH⋅H2O, Li2CO3) (Ning et al.,

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Milling/40 hours

Ball milling + 600°C/2 hours Milling + 600°C/4 hours Milling Ball milling + 850°C/24 hours Ball milling

LiOH + Co(OH)2

LiOH⋅H2O + Co(OH)2

LiCoO2 + silica, TiO2 and ZrO2

HT-LiCoO2

Li2O + CoO LiOH⋅H2O + Co(OH)2

Milling

Milling/10 hours

LiOH + Co(OH)2

LiOH + Co(OH)2(CoOOH)

Process

Raw materials

LiCoO2 + disordered LixM1−xO Coated LiCoO2

Disordered, dispersed HT-LiCoO2 Spinel LiCoO2 HT-LiCoO2

–

HT-LiCoO2 + spinel Co3O4 LiCoO2

Excellent

Poor

Poor Good

Intermediate

–

–

Electrochemical performance

HT-LiCoO2

Structure

Table 15.1 Some LiCoO2 materials prepared by mechanochemical processes

Fey et al., 2004, 2005, 2006, 2007

Obrovac et al., 1998

You et al., 1998 Jeong and Lee, 2002

Kosova et al., 2002

Fernandez-Rodriguez et al., 1987 Fernandez-Rodriguez et al., 1987 Jeong and Lee, 2001

References

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2004). On the one hand, the oxidation state of manganese greatly influences the kinetics of MC reactions, for example, MnO2 reacts almost completely with Li2CO3 to produce LiMn2O4. However, there is no evident interaction between Mn2O3 or MnO with Li2CO3. On the other hand, different crystal structures and mechanical properties of the initial lithium compounds result in different mechanisms of MC action on the activated mixtures. LiOH has a layered structure and exhibits good plasticity; the chemical interaction during MC processing between MnO2 and LiOH is preceded by a stage of molecular-dense aggregates (mechanocomposites) formation through the action of adhesive forces and thus the surface of the mechanocomposite particle is covered by an amorphous LiOH layer. In the case of Li2CO3, which is a typical ionic compound but more brittle, the chemical interaction is preceded by a process of brittle fracture of the components of MnO2 and Li2CO3 (Kosova et al., 2000). A longer milling time results in amorphization and further decomposition of the spinel LiMn2O4 into Mn2O3 and Li–Mn–O (Soiron et al., 2001). The spinel LiMn2O4 resulting from ball milling the mixture of Li2O and MnO2 is highly disordered, with nanocrystalline sizes of less than 25 nm, and has much strain variance or defects (Choi et al., 2001; Soiron et al., 2001). The existence of the highly disordered structure could accommodate the Jahn–Teller distortion, which comes from the transition of Mn4+ into Mn3+ in the spinel structure during Li+ intercalation around the 3 V region (Choi et al., 2001). In addition, Li-doped LixMn2O4 (x > 1) spinels can also be prepared by MC activation since doping lithium can further improve the cycling performance of this cathode material (Kosova et al., 2000). The composition and lattice constants of the final products are affected by the lithium content. The intergrain resistance rather than the bulk properties of the spinels, including the starting reagents and molding pressure, determines their conductivity. The activation energy of conductivity (Ea) of LixMn2O4 does not depend on x over a wide composition range (0.21 ≤ x ≤ 1.21) and is 0.36 ± 0.04 eV (Kosova et al., 2000). Instead, Li doping increases the valence of Mn above 3.5, which will be good way of inhibiting the Jahn–Teller distortion. When micrometre LiMn2O4 powder is ball milled, the particles are broken into nano particles. On further milling, the nano particles stick back together again as hard agglomerates and many nanocrystallites (20–40 nm) are generated within a large crystallite by the action of defects such as dislocation and strain at the grain boundaries. Partial oxidation of manganese ions also occurs during ball milling for less than 120 min (Kang et al., 2001). Of course, the Jahn–Teller effect associated with the spinel structure still exists during lithium deintercalation and intercalation.

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Table 15.2 Some LiMn2O4 materials prepared by mechanochemical processes

Raw materials

Process

LiOH (LiOH⋅H2O, Li2CO3) + MnO2 LiOH + Mn2O3 (MnO) Li2O + MnO2

Milling

LiMn2O4 from solid state reaction LiMn2O4 + carbon

Ball milling

Milling Grinding

Ball milling

Features of structure

Electrochemical performance

Nanocrystalline LiMn2O4 Disordered LiMn2O4 Disordered nano spinel LiMn2O4 with strain or defects Nanodomains, strains, defects Nanocomposite of LiMn2O4 with carbon

4 V, good – 3 V, 4 V plateaus, good cycling

References Kosova et al., 2000 Kosova et al., 2000 Choi et al., 2001; Soiron et al., 2001

3 V, good cycling

Kang et al., 2001

High reversible capacity of 200 mAh g−1, excellent cycling

Im and Manthiram, 2003

However, the net deformation of a particle consisting of small grains is less anisotropic than a particle with large grains and the possibility of particle fracture caused by the tetragonal distortion will be lower. In addition, the strain from the tetragonal distortion is expected to be accommodated by the already existing strain in the particles. Some selected results are summarized in Table 15.2 and it is clear that MC methods provide an easy way of preparing nanocrystalline LiMn2O4 with good accommodation of strains.

15.3.3 Other kinds of cathode materials Reports on the preparation of LiNiO2 using MC methods are rare (Obrovac et al., 1998). However, the mixing of primary materials such as Ni(OH)2 and Li2CO3 usually utilizes MC methods like ball milling (Wu et al., 2002a). The surface stability with electrolytes can be improved by coating, for example, LiNi0.8Co0.2O2 can be coated with ZrO2 by simple grinding (Lee et al., 2006). Layered LiNi1/3Co1/3Mn1/3O2 presents better safety performance and higher reversible capacity than LiCoO2, and has become another popular cathode material after LiCoO2. It can be prepared by ball milling, which favours solid state metathesis reactions. For example, in the case of the mixture of LiAc⋅2H2O, Ni(Ac)2⋅4H2O, Co(Ac)2⋅4H2O, Mn(Ac)2⋅4H2O (Ac = acetate) and excess H2C2O4⋅2H2O without any solvent, after ball-milling and a fol-

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lowing heat treatment, layered LiNi1/3Co1/3Mn1/3O2 is prepared (He et al., 2007a). Of course, F can be doped by adding LiF during this process to obtain layered LiNi1/3Co1/3Mn1/3O2−zFz (0 ≤ z ≤ 0.12) cathode material. The primary particle size of LiNi1/3Co1/3Mn1/3O2−zFz gradually increases with the fluorine content (He et al., 2007b) so that a better cycling performance is expected. Olivine LiFePO4 has an advantage over other cathode materials since sources of iron raw materials in the Earth are very rich. Its preparation requires heat treatment in an inert atmosphere to ensure that the iron exists in the 2+ form. Its preparation by ball milling alone has not been reported since the primary material usually exists in the ferric (+3) state. Consequently, ball milling of FePO4 and Li3PO4 is combined with heat treatment, for example microwave heating. In addition, during the preparation process, the addition of conductive carbon as a coating after the heat treatment is necessary since the electronic conductivity of LiFePO4 is very low. The existence of Fe2P after heat treatment enhances the electronic conductivity of the LiFePO4/C composite. Of course, when the amount of Fe2P is too high, it blocks the one-dimensional pathways for Li+ in LiFePO4 and might hinder Li+ movement in LiFePO4 (Song et al., 2008). Other cathode materials, such as LiMnO2, can also be prepared by ball milling a mixture of transition metal and lithium oxides; short-range ordering of cations such as Li+ and Mn3+ in the crystal structure still occurs even after extended milling (Obrovac et al., 1998). When lithium manganese oxide (LixMnOz) is ball milled with conductive carbon, nanocomposite powder particles in which the components are intimately mixed on a nanometer level are prepared. Although nanostructured oxide particles may be intrinsically prone to poor electronic conductivity arising from high grain boundary areas and side surface reactions arising from the high surface area, results show that ball milling with carbon is an effective and simple way to avoid these problems owing to the presence of conductive carbon (Im and Manthiram, 2003). Layered LiNi0.5Mn0.5O2, another promising cathode material, can also be produced by ball milling LiOH, NiO and MnO2 (Xia et al., 2008a). LiMnPO4 and Li3Fe2−2xTixMnx(PO4)3 (NASICON-type structure) can be separately treated by ball milling to reduce the particle size (Drezen et al., 2007; Ning et al., 2004; Sun et al., 2009). For example, as shown in Fig. 15.2, the particle size of Li3Fe2−2xTixMnx(PO4)3 can be reduced from 10–20 μm to 100–500 nm by ball milling at a rotation speed of 220 rpm for 10 hours (Sun et al., 2009). In addition, conductive carbon can be further introduced by ball milling with acetylene carbon at a rotation speed of 220 rpm for 10 hours to increase the electrical conductivity of the powders (Sun et al., 2009).

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(b)

10 μm

100 nm

15.2 Scanning electron micrographs of (a) Li3Fe1.8Ti0.1Mn0.1(PO4)3 and (b) ball milled Li3Fe1.8Ti0.1Mn0.1(PO4)3/carbon mixture (reprinted from Sun et al., 2009 with permission from Elsevier BV).

15.3.5 Electrochemical performance of cathode materials produced using mechanochemical methods As mentioned above, after ball milling there are distortions and defects in the crystal lattice, so that the prepared LiCoO2 powder does not give a good electrochemical performance. Only in combination with subsequent firing or heat treatment at 850°C, will a well-ordered LiCoO2 achieving good electrochemical performance be obtained. When LiCoO2 particles are coated, the reactivity of Co4+ with liquid electrolytes is reduced. As a result, the cycling behaviour is greatly improved (Li et al., 2006a). For example, the cycling life of LiCoO2 with a coating level of 1.0 wt% Al2O3 was improved three and nine times, respectively, for two different commercial LiCoO2 samples (Fey et al., 2004). Impedance spectra of the coated cathode powder suggest that the surface coating on LiCoO2 reduces changes in the surface film during electrochemical cycling, resulting in a longer cycle life. The results are complemented by slow scan cyclic voltammetric profiles, which show that the alumina nano coatings lead to a suppression of the cycle-limiting phase transitions from hexagonal to monoclinic to hexagonal (Fey et al., 2006, 2007), which is due to slower impedance growth and increased resistance to Co dissolution into the electrolyte during the (de)intercalation processes. The above-mentioned coating technology is also effective for LiNiO2. For example, ZrO2 coating on the LiNi0.8Co0.2O2 cathode improves its cycling stability considerably owing to the suppression of the impedance growth during charge–discharge cycling (Lee et al., 2006). Milling time influences the charge and discharge curves of LiMn2O4 since mechanical energy input during milling is comparable to treatment at high temperature, which will influence its crystal structure. After an intermediate milling time of about 8 hours, the resultant LiMn2O4 spinel powder shows the charge and discharge curves typical of spinel LiMn2O4 prepared by solid

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state reaction at 800°C for 1 hour, that is plateaus at 3 V and 4 V (Soiron et al., 2001). Milling can produce highly disordered structures, which can alleviate the Jahn–Teller distortion around the 3 V region that always exists during lithium deintercalation and intercalation. The intercalation of Li+ takes place with an initial capacity of up to 167 mAh g−1 in the 2.5–4.3 V range with a steady slope instead of two plateaus at around 3 V and 4 V in spinel structure. In addition, the cycling behaviour is much improved compared to the well-ordered crystalline LiMn2O4 powders (Kang et al., 2001; Choi et al., 2001). As shown in Fig. 15.3, in a 50-cycle test at 0.5 mA cm−2, a sample ball milled for 1 hour gives a constant capacity of 122 mAh g−1 between 2.4–3.4 V versus lithium and the cycling behaviour at elevated temperatures is also satisfactory. In addition, polarization during cycling is also greatly decreased (Choi et al., 2001). In the case of nano LiMn2O4 powder obtained from ball milling, excellent capacity retention in the 3 V range at room temperature has been achieved. The main reason for this capacity retention is that the net deformation of a particle with small crystallites is less anisotropic than deformation with the large ones and there is less possibility of fracture of particles caused by the tetragonal distortion. In addition, the strain imposed by the formation of the tetragonal phase can be accommodated by the already existing strain in particles. Incidentally, ball milling provides intimate mixing that improves the electrical contact between spinel particles and carbon (Kang et al., 2001).

(b)

180

Capacity (mAh g–1)

Capacity (mAh g–1)

(a)

150 120 90 Cut-off voltage: 2.4–3.4V Current density: 0.5 mA cm–2

60 0

10

20 30 40 Cycle number

175 150 125 100

Room temperature 60°C 80°C

75 50

50

0

10

20 30 40 Cycle number

50

15.3 (a) Effects of ball-milling time on cycling behaviour of LiMn2O4; 䊊, n = 5 h; 䊉, n = 15 h; ⵧ, n = 30 h; 䊏, n = 60 h; 䉭, n = 120 h. (b) Cycling behaviour of one LiMn2O4 sample after 1 h milling at different temperatures; 䊉: room temperature; 䉭, 60°C; 䊊, 80°C – both with a voltage range of 2.4–3.4 V and a current density of 0.5 mA cm−2 (reprinted from Kang et al., 2001 with permission from the American Chemical Society).

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Good electrochemical performance is provided by LiMnO2 obtained by ball milling (Obrovac et al., 1998). The nanocomposite particles obtained by ball milling of nanostructured lithium manganese oxide (LixMnOz) with acetylene black exhibit high capacity of over 200 mAh g−1 with excellent cyclability and charge efficiency, which are attributed to increased electronic conductivity and decreased surface area (Im and Manthiram, 2003). Layered LiNi0.5Mn0.5O2 obtained from ball milling can present a reversible capacity of 120 mAh g−1 over the range of 2.5–4.3 V (versus Li+/Li) (Xia et al., 2008a). The electrochemical performance of LiNi1/3Co1/3Mn1/3O2 obtained by ball milling is excellent. A steady discharge capacity of about 140 mAh g−1 after 100 cycles at a rate of 1 C (160 mA g−1) in the voltage range 3–4.5 V (versus Li+/Li) is achieved and the capacity retention is about 87% at the 350th cycle (He et al., 2007a). When F is doped in the ball milling process, the fluorine-substituted LiNi1/3Co1/3Mn1/3O2−zFz (0 ≤ z ≤ 0.12) presents a lower initial discharge capacity. In the case of a small amount of fluorinesubstituted LiNi1/3Co1/3Mn1/3O2−zFz (z = 0.04 and 0.08), this material exhibits excellent cycling and better rate capability compared with fluorine-free LiNi1/3Co1/3Mn1/3O2 (He et al., 2007b). LiFePO4/C composite with Fe2P below a critical concentration provides a very high discharge capacity of 165 mAh g−1, excellent rate capability (85.4% of C/50-rate discharge capacity at 2C) and stable capacity retention for 250 cycles. Above the critical concentration, the electrochemical performance deteriorates. In order to obtain a LiFePO4/C composite with good electrochemical performance, the amount of Fe2P should be carefully controlled (Song et al., 2008). After ball milling the initial LiMnPO4 powder, reversible capacities of 156 mAh g−1 at C/100 and 134 mAh g−1 at C/10 have been obtained, which are 92% and 79% of the theoretical values, respectively. At higher charging rates, the electrochemical performance is also improved since the particle size of LiMnPO4 is reduced compared to that without ball-milling (Drezen et al., 2007). Improved cycling behaviour is achieved if LiMnPO4 powder particles are coated by carbon via the milling process (Sun et al., 2009). In summary, compared with conventional solid state reactions at high temperature, MC methods appear to accelerate and simplify the synthesis process and decrease the energy expenses as well as the cost of the final product. Consequently, based on the latest developments in the technology to prepare cathode materials, further improvements can be expected. Obviously, when combined with some other methods such as high temperature heat treatment and adding a solvent such as water during milling (Kakuda et al., 2007), different cathode materials with even better electrochemical performance will become available (Ning et al., 2004).

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15.4

Mechanochemical preparation of anode materials

Various kinds of anode materials such as carbonaceous and alloy-based powders have become available (Wu et al., 2002a). Some have been prepared or modified by MC methods and these present some interesting changes and properties.

15.4.1 Carbonaceous materials Carbonaceous materials were the first available and major anode materials to come on the market. Since the highest reversible capacity of graphite is 372 mAh g−1, many methods have been tried to break through this limit, for example doping with heteroatoms (Wu et al., 2002a). Recently, it was found that milling or grinding was a good way to modify the electrochemical performance of carbon materials. The structure changes of carbonaceous materials are dependent on the ball milling period and mode. These changes mainly include the following aspects as shown in Table 15.3: particle size, surface area, surface structure, electronic self-spinning of radicals, microstructure such as defects and nanocavities or voids – and crystal structure. The effects of milling and grinding mainly on the structure and electrochemical performance of graphite are shown in Table 15.3. During milling, the graphite particle size will change. In general, bondbreaking of graphene layers leads to smaller particles (Disma et al., 1996). For example, after 150 hours ball milling, well-graphitized graphite is pulverized into small particles with a size of about 50 nm. Owing to the large surface energy, the merging of single particles is favoured and results in the formation of agglomerates with an average size of about 1 μm. In the case of impact or shock type ball milling, interactions vertical to the graphite cleavage planes are so strong that graphite particles are torn into pieces of disordered carbons. Although jet milling and turbo milling are soft and mild techniques, they also cut graphite particles into smaller ones in directions both perpendicular and parallel to the basal planes (Wang et al., 1999). Since the graphite particle size becomes smaller during milling, the specific surface area of the material increases with milling, although the degree depends on the milling mode (Disma et al., 1996; Natarajan et al., 2001). The surface structure especially the composition of surface groups is greatly changed. Different ambient atmospheres, for example air, high purity argon, nitrogen, and CO2 produce different effects. The active surface sites formed by milling react readily with active gases or vapours, for example O2, CO2 and H2O, to produce oxygen-containing highly reactive surface groups (Aurbach et al., 2002). Reactive milling can be used to develop bonding between carbon and lithium atoms and to form a lithiated surface which decreases the irreversible capacity of carbon in the first cycle.

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Torn into pieces of disordered carbon

Impact/shock-type mechanical milling/ grinding

Cut into smaller particles

Particle

Nano agglomerates

Drastic, strong

Soft, mild

Type of milling action

Ball milling

Jet- and turbo-milling

Mechanochemical methods

Increase

Increase

Increase

Specific surface area

O-terminated carbon

Surface structure

Localized spins, smaller g-value

Localized spins, larger g-value

Radicals

Large ratio of disorder such as vacancies, microcavities, and voids

Microstructure

Changes in structure parameters

Increase or little change in d002, and decrease in content of hexagonal phase Decrease in La and Lc

Little change in d002, decrease in Lc

Crystal structure

Increased reversible capacity mainly associated with the charge slope >1.0 V, increase in irreversible capacity, capacity fading

Decrease in irreversible capacity, increase in coulomb efficiency under optimal conditions Increase in reversible capacity and coulomb efficiency, improved cycling

Electrochemical performance

Disma et al., 1996, Wang et al., 1999, SalverDisma et al., 1997, 1999

Natarajan et al., 2001

Wang et al., 1999, SalverDisma et al., 1997

References

Table 15.3 Selected results of the effects of mechanochemical reactions on structure and electrochemical performance of graphite

© Woodhead Publishing Limited, 2010

Decrease in particle size

Increase

Oxygencontaining groups

Slight increase

O2, CO2, H2O

Reactive atmosphere

Jet milling

Lithiated surface

Increase

Graphite + lithium

Reactive milling

O-terminated carbon

No evident change

Shear force

Surface structure

Shear-type grinding

Particle

Specific surface area

Type of milling action Radicals

Disorder present

Less disorder or disorganization such as vacancies, microcavities

Microstructure

Changes in structure parameters

Mechanochemical methods

Table 15.3 Continued

Increase in rhombohedral (3R) phase

Little change in d002 and crystal size

Little change in crystal size and d002

Crystal structure

Spontaneous formation of surface passivating film, high reversible capacity, lower hysteresis Loose surface passivating film, exfoliation of graphite, capacity fading, slight increase of reversible capacity High initial capacity

Increase in reversible and irreversible capacity

Electrochemical performance

Herstedt et al., 2003

Aurbach et al., 2002

Tossici et al., 2003

Salver-Disma et al., 1997, 1999

References

Materials for lithium-ion batteries by mechanochemical methods

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ESR spectra of natural graphite derived powder vary with grinding methods. The line width which is due to conduction electron spins broadens for all ball milled specimens except for those ground by a ball mill down to 1 μm diameter. The contribution of localized spins produced in the mechanical grinding process becomes predominant over that of conduction electron spins in the range of T < 50 K (Salver-Disma et al., 1997). On the other hand, the absorption intensity of ball milled material satisfies the Curie law. Samples ground by jet milling have larger g-values compared with those prepared by ball and colloidal milling (Ning et al., 2004). However, for each grinding mode, the change is totally independent of the nature of the precursors used (graphite, mesocarbon microbead, coke) and/or its morphology (layers, microbeads or fibres) (Salver-Disma et al., 1997). The microstructure is also affected by the MC treatment. For example, shock-type mechanical milling generates a large ratio of disorder within the powder and small stacks of 2–3 parallel fringes of less than 1 nm are formed. These disordered domains are misoriented and distributed at random to form mesopores and no single layer is observed. Shear-type grinding generates a weaker mechanical strain than the shock mode and the carbon material is much less damaged, with fewer defects in the directions both perpendicular and parallel to the basal planes (Wang et al., 1999). Depending on the type of grinding, energetic interactions of different intensity are generated resulting in disorder or disorganization of the carbonaceous materials (Disma et al., 1996; Natarajan et al., 2001; Salver-Disma et al., 1999), observed by transmission electron microscopy (Salver-Disma et al., 1999). In addition, many vacancies, microcavities (or voids) and metastable carbon interstitial phases with sizes around 1.3 nm are formed among the agglomerated particles. Different atmospheric conditions also result in different microstructures. For example, in the case of a planetary mill where the deformation forces are mainly shear in nature, mechanically induced oxidation on the surface, probably along the edges of the graphene planes, suppresses the fracture rate and preserves the crystallinity of natural graphite milled in oxygen (Ning et al., 2004). Generally, after milling, the crystal structure of carbonaceous materials will be changed in the same way as it is after MC treatment and this is mainly reflected in disordering of the structure, changes in the interlayer distance d002, and the content of crystal phases (Wu et al., 2002a; Disma et al., 1996; Wang et al., 1999; Natarajan et al., 2001). The interlayer distance d002 usually increases with milling time owing to the introduction of voids and vacancies. However, in the case of mild milling such as turbo and jet milling, little change in d002 of the milled carbons is observed owing to the less destructive effect on graphite particles (Wang et al., 1999). The graphite lattices become thinner owing to expansion of graphite layer interspacing (Disma et al., 1996). As a result, the content of the hexagonal phase

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decreases and that of the rhombohedral phase (3R) increases (Wu et al., 2002a). However, there is a conflicting report that the rhombohedral phase fraction decreases with the milling time and then stabilizes at about 10% (Natarajan et al., 2001). Perhaps this depends on the species of graphite used. Similarly, in the case of other kinds of carbons such as non-graphitic carbon, the milling affects the particle size distribution, BET specific surface area, interlayer distance and other structural characteristics (Ning et al., 2004).

15.4.2 Alloy-based anode materials The first alloy-based powders for anode materials were produced by ball milling in the late 1970s. However, the problem associated with the growth of dendrite, which can easily lead to short circuit of the batteries, has not been completely overcome. As a result, alloy-based materials gave way to graphitic carbons in the early 1990s (Wu et al., 2002b). Recently, however, it was found that introduction of a conductive or inactive agent as a buffer could lead to an improvement in electrochemical performance of alloybased anodes. Consequently, the focus has again moved to novel alloybased anode materials, which mainly include tin-based, silicon-based and antimony-based alloy anode materials. Some selected results are summarized in Table 15.4. In the case of pure tin, since it can form alloys with Li up to Li22Sn4, the volume change can be up to 376%, resulting in serious cracking or crumbling of Li–Sn alloys and rapid capacity fading (Wu et al., 2002b). As a result, some buffering components, such as Li2O, conductive metals or graphite should be added to prevent or buffer the drastic volume change. Composites consisting of Sn and Li2O clusters can be prepared directly by ball milling tin and Li2O (Ning et al., 2004) or SnO and lithium metal as precursors (Wang et al., 2008b). In these clusters, tin particles are uniformly distributed within a lithium oxide matrix with the majority size of the tin particles in the order of 100 nm or less, which can be seen clearly from the transmission electron micrograph shown in Fig. 15.4. Besides Li2O, Al2O3 can also be used as a matrix. Finely dispersed tin alloy/oxide composites have been synthesized via the reduction of SnO by aluminium employing high-energy ball milling (Al powder and Co or Ni powder as precursors in a molar ratio of SnO :Al : Co = 1 : 0.878 : 0.333 or SnO :Al : Ni = 1 : 0.878 : 0.75). During milling process, the following reactions take place: 3SnO + 2Al → 3Sn + Al2O3

[15.4]

3Sn + Co → Sn3Co

[15.5]

4Sn + 3Ni → Sn4Ni3

[15.6]

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Ni3Sn4 Composite such as C0.9Sn0.1 or C0.8Sn0.2 SnMn3C

Sn31Co28C41 Sn3Co/Al2O3 Ag4Sn coated graphite Si Si/Li4SiO4 + Li2O

Mixture of SiO and graphite Si + Si3N4

Si core + amorphous carbon

Sn + Ni Sn + graphite

Sn + Co + C SnO + Co + Al Sn + Ag + graphite

SiO + graphite

Si + C

Si + Si3N4

SiBr4 + Mg SiO + Li

Sn + Mn + C

Cu6Sn5 Mg2Sn

Sn/Li2O

SnO + Li

Sn + Cu Sn + Mg

Composition

Precursors

Nano sized Monoclinic Li4SiO4 + nano Si Decrease in crystal size of graphite Nano crystal or amorphous Si Core-shell

Crystal Sn + amorphous Li2O Hexagonal structure Cubic + orthorhombic phase Nanocrystalline Amorphous graphite + nanocrystalline Sn Perovskite, nano particles Crystal Nano crystal of Sn3Co Amorphous graphite

Structure

Good

Good

>320 1000

Good

Fair to good Good

Good Good Fair to good

Good

Good Good

Improved but fair Good Good

Cycling

696

1375 770

500 540 >600

150

125–200 400–600

200 460

>500

Reversible capacity (mAh g−1)

Electrochemical performance

Table 15.4 Selected results of alloy-based anode materials prepared by mechanochemical methods

Liu et al., 2004b

Zhang et al., 2007b

Doh et al., 2008

Sandu et al., 2007 Yang et al., 2007

Hassoun et al., 2007 Zheng et al., 2007a Wang et al., 2008a

Beaulieu et al., 2000

Lee et al., 2002 Wang et al., 2001

Xia et al., 2001 Kim et al., 2001

Wang et al., 2008a

References

© Woodhead Publishing Limited, 2010

Co + Sb Zn4Sb3 + graphite

β-Zn4Sb3

Si + Sn + graphite

Fe80Si20 + graphite Si + Ni + graphite

Si + graphite

Si + copper(II) D-gluconate Si + Ni (Fe)

Si + TiN

85C–15[Si0.66Sn0.34] composite ZnSb + unknown structure CoSb3 Composite

Composite such as C0.8Si0.2 Composite Si/C alloy

NiSi (FeSi)

Si + Cu3Si + Carbon

Composite of Si and PPy

SiAg

Si + Ag

Si + PPy

Composition

Precursors

Table 15.4 Continued

Fine powder, <100 nm

Si + graphite with NiSi and NiSi2

Nanocrystalline or amorphous Si dispersed in TiN Carbon-coated composite Mixture of NiSi (FeSi) and Si

Nanosized Si in Ag matrix PPy coated Si

Structure

586 581

Poor Good

Poor

Good

>500 Increased to 566

Good Excellent

Satisfactory

Poor

Good

Good

Fair

Good

Cycling

ca. 600 780

1039

600–1000

1000

ca. 300

>1000

ca. 280

Reversible capacity (mAh g−1)

Electrochemical performance

Zhang et al., 2002 Zhao et al., 2001

Rock and Kumta, 2007 Cao et al., 2001

Lee and Lee, 2002 Park et al., 2005

Wang et al., 1998

Wang et al., 2000

Yoon et al., 2006

Kim et al., 2000

Guo et al., 2005a

Hwang et al., 2001

References

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0.2 μm

15.4 Transmission electron micrograph of Sn and Li2O composite particles (reprinted from Wang et al., 2008b with permission from Elsevier BV).

As a result, Sn3Co/Al2O3 and Sn4Ni3/Al2O3 composite powders were formed (Zheng et al., 2007a). Cu–Sn is the first reported novel alloy anode material (Wu et al., 2002b). It was found that Cu–Sn alloys could be prepared by MC methods with flake sizes below 1 μm (Xia et al., 2001). Mechanochemical reaction of Mg and Sn powders produces another tinbased alloy, Mg2Sn. The structure obtained changes with milling time and a mixture of cubic and orthorhombic phases of Mg2Sn is achieved, which provides some interstitial space that can accommodate possible stress during electrochemical reactions. In the case of cubic Mg2Sn, although some intercalation sites do exist, they appear to be too narrow. However, the orthorhombic phase has a layered structure and the distance between the layers is sufficient to accommodate Li atoms, so that lithium can insert into the layers (Ning et al., 2004). Apart from Cu and Mg, Ni can also form alloys with Sn and a nanocrystalline Ni3Sn4 alloy can be prepared by high energy ball milling (Lee et al., 2002). In addition to conductive metals such as Mg and Ni used as buffer or additive, graphite, which is an active material for lithium-ion insertion, can also be mixed with Sn. After intensive ball milling, composite particles of graphite and tin are obtained, with Sn encapsulated in the ductile graphite matrix at the nanometre scale, which can be seen in Fig. 15.5. During milling, graphite becomes amorphous and tin becomes nanocrystalline. In contrast with the ball milling of graphite alone, the graphite in this composite con-

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C

Sn

5 nm

15.5 High-resolution electron micrograph of ball milled C0.8Sn0.2 composite (reprinted from Wang et al., 2001 with permission from Elsevier BV).

tains fewer interstitial carbon atoms owing to the existence of the soft Sn particles. Tin particles are broken down to 15–20 nm and still attain an ordered crystalline state. Also in this material there are some interstitial sites such as vacancies and disorganized regions (Wang et al., 2001). On the basis of the composite of Sn with graphite, further additives such as metallic Mn, Fe, Co and Ag can be introduced into Sn particles. In the case of Mn, an intermetallic compound, SnMn3C, which has a perovskite structure, is obtained. After ball milling, the resultant SnMn3C powder comprises nanostructured particles with different crystallographic orientations and grain boundaries (Beaulieu et al., 2000). As for SnFe3C, its structure is similar to SnMn3C after ball milling. The Sn–Co–C composites are prepared by ball milling commercially available micrometresized powders of Sn, Co and C (graphite). The Sn : Co : C molar ratio can be varied, such as 0.31 : 0.28 : 0.41, 0.3 : 0.3 : 0.4 or 0.36 : 0.41 : 0.23. Composites of Sn30Co30C40 or Sn36Co41C23 with the same nanostructure from cosputtering can also be prepared by high energy ball milling and vertical-axis attritor milling (Hassoun et al., 2007; Ferguson et al., 2008). The tin–graphite–silver (Sn–G–Ag) composite powder obtained from high-energy mechanical milling comprises electrochemically active Sn and graphite, as well as Ag4Sn phases which are uniformly distributed on the surface of the graphite particles (Wang et al., 2008a). New tetragonal phases such as those of Ni, Sn and P can also be directly prepared via MC reactions (Xia et al., 2008b).

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Silicon, similar to tin, can form alloys up to Li22Si4 (Wu et al., 2002b). As a result, Si-based alloy anode materials are also promising owing to their high reversible capacity. However, the volume change during the charge/ discharge process is also very large, even more than that of Sn and capacity fades very quickly. Consequently, some measures are used to buffer or alleviate the drastic volume change. The first measure is to introduce inert oxides into Si in order to decrease volume expansion. Composite particles consisting of nanosized silicon, Li4SiO4 and other lithium-rich components are synthesized using reactive milling of SiO amorphous powder with lithium metal, resulting in the oxidation of lithium and silicon and the reduction of SiO to Si (Yang et al., 2007). Nanosized silicon particles have been obtained by reaction between SiBr4 solution and Mg granules during a ball milling process. After the MC reaction, crystallized silicon was obtained together with MgBr2, non-reacted Mg and probably amorphous SiOx. When the residual products were removed by cleaning processes, silicon powder consisting of nanosized Si particles with a limited oxidized fraction in regard to its high surface area was obtained (Sandu et al., 2007). Instead of Li and Mg, Al can be used as a reducing agent. By milling a mixture of amorphous SiO (325 mesh) and aluminium (100–200 mesh) powders with a molar ratio of SiO :Al = 1 : 1 at a speed of 400 rpm for 15 hours in an argon atmosphere, a Si/Al2O3 composite was prepared. The composite was added slowly to hydrochloric acid while stirring to remove Al2O3 and nano-porous Si powder particles were formed together with a thin layer of oxides. On the basis of the nanoporous Si particles, carbon can be further incorporated by using an approach such as ball milling to obtain nanoporous Si/graphite composite powder (Zheng et al., 2007b). Ballmilling of SiO can lead to a reduction in particle size (Doh et al., 2008). Of course, graphite or pyrolyzed carbon can also be added (Kim et al., 2007). The second way to reduce volume expansion is by adding some electrochemically inactive or inert compounds such as TiN, TiC or Si3N4. Nanocomposite particles of silicon and TiN from high-energy mechanical milling are composed of nanosized TiN containing a uniform dispersion of Si independent of the composition. The very small size of the Si crystals prevents the formation of a distinct phase boundary between Si and TiN. As a result, the volume changes continuously rather than abruptly and discretely (Kim et al., 2000). The nanocomposite particles of silicon/titanium carbide (TiC) obtained by high-energy ball milling consist of amorphous silicon and nanocrystalline titanium carbide (Guo et al., 2005b). In the case of the composites of silicon/Si3N4 nanoparticles, the amorphous Si3N4 particles could be the inactive buffer matrix supporting Si particles, which could prohibit further agglomeration of Si particles during cycling and ensure cycling stability (Zhang et al., 2007b).

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The third approach is to add some metals such as Mg, Ni, Fe, Mn, Ag, or alloy compounds (Wu et al., 2002a; Hwang et al., 2001) in order to decrease the volume expansion. In the case of Ni, Fe and Mn, they form nanosized intermetallic alloy powders with Si such as NiSi and FeSi, FeSi2 and MnSi after MC reactions, with Si still acting as an active material for lithium storage (Wang et al., 2000; Zuo et al., 2006). SiAg powder particles formed by the mechanical alloying process appear to contain a uniform dispersion of Si in the ductile Ag particle matrix. Ag not only provides a buffering environment for nanocrystalline Si but also serves as a conducting matrix. Of course, longer milling time leads to better dispersion of Si in the Ag matrix (Hwang et al., 2001). In the case of alloy compounds, a typical example is AB5 (e.g. MmNi3.6Co0.7Al0.3Mn0.4). In the Si–AB5 composites, Si nanocrystals are distributed homogeneously on the surface of the AB5 matrix and the inactive AB5 alloy can accommodate the large volume changes in the Si nanocrystals (Zhang et al., 2007a). The fourth technique is to add conductive agents such as carbon and polypyrrole. During the mechanical milling of Si and graphite powders, the crystal size of graphite increases and that of silicon decreases with the reducing silicon content (Wang et al., 1998). Ball milling of silicon powder with single wall carbon nanotubes (SWCNTs) functionalized by LiOH can provide good electrical contact between the electrochemically active particles (Wang and Kumta, 2007). In the composite powders of silicon with polypyrrole obtained by high-energy mechanical milling, polypyrrole acts as a good matrix to hold the active silicon grains and buffer the possible volume change, which is schematically shown in Fig. 15.6 (Guo et al., 2005a). Of course, the above approaches can be combined to obtain better composites of Si powder particles (Yan et al., 2008). For example, graphite can be ball milled with alloys of Si such as Fe20Si80, FeSi6 and MnSi to prepare composite powders (Lee and Lee, 2002; Li et al., 2008; Zuo et al., 2006). Electrochemically active Si0.66Sn0.34 (SiSn) composite alloys can be dispersed in a carbon (graphite) matrix using both wet and dry high-energy mechanical milling (HEMM). The resultant composite powder particles consist of amorphous carbon (in the case of dry HEMM) or crystalline carbon (in the case of wet HEMM) and crystalline silicon and tin (in both cases) (Rock and Kumta, 2007). Si–M–C (C is disordered carbon, and M is a hard comilling component like TiB2, TiN or Ni) composite powders can be prepared by the HEMM process. A large number of pores are produced, which can accommodate volume changes effectively and are essential to prevent the formation of inactive secondary phases between Si and M. In the case of Si/Ni alloy-C composites, inactive secondary phases such as NiSi and NiSi2 produced during the arc-melting process act as buffers to reduce pulverization of the

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PPy

Si

15.6 Schematic diagram of the Si/PPy composite particle: Si crystals act as the active phase and PPy as the inactive phase to buffer volume change (reprinted from Guo et al., 2005a with permission from Elsevier BV).

electrode (Park et al., 2006b). A Si–Zn–C composite powder is prepared by mechanical ball-milling using Zn as third element (Yoon et al., 2007). When the above processes are combined with heat treatment, new composite anode materials based on Si can be prepared. For example, a carbon–Si–Cu3Si composite powder was prepared using silicon and copper(ii) d-gluconate powders as precursors via a MC reaction and subsequent pyrolysis. In this process, Cu3Si and pyrolysed carbon uniformly adhere to the surface of the silicon particles (Yoon et al., 2006). Combination of HEMM with thermal pyrolysis may provide silicon-disordered carbon composite particles, in which multiphase Si–C cores are homogeneously distributed within the pyrolysed carbonaceous matrix (Liu et al., 2004c). In the Si/Ni alloy and graphite composites obtained by arc-melting followed by HEMM, alloy particles consisting of NiSi2, NiSi and Si phases are distributed finely and uniformly on the graphite surface. Fourier transform infrared spectroscopic (FTIR) analysis confirmed that some bonds were formed between alloy and graphite after HEMM, which appear to retain the electrical connection between alloy and graphite (Park et al., 2005). Si/G/PAN-C composite powders can be prepared by mechanically milling mixtures of elemental powders of synthetic graphite (1–2 μm), Si (−325 mesh) and PAN with a high energy shaker mill for up to 15 hours in a stainless steel (SS) vial using 20 SS balls of 2 mm diameter (total ball mass 20 g) with a ball to powder weight ratio 10 : 1 (Datta and Kumta, 2006).

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Since the chemical properties of Sb are similar to those of Si and Sn, and lithium insertion can reach the level of Li3Sb, Sb-based alloys can also potentially be used as the anode for lithium-ion batteries including β-Zn4Sb3, CoSb3 and InSb, when they are prepared using MC methods, as shown in Table 15.4. β-Zn4Sb3 has a rhombohedric crystal structure with a = 1.2231 nm and c = 1.2428 nm; its density (6.077 g cm−3) is lower than that of Zn (7.14 g cm−3) or Sb (6.684 g cm−3), suggesting that there is some space available in the structure for lithium storage. After high-energy ball milling, β-Zn4Sb3 alloy powders change into a ZnSb structure together with another unknown structure (Cao et al., 2001). CoSb3 powder with different particle sizes can also be prepared by ball milling Co and Sb. As can be expected, the particle size decreases with an increase in the ball milling time (Zhang et al., 2002). Heteroatoms such as Fe can be incorporated by high-energy ball milling to obtain ultrafine CoFe3Sb12 (Zhao et al., 2001). InSb can be prepared by ball milling In and Sb, the purpose of which is similar to that of the introduction of inert components in Si (Hewitt et al., 2001). Like Sn and Si, conductive agents acting as a buffer, and a conductive matrix such as graphite or carbon can also be introduced (Wang et al., 2001; Lee and Lee, 2002; Zhang et al., 2005a). For example, composites such as CoFe3Sb12–C16 (Zhao et al., 2001) and Zn4Sb3–C7 (Cao et al., 2001) have been successfully prepared. Of course, other metals may also be used to form alloys with Sb by MC reactions if they present good conductivity and good stability during cycling.

15.4.3 Anode materials from nitrides The application of nitrides as anode materials is due to the high ionic conductivity of Li3N. After the introduction of transition metals (TM) to partially substitute Li, Li–TM–N has shown promise as anode materials (Wu et al., 2002a). Partial results of mechanochemical reactions to prepare anode materials by using nitrides are summarized in Table 15.5. Li3N reacts with SnO during the milling process according to Equation [15.7] (Foster et al., 2000): 2Li3N + 3SnO → 3Li2O + 3Sn + N2↑

[15.7]

Ball milling SnO with Li3N produces composite particles of Li2O and Sn, wherein Sn is homogeneously distributed, with a particle size of about 100 nm or less. During the first charge–discharge cycle, the irreversible capacity is greatly decreased owing to oxidation of Li3N. Lithium-rich silicon nitride composites can be prepared by ball milling a mixture of Si powder (<100 nm) and Li3N powder in a mole ratio of 1 : 5. Nanometre precursor powder is recommended to accelerate the reaction and improve

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Table 15.5 Summary of the effects of ball milling on anode materials based on nitrides

Precursors

Compounds

Li3N + SnO

Sn + Li2O

Li3N + SiO2

Non-stoichiometric Li–Si–N Amorphous lithium silicon nitride composites Composite

Li3N + Si

Li–TM–N + carbon

Li2.6Co0.4N + Cu (Ni,Fe)

Li2.6Co0.4−xCux−yMyN (M = Ni and Fe)

Effects

Electrochemical performance Decrease in irreversible capacity and improved cycling behaviour Improved Improved

Ultrafine Li–TM–N powder is dispersed and adhered compactly on the surface of carbon Doping

References Foster et al., 2000

Wen et al., 2008 Wen et al., 2006

Increase in reversible capacity and initial coulomb efficiency

Hanai et al., 2008; Liu et al., 2004b

Slight decrease in reversible capacity and stable cycling behaviour

Liu et al., 2004a

the homogeneous property of the product. The structure of the composite particles is stable. Lithium silicon oxynitride composite can be prepared by high-energy ball milling of Li3N with SiO2 and the reactions are similar to that of the Equation [15.7] (Wen et al., 2008). When Li–TM–N is ball milled with carbonaceous materials such as disordered or graphitic carbon, an improvement in electrochemical performance in comparison to that without ball milling is observed (Hanai et al., 2008). Furthermore, Li–TM–N, like Li2.6Co0.4N, can be doped with Cu and Fe by HEMM (Liu et al., 2004b, c).

15.4.4 Other kinds of anode materials Other kinds of anode materials mainly include tin-based oxides since it is known that Sn can be an active host for reversible lithium storage. As mentioned above, in the case of SnO and SnO2, the volume change during insertion and deinsertion of lithium is also very large and the main method

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of absorbing this volume change is to introduce buffers such as non-metal and metal oxides. The system SnO–B2O3–P2O5 can be synthesized by mechanical treatment of starting oxides in a dry N2 atmosphere at room temperature (Wu et al., 2002a). Whether Li2O is present or not, the product is amorphous consisting of fine particles smaller than 1 μm in diameter. Metal oxides such as ZnO and MoO2 can be incorporated into SnO2 by ball milling (Wu et al., 2002a; Martos et al., 2002). For example, Mo-doped Sn1−xMoxO2 mixed oxides exhibit low crystallinity with a Li–Mo–O oxide conductive matrix. They are solid solutions with a cassiterite-type structure, with Mo in the tetravalent oxidation state (Martos et al., 2002). Nano SnS particles can be prepared by ball milling Sn (74 μm) and S (50 μm) powders in an atomic ratio of 1 : 1. When the ball milling time increases, the Bragg diffraction peaks of SnS become evident and show broadening and those of Sn and S gradually decrease (Li Y, et al., 2006b). Another active host for lithium is amorphous SiO powder since the Si can reversibly alloy and dealloy with Li. When amorphous SiO is mechanically milled with porous carbons such as carbon cryogel or carbon xerogel, some SiO is turned into Si and the particle sizes of silicon and carbon decrease with more uniformity, as shown in Fig. 15.7 (Yuan et al., 2007). Moreover, combined with the active SnO, an amorphous 50SiO⋅50SnO (mol%) composite powder can be obtained by mechanical treatment in a dry N2 atmosphere at room temperature (Morimoto et al., 2001). When α-FeOOH is mechanically milled, it turns into α-Fe2O3 (Morimoto et al., 2005). For example, after milling with a planetary ball mill at 250 rpm in the air for 100 hours, there is no clear evidence of α-FeOOH from X-ray diffraction; only crystalline α-Fe2O3 is obtained without heat treatment. This observation suggests that crystalline α-FeOOH is dehydrated during mechanical milling according to the following reaction: α-FeOOH → 1/2 Fe2O3 + 1/2H2O (a)

[15.8]

(b)

1 μm

1 μm

15.7 Scanning electron micrographs of (a) carbon xerogel (CX) and (b) CX–SiO from ball milling of CX and SiO for 10 hours at 400 rpm in a planetary ball milling machine (reprinted from Yuan et al., 2007 with permission from Elsevier BV).

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Beside the chemical change in the compound, the morphology of FeOOH particles also changes from needle-shape to fine aggregated particles. Phosphides MPx (M = V, Fe, Co, Ni and Cu) may also be used as anode materials for lithium-ion batteries; they can be prepared by mechanically milling red phosphorus and metals (Gillot et al., 2007). If Li3N is added during the milling process, a lithiated product, LiyMPx can be prepared. With increasing milling time, the phosphide CoP3 is converted to a CoP phase (Bichat et al., 2004; Zhang et al., 2005b). Even more anode materials such as FeS (Kim et al., 2005) and carboncoated Mg composite particles can also be obtained by MC reactions (Park et al., 2006a), which increase the electronic conduction and decrease the particle size.

15.4.5 Electrochemical performance of anode materials obtained by mechanochemical methods Carbonaceous anode materials As mentioned earlier, different changes in structures have been observed in accordance with different milling ways and conditions. As a result, the electrochemical performance of these anode materials also varies greatly and some selected results are shown in Table 15.3. After milling, the reversible capacity of graphite is increased up to 700 mAh g−1 (Salver-Disma et al., 1997, 1999). Of course, the increase is dependent on the type or mode and time of milling and is due to several factors: (i) the production of a large number of edge planes during milling, which are highly active and reactive with lithium during charge process, (ii) vacancies, microcavities or voids, and interstitial phases produced at the edges of the metastable carbon, where lithium can be absorbed (SalverDisma et al., 1999) and (iii) other disordered structures that may be present. As shown in Table 15.6, the particle size also exerts an influence on the initial charge capacity (Ning et al., 2004). Smaller particle sizes provide more inlets and outlets for lithium at the surrounding graphite crystallites and edge planes. As a result, the reversible capacity of lithium storage is increased, above the theoretical value of graphite. Since considerable defects or disorders are introduced into the milled graphite, a high reversible capacity is usually accompanied by a large irreversible capacity (Salver-Disma et al., 1997). Active structures such as the large number of edge planes, vacancies, microcavities or voids and interstitial phases at the edges of the metastable carbon (SalverDisma et al., 1999) are the main contribution to irreversible capacity. Of course, the increased surface area also contributes to irreversible capacity since the surface is more susceptible to irreversible reactions in the first

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Table 15.6 Effect of particle size of graphite anode materials prepared by a mechanochemical method on their reversible lithium capacity

Sample LONZA KS-6 LONZA KS-25 LONZA KS-44

Average diameter of particle (μm)

Specific surface area (m2 g−1)

Active specific surface area (m2 g−1)

22

0.44

65

880

10.5

13

0.09

90

550

20.25

10

0.04

100

250

3.34

Lc* (nm)

Reversible capacity (mAh g−1)

* Crystal size along c axis.

charge–discharge cycle. It is usually observed that the enhanced capacity is associated with a pronounced charge potential hysteresis at >1.0 V (Salver-Disma et al., 1999). During charge–discharge cycles, the reversible capacity above 1 V decreases rapidly. Another report mentions that the charge slope at >1.0 V converts to a plateau at <0.25 V after repeated cycling or additional heat treatment at 1000°C, suggesting the disordered structure can transform into a more ordered graphitic one upon cycling or heat treatment (Ning et al., 2004). Although the reversible capacity of the milled graphite is high, its cycling behaviour is usually not satisfactory (Salver-Disma et al., 1999). The capacity fading is perhaps due to coalescence of edge planes which leads to a decrease in the number and activity of the edge planes with solvents. In addition, the filling of some vacancies and microcavities or voids by some removable bound interstitial carbon and the gradual penetrating of electrolytes into voids during the Li insertion and deinsertion process also result in capacity fading (Ning et al., 2004). Milling should not be too drastic or long in order to obtain optimal performance in terms of reversible capacity, coulomb efficiency in the first cycle and cycling behaviour. If milling is too long or drastic, unfavourable effects such as increase in irreversible capacity and decrease in coulomb efficiency will occur (Natarajan et al., 2001). Upon immersing the milled graphite powders in the electrolyte, a solid electrolyte interface (SEI) is formed spontaneously (Tossici et al., 2003). Different atmospheres including air, argon, nitrogen and CO2 produce different surface chemistries on the milled carbon, which influence the formation of SEI and the kinetics of the Li insertion process including Li+ migration through the surface films. Highly reactive carbon sites and highly reactive oxygen-containing surface groups participate in the formation of

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the surface film and do not result in a dense film that could prevent exfoliation, consequently capacity fades during cycling (Aurbach et al., 2002). In the case of reactive milling of graphite with lithium, since the surface is lithiated and bonding between carbon and lithium atoms is developed, intercalation compounds with Li/C molar ratios higher than 1/6 (Li/C = 1/6, 1/4 and 1/2) can be prepared (Tossici et al., 2003). These materials show higher reversible capacities than graphite (372 mAh g−1), weaker hysteresis and smaller irreversible capacities than milled graphite under the same conditions. A primary capacity very close to 1115 mAh g−1 is observed during the first deintercalation cycle at constant current. The following intercalation–deintercalation cycles yield capacity close to the theoretical value of 372 mAh g−1, typical of natural graphite (Tossici et al., 2003). As mentioned earlier, carbon anode materials manufactured from nongraphitic carbon also display structural changes after milling. As a result, the electrochemical performance will also change. Petroleum cokes, prepared by milling for 12–24 hours with an average particle size of 6–8 μm, show the best electrochemical characteristics. Tin-based anode materials The electrochemical performance of tin-based anode materials obtained by MC reactions is dependent on the processes and the additives used, as partially summarized in Table 15.4. The Sn3Co/Al2O3 composite powder has a higher initial coulombic efficiency and better cycle performance than Sn/ Al2O3. The electrode containing LA132 binder exhibits a reversible capacity of about 540 mAh g−1 with good capacity retention, as shown in Fig. 15.8 1000 100

80 600 60

400 Discharge Charge Efficiency

200 0

0

5

10

15 20 Cycle number

Efficiency (%)

Capacity (mAh g–1)

800

40

25

30

20 35

15.8 Cycling performance and efficiency of Sn3Co/Al2O3 obtained by ball milling using LA132 as a binder (reprinted from Zheng et al., 2007a with permission from Elsevier BV).

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(Zheng et al., 2007a). The Cu–Sn alloy particles obtained by ball milling show better electrochemical performance than those obtained by other methods such as the gas-atomizing and melt-spinning techniques (Xia et al., 2001). Mg2Sn alloy powder delivers a reversible capacity in the range of 250–300 mAh g−1 even after 20 cycles. Presumably, the interstitial space is mainly responsible for lithium intercalation and will not result in destruction of the phase composition, as happens when alloying of Sn with Li. The capacity of the mixture may be due to the layered structure of orthorhombic phase and the coexistence of cubic phase decreases the generated internal stress as lithium is inserted (Ning et al., 2004). However, it has also been reported that lithium can alloy with Sn in Mg2Sn as well as inserting into the Mg2Sn lattice. As a result, the initial reversible capacity is higher, 460 mAh g−1. Mg in these alloys does not participate in the alloying reaction with Li. When conductive metal and/or graphite are ball milled together with tin, the prepared composites deliver better cycling performance (Beaulieu et al., 2000; Lee et al., 2002; Wang et al., 2001). For example, the case Sn31Co28C41 composite presents a reversible capacity of 500 mAh g−1 without quick capacity fading (Hassoun et al., 2007). In summary, tin-based alloys prepared by MC methods may be used as anode materials for lithium-ion batteries, with better cycling behaviour than that of pure Sn particles, and their electrochemical performance is comparable with those obtained via chemical reaction. MC methods affect the structure of the prepared composites and their electrochemical performance. With further progress, novel tin-based alloy anode materials can be expected to appear in combination with other preparation techniques. Silicon-based anode materials It has become evident that composites based on Si obtained by MC reactions show better cycling behaviour, leading to promising anode materials; their electrochemical performance is also dependent on processing and additives, as partially summarized in Table 15.4. The composite particles consisting of nanosized Si, Li4SiO4 and other lithium-rich components from reactive milling of amorphous SiO with lithium metal demonstrate higher capacity and superior cyclability even at a high C-rate in comparison with Si nano powder (Yang et al., 2007). The nano-porous Si–graphite–C composite particles obtained using Al as the reductant exhibit a reversible capacity of about 700 mAh g−1 with no capacity loss up to the 120th cycle at a constant current density of 0.2 mA cm−2, as shown in Fig. 15.9.

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Capacity (mAh g–1)

1000

800

600

400

NP-Si NP-Si/graphite NP-Si/C NP-Si/graphite/C

200

0

0

20

40

60 80 Cycle number

100

120

15.9 Cycling performance of a series of nanoporous (NP) anode materials based on silicon prepared using high-energy ball milling (reprinted from Zheng et al., 2007b with permission from Elsevier BV).

The superior electrochemical characteristics are attributed to the nanosized pores in the Si particles, which suppress the volume effect, as well as to the buffering action and the excellent electronic and ionic conductivity of carbon materials (Zheng et al., 2007b). The silicon monoxide (SiO)– carbon composite powder obtained by ball milling following pyrolysis exhibits charge and discharge capacities of about 1050 and 800 mAh g−1, respectively, with a coulombic efficiency of 76% in the initial cycle. It shows high reversible capacity retention of 710 mAh g−1 over 100 cycles. The improved cycling performance is attributed to the stable microstructure, enhanced electrical contact afforded by the pyrolysed carbon, and the amorphous phase of the active material which accommodates volume expansion during cycling (Kim et al., 2007). If there is no pyrolysis process, the anode composite presents a poorer electrochemical performance (Doh et al., 2008). In the nanocomposites of silicon with electrochemically inactive or inert compounds such as TiN, TiC and Si3N4, the Si alloys and dealloys with lithium during cycling, while TiN, TiC and Si3N4 remain inactive, providing structural stability by acting as a buffer matrix. The very small size of the Si crystals prevents the identification of a distinct phase boundary between Si and inert compounds and the volume changes continuously rather than abruptly and discretely. As a result, excellent cycling behaviour is achieved. For example, composite powder particles containing 40 mol% silicon and 60 mol% TiC exhibit a stable capacity of around 380 mAh g−1 (Guo et al.,

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2005b). In the case of intermetallic alloys of Si with metals such as Mg, Ni, Fe, Mn, Ag and alloy compounds, Si provides active centres, which react with Li to form LixSi alloys, and the metal acts as a conductive agent buffering the volume change during the charge and discharge process, leading to excellent cycling behaviour (Hwang et al., 2001). The cycling behaviour depends on the milling time. The longer the milling time, the better dispersion of Si in the Ag matrix and the better cycling performance (Zhang et al., 2007a). Composites of Si with conductive agents such as carbon and polypyrrole present better cycling behaviour. In composite particles of silicon and polypyrrole which are obtained through mechanical milling, polypyrrole decreases the initial irreversible capacity of the silicon anode owing to the reduction in the thickness of the formed SEI layer, and the composite exhibits good reversibility, higher coulombic efficiency and a better cycle life than the bare silicon (Guo et al., 2005a). When more components are introduced, the electrochemical performance of the composites with Si from MC reactions can be further ameliorated. The Fe20Si80 alloy– graphite composite shows good cyclability and a high reversible capacity of about 600 mAh g−1 (Lee and Lee, 2002). When ball milling is combined with heat treatment, the electrochemical performance of the new composite particles based on Si presents great promise for their application as anode materials for lithium-ion batteries. For example, the Si/C/ PAN–C composites exhibit a reversible capacity of about 660 mAh g−1 with excellent capacity retention displaying almost no fade in capacity when cycled at a rate of about C/4, shown in Fig. 15.10 (Datta and Kumta, 2006).

(b) 32513 s

Potential (V)

1.2

62036 s

1.0 0.8 0.6

C/3.84

C/3.98 C/4.2

C/4.2

0.4 0.2 0.0

18657 s 0

47700 s

77223 s

20000 40000 60000 80000 10000 30000 50000 70000

Specific capacity (mAh g–1)

(a)

900 800 700 600 500 400 300 200 100 0

Discharge Charge

0

5

Time (s)

10

15

20

25

Cycle number

15.10 (a) Discharge/charge profiles and (b) cycling behaviour of Si/C/PAN–C composite electrode cycled at a constant current of 160 mA g−1 (reprinted from Datta and Kumta, 2006 with permission from Elsevier BV).

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Antimony-based anode materials Ball-milled β-Zn4Sb3 powder shows an initial reversible capacity of 566 mAh g−1 from 503 mAh g−1, compared with β-Zn4Sb3 without ball milling, and the plateau of lithium de-insertion from the alloy decreases from 0.95 V to 0.85 V, indicating a lower resistance for lithium ion in the milled materials. However, the cycling stability of the milled β-Zn4Sb3 is still not satisfactory, because during cycling several lithium-containing compounds such as LiZnSb, Li3Sb and LiZn are formed successively in the process of lithium insertion into Zn4Sb3 (Cao et al., 2001). The initial reversible capacity of CoSb3 powder depends on the ball milling time or the particle size. For example, the reversible capacity changes from 385 mAh g−1 for the coarse powders (about 6 μm) to 586 mAh g−1 for the ball milled fine powders (smaller than 100 nm). Although the capacity retention of the ball milled materials during cycling is poorer than expected owing to granular structures, after 50 cycles its volume reversible capacity still remains above 1600 mAh cm−3, which is almost twice that of graphite (Zhang et al., 2002). In the case of InSb powder obtained by ball milling, there is a reversible reaction between InSb and Li as shown in Equation [15.9]: xLi + InSb X LixInSb

(x ≤ 0.27)

[15.9]

When x > 0.27, the following reaction will take place: 3Li + InSb → Li3Sb + In

[15.10]

The voltage plateau of this process is above 0.65 V and the cycling behaviour is good. At a lower potential, the In produced reacts with Li and forms alloys with a composition InLix, resulting in poor cyclability (Hewitt et al., 2001). Perhaps at the first stage, In can act as a buffer to resist cracking or volume expansion. The composites of Sb-based alloys with carbon lead to better cycling behaviour in comparison with virginal alloys. For example, composite material CoFe3Sb12–C16 obtained via ball milling shows superior cycling even at a large current density of 100 mA g−1 (Zhao et al., 2001). Compared with Zn4Sb3 powder, composite Zn4Sb3–C7 powder prepared by ball milling possesses higher initial reversible capacity (580 mAh g−1), smaller voltage hysteresis and better capacity retention, indicating another promising anode material (Cao et al., 2001). Anode materials from nitrides In the product obtained by ball milling SnO with Li3N, the irreversible capacity is greatly decreased and the cycling behaviour is improved. The main reason for the improvement seems to be the existence of Li2O as a buffer and the disappearance of SnO (Foster et al., 2000).

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Lithium-rich silicon nitride composite particles possess a reversible capacity of about 500 mAh g−1 and the retained capacity after 50 cycles is above 400 mAh g−1, which is 80% of the reversible capacity of the second cycle. The structure of the composite particle is stable even at a high lithium extraction state and can still maintain excellent cycling performance (Wen et al., 2006). The lithium silicon oxynitride composite powder obtained via ball milling presents capacity retention of 488 mAh g−1 (95% of the second cycle) after 50 cycles. It is superior to many other high capacity anode materials for lithium-ion batteries. The capability for charging–discharging within a wide voltage cut-off window makes it suitable to compensate lithium-ion sources for different electrode systems of either cathode or anode (Wen et al., 2008). In Li–TM–N composites with carbonaceous materials obtained by mixing or ball milling, both Li–TM–N and the carbonaceous materials are electrochemically active within the potential window of 0–1.4 V versus Li/Li+. The charge and discharge behaviour of the composite powder reflects the mixed electrochemical characteristics of the two active hosts. Additionally, the carbonaceous hosts can function as conducting materials, leading to an improvement in the charge rate of Li–TM–N. High cycling stability is feasible owing to the low volume effects upon Li insertion and extraction. However, voltage hysteresis was found during charge and discharge in the electrode owing to the different reactive potentials of the two active hosts (Hanai et al., 2008). This research indicates that ball milling with nitrides such as Li3N and Li–TM–N is a good way to decrease the irreversible capacity. Doping of Li–TM–N powders by metals such as Cu and Fe leads to a slight decrease in capacity, an improvement in electrochemical kinetics associated with the modification of morphology characteristics of the active particles and a reduction in the potential hysteresis between the charge and discharge process (Liu et al., 2004b, c). The lithium transition metal nitrides therefore appear to be promising anode materials for lithium-ion batteries. Other kinds of anode materials The SnO–B2O3–P2O5 system prepared by the MC method has a discharge capacity in the first cycle of more than 500 mAh g−1 at a constant current of 1.5 mA cm−2. The charge–discharge curves of these materials from the MC method are similar to those for the glassy powders in the SnO–B2O3– P2O5 system prepared by a melt quenching procedure (Morimoto et al., 1999). By selecting the cutoff voltage, good cycling performance is achieved. The amorphous 50SiO⋅50SnO (mol%) obtained through ball milling the composite shows a high capacity of over 800 mAh g−1 in the potential range of 0–2.0 V versus Li+/Li with a current density of 1.5 mA cm−2. Its first dis-

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charge capacity is much larger than that of the amorphous SnO–B2O3–P2O5 materials obtained via mechanical milling. In addition, its irreversible capacity is also smaller (Morimoto et al., 2001). In the case of SnO2 doped with metal oxides, for example Sn1−xMoxO2, the metal oxides additives present two favourable effects, namely: (i) an increase of the discharge capacity and (ii) an improvement of capacity retention during cycling owing to the Li–M–O oxide conductive matrix (Martos et al., 2002). The electrochemical performance of nanosized SnS particles is dependent on the particle size, inactive buffer phase of Li2S and S and the ratio of Li2S to S. The irreversible capacity loss in the first discharge is attributed to the formation of Li2S, which plays an important role in the following discharge–charge cycles in preventing volume expansion and stabilizing the electrode structure. Li2S is more appropriate for buffering the volume change than S. In the following cycles, the nanosized SnS anode material synthesized via ball milling delivers high discharge capacity and exhibits good cycling stability (Li Y, et al., 2006b). Composite powders of SiO with porous carbons prepared via mechanical milling present high reversible capacity, excellent cycling behaviour and excellent high-rate discharge capability even at a discharge current density of 600 mA g−1. The main reason for this seems to be the embedding of Si–SiO in the continuous porous structure of carbon (Yuan et al., 2007; Chao et al., 2008). In the case of α-Fe2O3 powder obtained by mechanical milling, the material displays a high insertion capacity of over 1000 mAh g−1 (Morimoto et al., 2005). The electrochemical performance of the phosphides MPx obtained by mechanical milling, especially the initial capacity and capacity retention, is strongly related to the powder morphology: small particle size favours high capacity and the operational scan rate affects the capacity depending on the degree of crystallinity of the powder. On the other hand, capacity retention is better than that of microsized powders (Bichat et al., 2004). In the case of CoPx (CoP3 + CoP) prepared by ball milling, the resulting coulombic efficiency in the first cycle is increased from about 74% to 95%, along with a sacrifice of part of the reversible capacity. In contrast, the lithiated product by thermoreaction results in poor electrochemical performance (Zhang et al., 2005b).

15.5

Solid electrolytes from mechanochemical methods

Solid electrolyte is divided into inorganic electrolyte and solid polymer electrolyte (Wu et al., 2002a). Although solid polymer electrolyte is usually prepared by polymerization, MC methods are also applied in its prepara-

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tion, such as the addition of nanofillers (Wu et al., 2002b). However, it is not the main or determining process and this aspect will not be expounded here. Inorganic electrolyte usually includes oxide and sulfide powders. In the application to lithium-ion batteries, it should be amorphous or noncrystalline. As a result, MC methods using milling or grinding are promising ways to prepare inorganic solid electrolytes such as oxides, sulfides and oxysulfides (Wu et al., 2002b). Amorphous oxide electrolyte powders can be prepared via MC reactions between Li2O and SiO2, or P2O5 and GeO2 owing to the amorphization effect during the milling process. Other oxides such as B2O3 and lithium salts such as Li3PO4 can be added (Wu et al., 2002a). Nanosized amorphous Li1.4Al0.4Ti1.6(PO4)3 powders can be prepared by mechanical milling, which will be converted into glass ceramics after heat-treatment at 700–1000°C (Xu et al., 2006). Lithium-ion conducting amorphous solid electrolytes based on sulfides are first prepared from SiS2 and other sulfides. For example, grinding crystal Li2S with SiS2, P2S5 or P2S3 produces an amorphous glass. The amorphousforming region is extended in comparison with that obtained by meltquenching, and the molar ratio of Li2S in amorphous xLi2S⋅(1 − x)P2S5 can be as high as 0.80 (Tatsumisago et al., 2002). During the formation of Li2S– P2S5-based glass ceramics, softening of the glass occurs concurrently, which must be used to form close solid/solid contact between electrolyte and electrode powders without any chemical reactions because of low heattreatment temperature (Tatsumisago, 2004). Superionic crystals, a series of sulfide crystalline solid electrolytes such as Li4GeS4–Li3PS4, have been successfully formed by the crystallization of mechanically milled Li2S–P2S5 glasses and GeS2 (Hayashi et al., 2003). Amorphous composite 75Li2S⋅xP2S3⋅ (25 − x)P2S5 (mol%) can be obtained in the range of x = 0–5. In composites with x ≥ 8.3, a crystalline Li2S phase remains after ball milling for 50 hours (Machida et al., 2005). Considering the experience with sulfides, oxides such as LixMOy (M = Si, P or Ge) and P2O5 can be further incorporated by ball milling (Hayashi et al., 2002; Ohtomo et al., 2005). The glassy composite powders obtained are pelletized and heated to various temperatures over their crystallization temperatures to form the glass–ceramic materials (Hayashi et al., 2002; Tatsumisago and Hayashi, 2008). During milling, the particle size at first decreases with milling time and later increases because the particles aggregate with each other. The local structure of the mechanically milled amorphous materials containing Li4SiO4 is similar to that of the corresponding melt-quenched glasses, in which SiS4 and SiOS3 tetrahedral units are mainly present. SiOnS4−n (n = 1, 2, 3) tetrahedral units are formed when adding Li4SiO4 and Li4GeO4 to the sulfide system, however these units are not present when Li3PO4 is added. The reactivity of LixMOy derived from its

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basicity affects the structure and formation process of the oxysulfide materials obtained by mechanical milling (Hayashi et al., 2002). As well as oxides, lithium compounds such as Li2S, Li3N, Li2O, Li3PO4, LiOH, LiCl and LiI can be added to sulfides and solid electrolytes are obtained through mechanical milling (Hayashi et al., 2004; Ning et al., 2004). Too long a grinding time will lead to partial crystallization (Ning et al., 2004). The electrochemical properties of the prepared solid electrolytes vary according to the materials and ball milling treatment employed. In the case of the Li2O–B2O3–P2O5 composite oxides, the total ionic conductivity can be as high as 9 × 10−5 S cm−1 (Wu et al., 2002a). The ionic conductivity of the glass–ceramic Li1.4Al0.4Ti1.6(PO4)3 increases with the increasing amount of crystalline phase and the reduction in grain size. The highest bulk conductivity (σb) is 1.09 × 10−3 S cm−1 at room temperature with an activation energy as low as 0.28 eV (Xu et al., 2006). Amorphous sulfide composites such as 60Li2S⋅40SiS2 (mol%) prepared via ball milling show ionic conductivities up to 10−4 S cm−1 at room temperature, comparable with those obtained by melt quenching, and the transport number of Li+ is nearly 1.0. When these composites are heated to suitable temperatures to obtain glass ceramics, the ionic conductivity will be further increased. For example, as shown in Fig. 15.11, the Li2S–P2S5 glass ceramic

0

Log (conductivity / S cm–1)

–1

Glass–ceramics

–2 –3 –4 Glass

–5 –6 –7 Solid-state reaction

–8 –9 1.8

2

2.2

2.4

2.6 2.8 1000 / T

3

3.2

3.4

15.11 Temperature dependence of the conductivities of the 70Li2S– 30SiS2 glass and glass–ceramics prepared by the MC method. The conductivity data for the sample prepared by solid state reaction are also shown for comparison (reprinted from Tatsumisago and Hayashi, 2008 with permission from Elsevier BV).

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shows ionic conductivity as high as 3.2 × 10−3 S cm−1 at room temperature (Tatsumisago and Hayashi, 2008). After the addition of oxides by the MC method, the prepared amorphous oxysulfide materials of (100 − x)(0.6Li2S⋅0.4SiS2)⋅xLi4SiO4 show an ionic conductivity above 10−4 S cm−1 and an electrochemical window up to 10 V (Hayashi et al., 2002). The addition of a small amount of P2O5 to the Li2S–P2S5 sulfide system in the ball milling process lowers the ionic conductivity, but enhances the electrochemical stability of the glass– ceramic electrolytes (Ohtomo et al., 2005). Amorphous solid electrolytes from SiS2 and various kinds of lithium compounds such as Li2S, Li3N and Li2O exhibit high conductivities of 10−5–10−4 S cm−1 at room temperature, while those obtained using Li3PO4, LiOH and LiCl present conductivities lower than 10−6 S cm−1. The conductivity of the 50Li2O⋅50SiS2 (mol%) composite drastically increases with the milling time from 0 to 5 hours. The enhancement is due to the formation of corner-shared and isolated SiS4 and SiOS3 tetrahedral units in the solid electrolyte, leading to lower activation energy for the movement of lithium ions (Hayashi et al., 2004). When these kinds of amorphous sulfides and oxysulfides are used as electrolytes and assembled into solid-state lithium-ion batteries, they exhibit excellent cycling performance (Hayashi et al., 2002; Machida et al., 2005). For example, an all-solid-state battery fabricated with LiCoO2 fine powder as cathode material, the glass ceramic 80Li2S⋅20P2S5 (mol%) as solid electrolyte and a metallic indium foil as anode shows a rechargeable capacity of 100 mAh g−1 based on the cathode material, and its cycle efficiency is almost 100% after the second cycle, with no capacity fading for 500 cycles (Machida et al., 2005).

15.6

Conclusions

In conclusion, MC methods are promising methods for preparing materials for lithium-ion batteries. Compared with conventional solid state reactions at high temperature, the MC methods appear to accelerate and simplify the synthesis process and decrease the energy expenses as well as the cost of the material. Furthermore, the prepared materials present good electrochemical performance. For example, cathode materials such as LiMn2O4 spinels present better cycling behaviour in comparison to those from solid state reactions because the highly disordered nanocrystalline particles can accommodate the Jahn–Teller distortion. Carbonaceous anode materials such as graphite show very high capacity due to the introduction of many defects, such as voids/nanocavities and the change into nanoparticles (Kosova and Devyatkina, 2004). In the case of alloy-based anode materials, their cycling behaviour is also improved since heteroatoms can be homo-

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geneously incorporated into the prepared nanocomposites and they buffer the drastic volume change. The prepared solid glass electrolytes show high conductivity (>10−4 S cm−1) and the fabricated all-solid state lithium-ion battery presents good cycling behaviour. In addition, MC methods can be combined with other techniques to prepare materials for lithium-ion batteries. As mentioned above, when MC methods are combined with heat treatment, glass–ceramic electrolytes can be prepared with high ionic conductivity, around 10−3 S cm−1 at ambient temperature (Tatsumisago et al., 2002; Akitoshi et al., 2003). Grinding in combination with subsequent firing at 400°C or higher can enable high purity spinel Li4Mn5O12 to be prepared which is difficult to obtain by heat treatment alone (Tanaka et al., 2003). Combining with electric discharge, which favours the formation of crystal structure, will lead to preparation of versatile electrode materials for lithium-ion batteries such as LiFePO4 (Needham et al., 2006). MC reactions cannot induce every material to display the best electrochemical performance since some actions are not favoured by ball milling. For example, during the mechanical milling process, sometimes the structure change will lead to poor cycling performance (Wang and Kumta 2007; Zhang et al., 2005a). One main reason for the deteriorated performance is the resulting amorphization. Another one is perhaps related to the change in morphology (Bichat et al., 2004). In the case of the composite of Si with graphite, a diffusion barrier such as polymethacrylonitrile must be added to suppress the possible MC reaction between silicon and graphite particles to form SiC and further to prevent the amorphization of graphite during extended milling (Wang and Kumta 2007). The agglomeration of the composite particles during the high energy mechanical milling process and low coulombic efficiency in the initial cycle still remain to be solved (Wang et al. 2008a). In some situations, good crystallinity could not be achieved since the ball-milling could not provide energies sufficient to arrive at a temperature above 400°C (Gillot et al., 2007). Furthermore, there is still much development expected from further research into the effects of MC methods on the materials for lithium-ion batteries. With further research and development, the application of MC reactions can be expected to expand even more (Myung et al., 2006).

15.7

Acknowledgement

Financial support from the National Basic Research Program of China (973 Program No: 2007CB209702), the Shanghai Committee of Science and Technology (09QH1400400) and the South African National Research Foundation (UID No. 67295) is gratefully acknowledged.

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Index

acetylene black, 373 acetylsalicylic acid, 230 acylation reactions, 226 adamantine, 318, 320 adhesion, 243 AGO-2, 11 Akashi Hardness Testing Machine, 270 Al7Cu2Fe, 155 Al13Cu4Fe3, 155 Al8Fe2Si, 155 alkaline niobates, 129, 130 aluminium plates, 261 Al sample with hammered Ti particles, 259 entirely coated Al surface with Ti particles, 262 ammonia, 316 amorphisation, 234 anastase, 305–6, 311 anchoring, 269 anode materials alloy-based anode materials, 378–86 ball milled C0.8Sn0.2 composite, 382 selected results, 379–80 Si/PPy composite particle, 385 Sn and Li2O composite particles, 381 carbonaceous materials, 374–8 cycling performance discharge/charge profiles of Si/C/ PAN–C composite electrode, 394 efficiency of Sn3Co/Al2O3, 391

nanoporous anode materials based on silicon, 393 electrochemical performance, 389–97 anode materials from nitrides, 395–6 antimony-based anode materials, 395 carbonaceous anode materials, 389–91 other kinds of anode materials, 396–7 silicon-based anode materials, 392–4 tin-based anode materials, 391–2 mechanochemical preparation, 374–97 from nitrides, 386–7 ball milling effects, 387 other kinds of anode materials, 387–9 carbon xerogel and CX–SiO, 388 anthracene, 81 arabinogalactan, 236, 237, 239 arc-melting process, 384–5 Archimedes method, 282 armoured surface, 260 artificial lattices, 63 ascorbic acid, 242 attritor mill, 137, 139, 186 β-cyclodextrin, 235, 236 β-sialon, 187

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410

Index

β-SiC, 177 β-Zn4Sb3 powder, 395 ball-impact energy, 114 ball-impact frequency, 114 ball milling, 10, 12, 94, 149, 252, 275 devices and summary of conditions investigated, 280 ball mills, 16, 46–7, 93 ball to powder mass ratio, 275 barium titanate, 117, 128 BaTiO3 see barium titanate benzoquinone, 81, 82 bicarbonate, 241 biexponential kinetics, 238 BiFeO3, 131–2 binary alloy systems, 152–5 Fe–X (X = Sn, Zn, Sb, Al) systems, 153 Ni–X (X = Sn, Sb, Zn, Al) systems, 153–4 other systems, 154–5 binary composite system, 322 Binol, 82 bismuth titanate, 122, 130–1 biuret, 304 Boltzmann factor, 33 borides, 292 Bragg diffraction peaks, 388 Bridgman’s anvils, 225, 226 brookite, 305–6 bulk composites, 193 calcination, 72 calcium carbonate, 243 carbide powders, 173 carbides, 173–9, 292 carbon, 384 carbon cryogel, 388 carbon monoxide, 57–9 carbon xerogel, 388 carbonaceous materials, 374–8 carbonato complex, 123 carbothermal reaction, 185 carbothermal reduction, 185–7 carboxylic acids, 244 catalytic hydrogenation, 57–9

cathode materials ball-milling time effects on LiMn2O4 cycling behaviour, 372 electrochemical performance, 371–3 Li3Fe1.8Ti0.1Mn0.1(PO4)3 and ball milled Li3Fe1.8Ti0.1Mn0.1(PO4)3/ carbon mixture, 371 lithium cobalt oxides, 365–6 mechanochemical preparation, 365–73 other kinds of cathode materials, 369–70 spinel lithium manganese oxides, 366–9 cellulose, 235 ceramic method, 347 ceramic oxides, 113–40 complex oxides with various properties, 126–39 future trends, 139–40 mechanisms and kinetics, 114–26 ceramics, 194 cermetals, 194, 217 charge transfer complex, 82 Chirchik Plant, 2 chitosan, 236 chlorinated aromatics, 50–3 chlorobenzene, 50 citric acid, 242 cobalt, 99 CoFe3Sb12, 386 (Co50Fe50)0.2(TiO2)99.8 catalyst, 58–9 combustion synthesis, 169 composite powder, 194–203 conductive carbon, 370 copper, 217 Coulomb interaction, 239 Coulomb law of friction, 12 coulombic efficiency, 397 covalent bonds, 167 critical loading conditions, 96, 97, 102 critical nuclei, 26 critical threshold value, 96 Cr3O3, 135 cryomilling technique, 297 CSEM Revest, 271 Cu/Al2O3, 219

© Woodhead Publishing Limited, 2010

Index Cu–Al/Al2O3 composite powder lamellar structure, 197 mechanosynthesis GTM results, 200 SHS GTM results, 201 structure homogeneity, 199 CuO–Al system Al2O3 particles embedded in Cu–Al matrix, 213 effects of mechanosynthesis, 211–13 XRD patterns after different treatment times, 212 Cu2(OH)2CO3–Al system after different treatment times, 209 after mechanochemical synthesis, 206 DTA curve, 202 effects of mechanosynthesis, 208–11 TG–DTA curves, 211 Curie law, 377 Curie temperature, 131 cyanuric acid, 316 cyclisation, 326 cyclodextin, 243 degree α, 27 Degussa P25, 311, 322 densification, 168 diamagnetic imidazoline, 229 Diels–Alder reaction, 64, 80–2 effect of alcohols on rate, 82 with a eutectic complex, 83 yields under various conditions, 81 differential scanning calorimetry, 234 differential thermal analysis, 69, 201 diffuse reflectance spectroscopy, 314–15 dimethylanthracene, 82 effervescent tablet, 241 El 2 × 150 mill, 20 electric current activated sintering, 276 electrochemical photolysis, 305 electron-hole separation, 322–3 electron probe microanalysis, 261, 262 electron-trapping effect, 322 electronic diffuse reflectance spectra, 345

411

ethanol, 318 etherification, 226 extractive metallurgy, 1 extruders, 225 Fermi level, 306 ferroelectric oxides, 126–33 Fe20Si80 alloy–graphite composite, 394 Finnigam Tracker Mass Spectrometer GC/MS PE 8420, 50 Fisons 8000, 50 Fisons Sorptomatic 9600, 50 flame pyrolysis, 317 flash point, 202 fluorides, 228 Fourier transform infrared spectroscopy, 385 Fritsch Pulverisette 4, 116 gas-atomising, 392 gas pressure–temperature measuring system, 204 Gibbs energy, 177 glass–ceramic electrolytes, 400 GN-2 ball mill, 175 grain boundary migration, 68 graphene, 374 graphite, 374, 400 mechanochemical reactions effect on structure and electrochemical performance, 375–6 particle size effect on their reversible lithium capacity, 390 GTM see gas pressure–temperature measuring system Halder–Wagner method, 281 hammering effect, 258 HBB see hetero-bridging bond Hertz–Routh theory, 14 Hertz theory, 12, 14, 16, 21–2 illustration, 13 hetero-bridging bond, 69, 71 hetero-metalloxane bond, 69 heteroatoms, 374, 386, 400 hexachlorobenzene, 50, 228, 229

© Woodhead Publishing Limited, 2010

412

Index

hexamethylenetetramine, 311, 316, 320 high-energy ball milling, 160, 195 see also mechanical alloying vs solid–liquid reaction ball milling, 163–4 high-energy mechanical milling, 384, 385 highest occupied molecular orbital, 68 HOMO see highest occupied molecular orbital hydrogen, 53 absorption on Mg2Ni/Ni composite, 53–7 hydrogen bond, 239 hydrometallurgy, 64–5

hydrogen moles absorbed as a function of time, 54 mechanically induced structural and chemical transformations, 45–60 mechanochemistry, 46–8 outline of experimental methods, 48–50 quantity as a function of scaled time, 56 as a function of time, 55 stainless steel reactor temperature, 51 Kyence Color 3D Laser Confocal Microscope VK9700, 270

ignition temperature, 199 ignition time, 198–9 INCO laboratories, 2 induction effect, 351 ingot metallurgy, 158 injected mechanical power, 117 inorganic compounds, 331–3 intermetallic compounds, 150 ion-plating technique, 271 iron, 153, 386

LA132 binder, 391 La-monazite coatings, 251, 253, 272–3 LaCoO3, 137 LaCrO3, 135 LaMnO3, 134, 135, 136 lantern units, 353 lanthanophosphate coatings characterisation, 264, 266–7 cross-section deposited at different levels, 265 stainless steel surface, 263 degree of coverage and coating microstructure, 261–4 EPMA mapping of coated samples, 263 microhardness and peeling force, 266 on stainless steel plates, 261–7 lanthanum oxyfluoride, 324 LaOF see lanthanum oxyfluoride LaPO4 coatings see lanthanophosphate coatings lead zirconate titanate, 128–9, 129 Lewis classification, 333 LiCoO2 see lithium cobalt oxides LiFePO4, 348–51 Li2FeSiO4, 351–3 LiMn2O4, 124, 333–6 LiNi0.5Mn1.5O4, 336–7 LiNi1–x–yCoxMnyO2, 347–8

Jahn–Teller distortion, 68, 83, 333, 368, 400 JEOL Cross Section Polisher SM-09010, 262–3, 270 jet milling, 374 JXA-8200, 262 K-type thermocouple, 256 KhK 871 mill, 19 kinetic behaviour carbon monoxide catalytic hydrogenation, 57–9 conversion data, 58 chlorinated aromatics degradation, 50–3 fraction α of reacted chlorobenzene, 53 future trends, 59–60 hydrogen absorption on Mg2Ni/Ni composite, 53–7

© Woodhead Publishing Limited, 2010

Index liquid-phase oxidation, 230 liquid reactants, 154 lithium cobalt oxides, 344–7, 365–6 materials prepared by mechanochemical processes, 367 lithium-ion batteries anode materials mechanochemical preparation, 374–97 alloy-based anode materials, 378–86 anode materials from nitrides, 386–7 carbonaceous materials, 374–8 electrochemical performance, 389–97 other kinds of anode materials, 387–9 cathode materials mechanochemical preparation, 365–73 electrochemical performance, 371–3 lithium cobalt oxides, 365–6 other kinds of cathode materials, 369–70 spinel lithium manganese oxides, 366–9 DTA curves of initial and activated mixtures Li2CO3 with FeC2O4·2H2O and SiO2·0.1 H2O, 353 LiOH with MnO2, 335 IR spectra LiOH + Co(OH)2 activated mixture, 346 LiOH + V2O5 activated mixtures, 343 Li NMR spectra of LiOH + MnO2 activated mixture, 336 LiOH + Co(OH)2 activated mixtures diffuse reflectance spectrum, 346 LiOH and Li2CO3 reactions with crystal hydrates and acidic salts, 348–54 LiFePO4 synthesis, 348–51

413

Li2FeSiO4 synthesis, 351–3 LiTi2(PO4)3 synthesis, 353–4 LiOH reactions with anhydrous oxides, 333–43 LiMn2O4 synthesis, 333–6 LiNi0.5Mn1.5O4 synthesis, 336–7 Li4Ti5O12 synthesis, 338–9 Li1+xV3O8 synthesis, 339–43 LiOH reactions with solid hydroxides, 344–8 LiCoO2 synthesis, 344–7 LiNi1–x–yCoxMnyO2 synthesis, 347–8 70 Li2S–30SiS2 glass and glass– ceramics conductivities temperature dependence, 399 materials for cathode, anode and electrolyte, 365 mechanochemical methods, 361–401 Mössbauer spectrum of Li2CO3, FeC2O4·2H2O and NH4H2PO4, 352 principle, 362–3 requirements for materials, 364–5 soft mechanochemical synthesis of materials future trends, 357 principles of solid inorganic compounds synthesis, 331–3 solid electrolytes from mechanochemical methods, 397–400 TG and DTA curves of initial and activated mixtures LiOH + TiO2 + NH4H2PO4, 354 LiOH + V2O5, 341 LiOH and TiO2, 339 LiOH with FeC2O4· 2H2O and NH4H2PO4, 350 LiOH with Ni0.8Co0.2(OH)2, 348 LiOH with NiO and MnO2, 337 X-ray patterns LiNi1-yCoyO2, 349 LiOH + Co (OH)2 activated mixtures, 345 LiOH + FeC2O4·2H2O and NH4H2PO4, 351

© Woodhead Publishing Limited, 2010

414

Index

LiOH + NiO + MnO2 activated mixture, 338 LiOH + TiO2, 340 LiOH + TiO2 + NH4H2PO4 mixture, 355 LiOH + V2O5 activated mixture, 342 LiOH activated mixtures with different manganese oxides, 334 lithium manganese oxide, 370, 373 lithium silicon oxynitride composite, 387, 396 lithium transition metal nitrides, 386–7, 396 Li4Ti5O12, 338–9 LiTi2(PO4)3, 353–4 Li–TM–N see lithium transition metal nitrides LixMnOz see lithium manganese oxide Li1+xV3O8, 339–43 lowest unoccupied molecular orbital, 68 LUMO see lowest unoccupied molecular orbital magnesium carbonate, 243 magnetic oxides, 133–4 malachite green oxalate, 314 matrix-isolated imidazoline radicals, 229 mechanical activation, 47–8, 134, 195–6, 275, 332 mechanical alloying, 65, 66, 149–50, 196, 252 cycle, 27 kinetic processes and mechanisms, 92–107 chemical mixing mechanisms, 103–6 collecting and analysing experimental data, 98–103 fundamentals of processes in ball mills, 93–5 future trends, 106–7

instantaneous temperature map for square domain, 106 local temperature time, 105 mass fractions as a function of the number of collisions, 100 Ni40Ti60 mass fraction, 101 NiTi2 phases mass fractions, 100 phenomenological model kinetics, 95–8 rotating atoms cluster, 105 square of thickness of amorphous layer, 104 vs solid–liquid reaction ball milling, 161–3, 163–4 mechanical diffusion, 65 mechanical doping synthesis anion-doped, titanium dioxide, visible-light induced photocatalysts, 310–25 photocatalysts synthesis, 304–26 mechanical energy, 2 mechanical stressing, 66 mechanically induced self-propagating reaction, 174 mechanochemical activation, 113–14 mechanochemical alloying, 1–2 mechanochemical processing, 1–4 composite oxides in electromagnetic applications, 68–80 bond overlap population, 72 HBB formation during milling, 70 homogeneity of starting mixtures for BaTiO3 preparation, 73 Mg(OH)2 and Mg(OH)2-SiO2 change in thermal gravimetryDTA profiles, 69 Mg(OH)2-SiO2 change in FT-IR spectra, 70 molecular orbital calculation, 71 percent perovskite against firing temperature, 80 PMN and PZN perovskite ceramics application, 76–80 quenched samples XRD profiles, 75 reaction mixture for PMN-xPT SEM micrographs, 76

© Woodhead Publishing Limited, 2010

Index soft mechanochemical effects principle, 68–71 definition, 1 degree α of transformation for BaCO3 + WO3 = BaWO4 + CO2 density of balls, 34 frequency of rotation, 34 number of balls, 35 radius of balls, 35 departure from equilibrium achieved from various processes, 3 effects on graphite structure and electrochemical performance, 375–6 manufacturing non-oxide powders, 169–80 ball size effect on milled powder lattice strain and crystallite size, 178 carbides, 173–9 milled and heat-treated powders, 180 milled silicon and graphite powders XRD patterns, 177 nitrides, 170–3 silicides, 179–80 silicon-based mixture maximum combustion temperature, 171 Ti-C powder TEM micrographs, 176 titanium and activated carbon reaction progress, 175 materials design, 63–85 forced diffusion schematic illustration, 65 future trends, 84–5 mechanochemical route benefits, 64–7 modern mechanochemistry theoretical aspects, 67–8 materials for lithium-ion batteries, 361–401 anode materials preparation, 374–97 cathode materials preparation, 365–73

415

lithium-ion batteries, 362–5 solid electrolytes, 397–400 mechanism and kinetics, 9–40 calculating expressions, 19–20 change in specific surface of barium carbonate and tungsten oxide, 31 dependence of real time of activation process, 29 general issues, 15–16 Hertz theory and Routh’s Ω-hypothesis illustration, 13 impact-friction contact kinematics and dynamics, 13 impact-friction contact of two particles, 18 kinetics, 27–32 mechanism, 22–7 milling tools impact on particle layers and lining, 17 modern vibratory and planetary mills, 11 NaNO3 + KCl = KNO3 + NaCl degree α of transformation, 25 planetary ball mill scheme for calculation, 12 silver oxalate change in specific surface and conformity, 31 stress wave propagation, 17 temperature impulse at impactfriction contact, 23 t–P–T conditions calculation, 16–22 non-oxide systems with highly covalent bonds, 167–88 basic properties, 169 reactive systems, 180–7 α-Si3N4 unit cell volume vs milling time, 184 carbothermal reduction, 185–7 defective β-Si3N4 particle TEM image, 184 gas-pressured sintered β-sialon microstructure, 186 precursor activation in subsequent densification, 181–5

© Woodhead Publishing Limited, 2010

416

Index

sialon precursor powder specific surface area vs milling time, 183 spontaneous chemical reactions organic synthesis and utilisation, 80–4 change in energy of system and dihedral angle of anthracene, 84 Diels–Alder addition reaction, 80–2 increase in XRD peaks with fraction of PZN, 81 intermolecular HOMO-LUMO gap, 85 molecular distortion and change in frontier orbitals, 83–4 use of kinetic equations, 32–40 BaWO4 experimental values for mechanical synthesis, 39 degree of transformation based on compressibility, 36 experimental kinetic curves, 36 mechanochemical activation conditions effect on kinetics, 38–40 reactions controlled by diffusion, 37–8 reactions controlled by kinetics, 38 substances mechanical activation, 33, 37 mechanochemical synthesis complex ceramic oxides, 113–40 BaCO3–TiO2 powder mixture XRD patterns, 118 mechanochemical reactions mechanisms and kinetics, 114–26 milling efficiency graphical representation, 116 milling parameters, 115 NiFe2O4 non-uniform core/shell structure, 125 nucleation and growth mechanism schematic, 120

PbO–MgO–Nb2O5–TiO2 quantitative phase composition, 121 complex oxides with various properties, 126–39 BaCO3–TiO2 powder XRD patterns, 129 complex oxides with various structures and properties, 127 Cr(OH)3-nH2O XRD patterns, 136 ferroelectric and related oxides, 126–33 La2O3–Mn3O4 XRD patterns, 137 magnetic oxides, 133–4 oxides with semiconducting and catalytic properties, 134–9 PMN–PT thick films XRD patterns, 132 Zn–Fe–O sorbents sulphur absorption capacity, 138 composite powder formation, 194–203 Cu2(OH)2CO3–Al mixture DTA curve, 202 Cu2(OH)2CO3–AL2O3 synthesised composite powder elemental microanalysis results, 198 mechanical activation, mechanical alloying, reactive milling, 195–203 mechanical treatment, 195 SHS mechanochemical reactions induction time intervals, 203 copper-based composite powders with Al2O3, 207–13 CuO–Al reagent system, 211–13 Cu2(OH)2CO3–Al reagent system, 208–11 metal oxides formation enthalpy values, 208 Cu–Al–O and Ni–Al–O elemental systems, 217–19 low-molecular organic compounds solid-state mechanochemical synthesis, 227

© Woodhead Publishing Limited, 2010

Index main sequence and products, 226 metallic–ceramic composite powders, 193–200 applied high-energy milling, 205–6 Cu2(OH)2CO3–Al system after mechanochemical synthesis, 206 nickel-based composite powders with Al2O3, 214–16 combustion aluminothermic reaction product structure, 218 NiO–Al reagent systems, 216 Ni(OH)2CO3·xH2O–Al reagent system, 214–16 organic compounds and rapidly soluble materials, 224–45 mechanochemically activated powders, 281–2 mechanochemistry, 1, 46–8, 65, 149, 195 theoretical aspects, 67–8 mechanosynthesis, 66 melt-spinning techniques, 392 melts, 150 Mendeleev periodic table, 356 metal carbonates, 241, 244 metallic–ceramic composite powders mechanochemical synthesis, 193–220 applied high-energy milling, 205–6 composite powder formation, 194–203 copper-based composite powders with Al2O3, 207–13 Cu–Al–O and Ni–Al–O elemental systems, 217–19 monitoring mechanochemical processes, 203–5 nickel-based composite powders with Al2O3, 214–16 metallothermic reactions, 206 metallothermic reduction, 206 methane, 57 methanol, 307 MFe2O4, 133 Mg2Ni/Ni composite, 53 Mg2NiH4, 53 Mg2Sn alloy, 392 microhardness, 264

417

milling map, 114 milling medium, 153, 154 milling parameters, 115 Mixer/Mill mod. 8000, 48 molecular mass distribution, 237 molybdenum, 179 molybdenum disilicide system, 187, 281, 282–7, 293 milling time and charge ratio effect on product formation, 285 Mo crystallite size dependence on milling time, 283 temperature and sample displacement, 286 XRD patterns co-milled Mo + 2 Si powders, 284 Mo as function of milling time, 283 starting mixture and final product, 287 monazite-type lanthanum phosphate, 260 MoSi2 see molybdenum disilicide Mössbauer spectroscopy, 350 Na2CO3, 122–3 NaNbO3, 114, 123, 130 nanocoatings, 267–72 characterisation, 270–2 nanofillers, 398 NASICON, 353 NbAl3 system, 293–5 milling time effect on composition of samples obtained by SPS, 294 temperature and mechanical load temporal profiles, 295 XRD patterns of reactants as function of milling time, 293 Néel temperature, 131 neutralisation, 326 Newton’s theory of collisions, 10 (NH4)2CO3, 325 NiAl, 214 Ni3Al, 214 Ni2Al/Al2O3, 216 nickel, 153, 214, 217 Nikolsky Plant, 2

© Woodhead Publishing Limited, 2010

418

Index

NiO–Al system effects of mechanosynthesis, 216 XRD patterns after mechanochemical treatment, 217 niobium aluminides, 277 Ni2(OH)2CO3·xH2O–Al system effects of mechanosynthesis, 215–16 XRD patterns after mechanochemical treatment, 216 NiTi2, 100 nitrides, 170–3 anode materials preparation, 386–7 nitrogen monoxide, 308, 315 non-linear elastic theory, 15–16 non-oxide systems compounds basic properties, 169 with highly covalent bonds mechanochemical processing, 167–88 powder manufacturing by mechanochemical processing, 169–80 reactive systems mechanochemical processing, 180–7 O2 plasma treatment, 326 O-ring systems, 49 octachloronaphthalene, 228 Olivine LiFePO4, 370 organic acid–metal carbonate, 231 organic compounds benzoic acid/BaSO4 DSC thermograms, 235 break area of aggregated particle sample, 233 changes in morphology of reagents particles, 231 citric acid/CaCO3 mixture microphotographs, 232 fluorinated products yields after 20 hours autoclave treatment, 229 low-molecular organic compounds solid-state mechanochemical synthesis, 225–9, 227 possibilities in various types of reactions, 225–8

solid reagents preliminary activation for subsequent heterophase reactions, 228–9 mechanochemical synthesis, 224–45 future trends, 244–5 solid particles aggregation on reactivity carbon acids neutralisation, 230–5 stable iminoxyl radicals formation, 229–30 solid particles aggregation on reactivity in mechanochemical reactions, 229–35 synthesis for pharmaceutical applications, 235–44 low molecular organic compounds and polysaccharide adducts, 235–41 pharmaceutical substances solubility in water solutions, 238 rapidly soluble composite materials mechanochemical preparation, 241–4 oxidation–reduction, 326 oxide-dispersion strengthening, 1, 2 oxygen, 262 oxysulfides, 400 particle-reinforced surface, 260 PCA see process control agent Pechini process, 332 pentachloropyridine, 228 Periodic law, 356 Perkin-Elmer 8600, 50 perovskite, 382 perovskite ceramics BaTiO3, 72–6 granulometric properties, 76 obtained by calcining, 75 perovskite ceramics application, 72–6 PMN and PZN, 76–80 PMN-0.1PT dielectric constant frequency dependence, 79

© Woodhead Publishing Limited, 2010

Index PMN-xPT lattice parameter and average grains size, 78 PMN-xPT reaction mixture x-ray photoelectron spectra, 77 PMN-xPT XRD profiles, 78 petroleum cokes, 391 phenol, 308 Phillips X’Pert diffractometer Cu K α, 208 phosphide CoP3, 389 phosphides MPx, 389, 397 photocatalysis, 325 photocatalysts environment purification system by photocatalytic reaction, 309 mechanical doping synthesis, 304–26 anion-doped, titanium-dioxide, visible-light induced, 310–25 future trends, 325–6 titanium dioxide and its photocatalytic applications, 305–10 photocatalytic system, 304 pilling force, 264 piroxicam, 236 planetary ball milling, 314, 315 planetary mills, 10, 67, 116, 118, 128, 130, 131, 139, 177, 182, 186, 225, 226, 236, 377 milling parameters, 115 scheme for calculation, 12 PMN, 119 PMN–PT, 119, 120 polyethylene glycol, 235, 243 polymer alloys, 65–6 polymorphism, 234 polypyrrole, 384 polysaccharide, 239 polysaccharide–arabinogalactan systems, 240–1 polytetrafluoroethylene, 178, 324 polyvinylpyrrolidone, 235, 243 potassium bicarbonate, 242 potassium fluoride, 229 potassium persulphate, 229 powder charge, 49, 96 powder mass, 49, 96

419

powder processing, 94 powder technology, 64 precoating, 261 process control agent, 157–8 proton affinity, 356 pulsed mechanical action, 225 Pulverissette 6, 208 pyrometallurgy, 64 pyrometer, 281 PZT see lead zirconate titanate rapidly soluble materials mechanochemical synthesis, 224–45 future trends, 244–5 pharmaceutical applications, 235–44 low-molecular-weight organic compounds and polysaccharide adducts, 235–41 rapidly soluble composite materials mechanochemical preparation, 241–4 possible components, 242 reactant, 152 reactive milling, 149–50, 150, 160, 198 reactor chamber, 49 Rietveld method, 99 Rietveld refinement analysis, 119 Rigaku D/Max diffractometer, 50 roll mills, 225 Routh’s ξ-hypothesis, 12 illustration, 13 rutile, 305–6, 311 Scion Image software, 256 self-propagating high-temperature synthesis, 169, 199, 204 shaker mill, 117, 122, 131, 133, 139 Sherrer’s formula, 183 shock-attrition mills, 225 SHS see self-propagating hightemperature synthesis Si/C/PAN–C composites, 394 sialon ceramics, 182–3 sialon powders, 172 Siberian larch, 236 silicides, 179–80 silicon, 383

© Woodhead Publishing Limited, 2010

420

Index

silicon carbide, 176–9 silicon monoxide–carbon composite powder, 393 silicon nitride, 170 silicon nitride ceramics, 181–2 silver oxalate, 38 single wall carbon nanotubes, 384 Si–Zn–C composite, 385 Sn–Co–C composite, 382 Sn31Co28C41 composite, 392 SnO–B2O3–P2O5, 388, 396–7 sodium bicarbonate, 242 soft heating conditions, 356 soft mechanochemical synthesis, 125 future trends, 357 materials for lithium-ion batteries, 331–57 principles of solid inorganic compounds synthesis, 331–3 LiOH and Li2CO3 reactions with crystal hydrates and acidic salts, 348–54 LiOH reactions with anhydrous oxides, 333–43 LiOH reactions with solid hydroxides, 344–8 soft-mechanochemistry, 67 solid activation, 202 solid electrolyte interface, 363, 390 solid–liquid reaction ball milling as-milled products obtained, 152–8 binary alloy systems, 152–5 Fe1.3Sn powder particle SEM and TEM images, 154 milled product with addition of different PCAs, 158–9 nano-size intermetallic compound powder formation mechanism, 160 process control agents effect, 157–8 ternary alloy powders TEM images, 156 ternary alloy systems, 155–7 XRD results, 157

intermetallic compound powders production, 149–64 equipment schematic diagram, 151 experiment equipment and methods, 151–2 reaction mechanism, 158–4 comparison with high-energy ball milling, 161–3 main advantages vs mechanical alloying, 163 solid-state mechanochemical synthesis low-molecular organic compounds, 225–9 illustration, 227 possibilities in various types of reactions, 225–6, 228 solid reagents preliminary activation for subsequent heterophase reactions, 228–9 SPALTAN ’02, 84 spark plasma sintering process, 275–97 apparatuses and summary of consolidation of conditions investigated, 281 future trends, 296–7 mechanochemically activated powders, 281–2 MoSi2 system, 282–7 NbAl3 system, 293–5 schematic representation, 276 selected systems investigated, 278–9 starting powders used, 280 starting reactants mechanochemical activation, 277, 280–1 TiC–TiB2 system, 287–93 Spex 8000, 11, 175 Spex CertiPrep Inc., 48 Spex Mixer/Mill mod. 8000, 95 spin-orbit coupling, 319 spinel ferrites, 133, 134 spinel lithium manganese oxides, 366–9 materials prepared by mechanochemical processes, 369 spin–spin relaxation, 238 SPS see spark plasma sintering process srilankite, 311 stainless steel plates, 261–7

© Woodhead Publishing Limited, 2010

Index start material, 152 substrate temperature, 25 sulphides, 400 supersaturation, 234 surface atomic density, 58 tartaric acid, 242 temperature impulse, 22 ternary alloy systems, 155–7 thermal decomposition, 350 thermoreaction, 397 thymol, 82 TiC–TiB2 system, 281, 287–93 effect of milling time crystallite sizes in SPS samples, 292 SPS samples density, 291 sample shrinkage and temperature time profiles, 289 SPS sample when starting from unmilled and milled reactants, 290 XRD patterns of reactants as function of milling time, 288 TiN nanocoating, 270 characterisation, 271 cross-sectional view on Al and steel substrates, 268 interface between TiN coating and Al substrate, 270 original particles TEM image, 269 tin–graphite–silver composite, 382 Ti3SiC2, 179 titanate, 304–5 titania, 305, 316 DeNOx ability, 321 diffuse reflectance spectra, 315 light wavelength and photocatalytic ability relationship, 316 milling time effect on rutile phase mole fraction, 312 NOx gas decomposition activity, 318 particle size distribution and medium diameter, 313 S 2p XPS spectra of mixtures ground for 20 and 120 minutes, 319 titanium, 259, 267

421

titanium carbide, 173–6, 179 titanium coatings, 261, 262 titanium dioxide, 304–5, 304–26 acetaldehyde photo degradations ability, 323 and BaCO3 mixtures TEM micrographs, 74 TG and DTG profiles, 73 C 1s and N 1s XPS spectra of P25, 320 doped TiO2 total densities, 310 electron-hole separation mechanism during photocatalysis, 323 energy diagram and photoinduced charge transfer process, 307 F1s of SrF2-doped SrTiO3 sample and SrF2 powder XPS spectra, 325 malachite green oxalate sorption and degradation, 314 photocatalytic applications, 305–10 highly active visible-light induced TiO2 and its applications, 308–10 photocatalytic reaction mechanism, 306–8 structure, 305–6 polymorphs with distinct structure, 306 synthesis by mechanochemical doping, 310–25 binary composite photocatalyst, 322–3 carbon-doped and its photocatalytic activity, 317–18 co-doping effect on TiO2 photocatalysis, 320–2 effect on structure, sorption, and physicochemical properties, 311–14 fluorine doping and multipleelement co-doped photocatalysts, 324–5 nitrogen-doped and its photocatalytic activity, 314–17

© Woodhead Publishing Limited, 2010

422

Index

sulfur-doped and its photocatalytic activity, 318–20 TiO2-xNy/TiO2 composite acetaldehyde photodegradation ability, 324 washing process effect on photocatalytic activity, 326 t–P–T conditions, 10, 16–22 transition metals, 386 transmission electron microscopy, 205 tumbler milling, 160 tungsten, 179 turbo milling, 374 ultrasonic mechanical coating and armouring, 253 future trends, 272–3 key concepts, 253–5 representation, 254 ultrasonic vibration future trends, 272–3 key concepts, 253–5 mechanochemical plating and surface modification, 251–73 nanostructured coatings production, 267–72 nanocoatings characterisation, 270–2 surface morphology and coating nanostructure after treatment, 268–70 treatment conditions, 267

precoated substrates surface treatment, 260–7 LaPO4 coatings on stainless steel plates, 261–7 material systems and treatment conditions, 260–1 Ti coatings on Al plates, 261 surface modification and coating treatment, 255–60 Al surface armoured with Ti particles, 257 degree of coverage, 256–8 experimental, 255–6 particle sizes distribution, 258 set-up, 255 surface morphology and coating microstructure after treatment, 258–60 ultraviolet, 305 Uni-Ball-Mill, 170, 175 urea, 316, 325 Vegard’s law, 210 vibration mill, 130, 131, 139, 180 vibromills, 67 visible light irradiation, 305 Wako Pure Chemical Industry, 317, 319 wave theory, 14 Williamson–Hall method, 281, 284 zinc, 385 ZnFe2O4, 138 ZrO2 balls, 261

© Woodhead Publishing Limited, 2010