Nanostructured silicon-based powders and composites

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Author: A. P. Legrand | C. Senemaud

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Nanostructured Silicon-based Powders and Composites Legrand, A. P.; Senemaud, C. CRC Press 9780203297551 9780203222294 English Nanostructures, Silicon--Industrial applications. 2003 QC176.8.N35N36 2003eb 620.5 Nanostructures, Silicon--Industrial applications.

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Page i Nanostructured Silicon-based Powders and Composites

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Page iii Nanostructured Silicon-based Powders and Composites Edited by André P. Legrand and Christiane Sénémaud

LONDON AND NEW YORK

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Page iv First published 2003 by Taylor & Francis 11 New Fetter Lane, London EC4P 4EE Simultaneously published in the USA and Canada by Taylor & Francis Inc, 29 West 35th Street, New York, NY 10001 Taylor & Francis is an imprint of the Taylor & Francis Group This edition published in the Taylor & Francis e-Library, 2003. © 2003 Taylor & Francis All rights reserved. No part of this book may be reprinted or reproduced or utilised in any form or by any electronic, mechanical, or other means, now known or hereafter invented, including photocopying and recording, or in any information storage or retrieval system, without permission in writing from the publishers. Every effort has been made to ensure that the advice and information in this book is true and accurate at the time of going to press. However, neither the publisher nor the authors can accept any legal responsibility or liability for any errors or omissions that may be made. In the case of drug administration, any medical procedure or the use of technical equipment mentioned within this book, you are strongly advised to consult the manufacturer’s guidelines. 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 has been requested ISBN 0-203-22229-6 Master e-book ISBN ISBN 0-203-29755-5 (OEB Format) ISBN 0-415-30113-0 (Print Edition) Cover: A laser irradiation cell observed through a KCl window

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Christiane Sénémaud was a CNRS Directeur de Recherches until she passed away in 1998. Her area of expertise was the determination of the electronic structure of complex materials by X-ray photoelectrons and soft X-ray spectroscopies. She had published more than one hundred papers and communications in international meetings, and owing to her open mind and vast knowledge, she skilfully established fruitful cooperation between foreign and French laboratories. Moreover, she directed the theses of many young students and contributed to graduate education at the university, more precisely the ‘Complex Materials’ doctoral school. Finally, she was the leader of the CNRS research group that was the origin of this book. Her merit was recognised with the attribution of the Charles-Louis de Saulces de Freycinet Award by the French Academy of Sciences. Christiane Sénémaud was above all smart, discrete, fair, active, efficient and kind. For these reasons, the authors would like to dedicate this book to her memory.

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Page vii Contents List of contributors Foreword Preface Acknowledgements 1 Objectives and state-of-the-art of nanocomposites MICHEL CAUCHETIER AND ANDRÉ PIERRE LEGRAND 1.1 Introduction 1.2 Nanostructured ceramic elaboration 1.3 Nanostructured Si3N4/SiC materials 1.4 Analytical methods 1.5 Conclusion References 2 Laser synthesis of nanosized powders MICHEL CAUCHETIER, EMMANUEL MUSSET, MICHEL LUCE, NATHALIE HERLIN, XAVIER ARMAND AND MARTINE MAYNE 2.1Generalities on the synthesis of nanosized powders by laser pyrolysis 2.2 Synthesis of Si, SiC, Si3N4and SiC—Si3N4mixtures with SiH4and SiH3CH3as silicon precursors 2.3 Synthesis of silicon carbonitride (Si/C/N) powders using hexamethyldisilazane as Si precursor 2.4 Silicon carbonitride (Si/C/N) powders containing in situ the elements of the sintering aids (Al, Y): pre-mixed powders 2.5 Conclusion References 3 Thermal behaviour of as-formed silicon-based nanopowders

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Page viii 3.1 Thermal behaviour in a high-temperature graphite furnace EMMANUEL MUSSET, MARTINE MAYNE, MICHEL CAUCHETIER, XAVIER ARMAND, MICHEL LUCE AND NATHALIE HERLIN-BOIME 3.1.1 Introduction 3.1.2 Weight evolution of SiC, Si/C/N and Si/C/N/Al/Y/O nanopowders 3.1.3 Chemical evolution during annealing in a nitrogen atmosphere 3.1.4 Structural changes 3.1.5 Information about changes in the grain sizes 3.1.6 Conclusion References 3.2 Thermal reactivity of silicon-based nanopowders DJAMILA BAHLOUL-HOURLIER, BENOÎT DOUCEY, JEAN-LOUIS BESSON AND PAUL GOURSAT 3.2.1 Introduction 3.2.2 Results and discussions 3.2.3 Conclusion References 4 Structure of some nanopowders and micro/nano composites by transmission electron microscopy MARC MONTHIOUX, CATHERINE BÉRAUD AND NAHIDA EL HORR 4.1 As-prepared and heat-treated nanopowders 4.2 As-prepared sintered composites 4.3 TEM sample preparation prior to investigations and operating conditions References 5 From short- to long-range order 5.1 Chemical order studied by solid-state nuclear magnetic resonance ANDRÉ PIERRE LEGRAND, JEAN-BAPTISTE D’ESPINOSE DE LA CAILLERIE AND YOUSSEF EL KORTOBI 5.1.1 Nuclear magnetic methods 5.1.2 SiC powders analysis 5.1.3 Si/C/N powders analysis 5.1.4 Pre-alloyed Si/C/N/Al/Y nanopowders 5.1.5 Conclusion References

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Page ix 5.2 Paramagnetic defect states in silicon-based nanopowders ABDELHADI KASSIBA AND STÉPHANE CHARPENTIER 5.2.1Introduction 5.2.2 Theoretical background 5.2.3 Experimental details 5.2.4 Si/C powders 5.2.5 Assignment of the paramagnetic species in Si/C nanopowders 5.2.6 Silicon carbonitride (Si/C/N) powders References 5.3 Electronic structure studied by X-ray photoemission (XPS) and soft-X-ray (SXS) spectroscopies ADRIANA GHEORGHIU DE LA ROCQUE, GEORGES DUFOUR, FRANÇOIS TÉNÉGAL AND CHRISTIANE SÉNÉMAUD 5.3.1Principles 5.3.2 Experiments 5.3.3 Study by XPS–SXS of Si, Si/C, Si/N and Si/C/N laser synthetised nanopowders produced from gaseous precursors 5.3.4 Effect of heat treatment on the electronic structure of nanometric Si/C/N ex-HMDS and prealloyed Si/C/N/Al/Y powders studied by XPS spectroscopy References 5.4 Local order in Si/C/N nanopowders studied by X-ray absorption spectroscopy: study at the Si—K edge FRANÇOIS TÉNÉGAL, ANNE-MARIE FLANK, ADRIANA GHEORGHIU-DE LA ROCQUE AND CHRISTIANE SÉNÉMAUD 5.4.1 X-ray absorption principles and data analysis 5.4.2 Study by EXAFS at the Si—K edge of Si, SiN, SiC and Si/C/N laser-synthetised nanopowders produced from gaseous precursors 5.4.3 Short-range order determination in amorphous SiCN-HMDS nanopowders: evolution as a function of the C/N ratio and annealing treatment 5.4.4 Local structure of pre-alloyed Si/C/N/Al/Y nanopowders studied by XAS at the Al—K edge using fluorescence yield detection References 5.5 Neutron and X-ray diffraction applied to the study of the medium-range order in SiCN nanopowders FRANÇOIS TÉNÉGAL, BRIGITTE BOUCHET-FABRE, JEAN DIXMIER AND ROBERT BELLISSENT 5.5.1Diffraction applied to the study of disordered compounds

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Page x 5.5.2 Early stages of the crystallisation of Si/C/N-HMDS nanopowders studied by X-ray and neutron diffraction: Nano-ordered Si/C/N structures References 6 Conductivity and dielectric properties of Si/C nanopowders ABDELHADI KASSIBA 6.1Introduction 6.2 Material specifications 6.3 Experimental results and interpretation 6.4 Conclusion References 7 Processing and tailoring of Si/C/N-based nanocomposites PAUL GOURSAT, DJAMILA BAHLOUL-HOURLIER, BENOÎT DOUCEY AND MARTINE MAYNE 7.1Introduction 7.2 Sintering of silicon nitride with additives 7.3 Sintering of silicon carbide 7.4 Si3N4/SiC nanocomposites 7.5 Conclusion References 8 Mechanical properties JEAN-LOUIS BESSON 8.1 Room temperature properties 8.2 High-temperature properties 8.3 Conclusion References 9 Theoretical investigations of Si/C/N-based alloys FABIO FINOCCHI 9.1 Introduction 9.2 Simulations of Si/C and Si/N disordered alloys 9.3 Simulations of ternary Si/C/N systems References Index

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Page xi Contributors Xavier Armand Commissariat à l’Energie Atomique, Service des photons, atomes et molécules, Bat 522. 91191 Gif-sur-Yvette Cedex, France Robert Bellissent Laboratoire Léon Brillouin, CEA-Saclay, batiment 563, 91191 Gif-sur-Yvette Cedex, France Catherine Béraud Laboratoire de Physique de la Matière Condensée, INA,135 avenur de Raugueil, 31077 Toulouse Cedex, France Jean-Louis Besson SPCTS, UMR CNRS 6638, ENSCI, 47-73, avenue Albert Thomas, 87065 Limoges Cedex, France Brigitte Bouchet-Fabre Laboratoire d’Utilisation du Rayonnement Electromagnétique, Bat. 209, Centre Universitaire Paris-Sud, 91405 Orsay Cedex, France Michel Cauchetier Commissariat à l’Energie Atomique Service des photons, atomes et molécules, Bat 522. 91191 Gif-sur-Yvette Cedex, France Stéphane Charpentier Laboratoire de Physique de l’Etat Condensé, Université du Maine, Faculté des Sciences, Boulevard Olivier Messiaen, 72085 Le Mans Cedex, France Jean Dixmier Laboratoire de Physique du Solide, 1 place Aristide Briand, 92195 Meudon Cedex, France Benoît Doucey SPCTS, UMR CNRS 6638, Faculté desSciences, 123, avenue Albert Thomas, 87060 Limoges Cedex, France Georges Dufour UMR CNRS7614, Laboratoire de Chimie-Physique, Université Pierre et Marie Curie, 11 rue P. M. Curie, 75231 Paris Cedex 05, France Nahida El Horr CEMES, BP 4347, 29 rue Jeanne Marvig, 31055 Toulouse Cedex 4, France Youssef El Kortobi FRE CNRS2312, Laboratoire de Physique Quantique, Ecole Supérieure de Physique et de Chimie Industrielles de la Ville de Paris, 10, rue Vauquelin, 75231 Paris Cedex 05, France Fabio Finocchi UMR CNRS7588, Groupe de Physique des Solides, UPMC, 2, place Jussieu, 75251 Paris Cedex 05, France Anne-Marie Flank Laboratoire d’Utilisation du Rayonnement Electromagnétique, Bat. 209, Centre Universitaire Paris-Sud, 91405 Orsay Cedex, France

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Page xii Adriana Gheorghiu-de La Rocque UMR CNRS7614, Laboratoire de Chimie-Physique, Université Pierre et Marie Curie, 11, rue P.M. Curie, 75231 Paris Cedex 05, France Paul Goursat SPCTS, UMR CNRS 6638, Faculté des Sciences, 123, avenue Albert Thomas, 87060 Limoges Cedex, France Nathalie Herlin-Boime Commissariat à l’Energie Atomique, DSM/DRECAM/SPAM-Laboratoire Francis Perrin URA CNRS2453, Bat 522. 91191 Gif-sur-Yvette Cedex, France Djamila Bahloul-Hourlier SPCTS, UMR CNRS 6638, Faculté des Sciences, 123, avenue Albert Thomas, 87060 Limoges Cedex, France Abdelhadi Kassiba Laboratoire de Physique de l’Etat Condensé, Université du Maine, Faculté des Sciences, Boulevard Olivier Messiaen, 72085 Le Mans Cedex, France Jean-Baptiste d’Espinose de la Caillerie FRE CNRS2312, Laboratoire de Physique Quantique, Ecole Supérieure de Physique et de Chimie Industrielles de la Ville de Paris, 10, rue Vauquelin, 75231 Paris Cedex 05, France André Pierre Legrand FRE CNRS2312, Laboratoire de Physique Quantique, Ecole Supérieure de Physique et de Chimie Industrielles de la Ville de Paris, 10, rue Vauquelin, 75231 Paris Cedex 05, France Michel Luce Commissariat à l’Energie Atomique Service des photons, atomes et molécules, Bat 522. 91191 Gif-sur-Yvette Cedex, France Martine Mayne Commissariat à l’Energie Atomique, DSM/DRECAM/SPAM-Laboratoire Francis Perrin URA CNRS2453, Bat 522. 91191 Gif-sur-Yvette Cedex, France Marc Monthioux CEMES, BP 4347, 29 rue Jeanne Marvig, 31055 Toulouse Cedex 4, France Emmanuel Musset Commissariat à l’Energie Atomique, Service des photons, atomes et molécules, Bat 522. 91191 Gif-sur-Yvette Cedex, France Christiane Sénémaud UMR CNRS7614, Laboratoire de Chimie-Physique, Université Pierre et Marie Curie, 11, rue P.M. Curie, 75231 Paris Cedex 05, France François Ténégal Laboratoire d’Utilisation du Rayonnement Electromagnétique, Bat. 209, Centre Universitaire Paris-Sud, 91405 Orsay Cedex, France

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Page xiii Foreword The preparation of inorganic materials by thermolysis of preceramic compounds has become of substantial interest for the production of innovative materials in recent years. The general idea of this type of process route is that element-organic precursor molecules are constituted of structural units of the inorganic material aimed for by this process, thus providing novel paths of controlling the atomic array, composition and microstructure of materials. Of special interest is the manufacture of amorphous or nano-crystalline covalently bounded inorganics on the basis of silicon, boron, carbon and nitrogen, revealing a unique potential of properties not known from conventionally prepared materials. It is the intention of this book to provide a comprehensive survey on the French contributions to the synthesis of nano-sized powders and the production of nano-structured silicon-based composites by liquidphase sintering using nano-sized powders where interfaces play a decisive role for the evolution of properties. The infrared laser thermolysis of gaseous or liquid precursors is a highly versatile method well-fitted for the production of silicon-based amorphous or crystalline nanopowders. Due to the purity level and degree of dispersion, such powders are good candidates for the application in ceramics. Apart from ‘pure’ powders consisting of silicon, carbon and/or nitrogen, such ones containing homogeneously distributed alumina and yttria are typically used as sintering aids for silicon carbide and silicon nitride based materials. It has been proven that there is a strong influence of the carrier gases used for the synthesis laser process on residual hydrogen bonded to silicon, carbon and/or nitrogen providing some functionality during further thermal treatments. For the analysis of the amorphous and nano-crystalline structure of the as-thermolysed and heat treated powders, morphological and analytical transmission electron microscopy studies are decisive. Chemical ordering of the material such as powders in particular is analysed by nuclear magnetic resonance studies. Further analytical investigations are electron paramagnetic resonance, X-ray photoelectron spectroscopy, Xray absorption spectroscopy, X-ray and neutron diffraction. Furthermore, numerical simulations are used for the study of the chemical disorder and to propose structural models for the structure of the amorphous state of the materials. The content of the book is – even though very comprehensive – very detailed and strongly interdisciplinary. Therefore, the book is valuable for a broad readership in research and education.

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Page xiv The editors are well-known experts in the field of research covered by the book. They can be congratulated for a proper selection of authors and for providing an excellent piece of work. Fritz Aldinger Max-Planck-Institut für Metallforschung, Stuttgart, Germany June 2001

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Page xv Preface Physical properties of solids are greatly influenced by the arrangements of atoms present in the network. Nanostructured systems can offer improved and new properties compared to those of bulk materials. Consequently, in such materials, the influence of interfaces and interphases are of noticeable importance. Such interactions between the grains can lead to new physical and chemical properties. In recent years, an increasing interest has been devoted to nanostructured composites. This attention is largely due to exciting possible applications ranging from new catalysts to the preparation of nanocomposite ceramics, with significant improvements in their properties. Ceramic composites, obtained by dispersing nanoscale particles of one constituent into larger particles of a second constituent, have shown such improvements in mechanical properties (micro/nano composites). In particular, superplasticity has been obtained by mixing two equally fine constituents (nano/nano composites). Different modes of preparation have been used to make nanosized particles in large enough quantities. Among them, laser ablation of solid targets, plasma- or laser-induced reactions of gaseous mixtures have been developed. Other methods are related to evaporation– condensation, ball milling, self-propagating high temperature synthesis, sol–gel, spray drying of solutions, aerosol pyrolysis and chemical vapour decomposition at low pressure. Due to the large variety of such materials, comparative analysis is questionable. The originality of the book consists principally of focusing on a series of identical materials including preparation conditions, physical and chemical analysis, preparation and mechanical property determination of silicon-based nanocomposites. André P. Legrand ([email protected])

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Page xvi Acknowledgements We thank the Chemistry Department of the Centre National de la Recherche Scientifique (CNRS), which endorsed this work through a Groupe de Recherche (CNRS GdR 1168), for its interest and partial support, and more particularly its Director Jean-Claude Bernier and Associate Director Pierre Henri Dixneuf. The Délégation Générale pour l’Armement contributed to the development of this research through its financial support. We express here our sincere thanks.

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Page 1 1 Objectives and state-of-the-art of nanocomposites Michel Cauchetier and André Pierre Legrand 1.1 Introduction Processes to make new ceramic materials are many and depend on the final desirable structure and macroscopic shape to be obtained. So silicon carbide can be used as an abrasive, refractory, structural material or semiconductor with a negative temperature coefficient. Processes developed to produce such materials are extremely diversified. A process is selected based on the required properties and structural shape of the parts to be obtained. As an example, the preparation of carbon fibres, which needs specific polymer precursors for spinning followed by heat and oxidation treatments, had been adapted to the preparation of silicon carbide fibres. Other processes of transformation, descended from physics or chemistry domains are chemical vapour deposition, chemical liquid deposition, etc. (Figure 1.1). They are able to present improved mechanical, thermal or electrical properties, not solely depending on a peculiar chemical composition but on the arrangement of the crystalline phases or the size of the grains. Nanocomposites constitute a new class of extremely diversified materials which have appeared in the last decade.

Figure 1.1 Schematic representation of ceramic preparation showing the feedback process leading to a new ceramic material involving three main steps: synthesis of a precursor; analytical methods; and sintering and validation of the physical and chemical properties.

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Figure 1.2 Ceramic/ceramic micro–nano and nano–nano composites’ aspects (from Figures 7.8 and 7.20(b)). Sajgalik et al. (1996) proposed a classification into two main systems: • micro–nano composites in which the matrix is constituted by micrometric crystals embedded into a nanometric phase. Moreover, Niihara (1991) distinguishes three different configurations: intragranular, intergranular, intra/intergranular (Figure 1.2); • nano–nano composites in which the nanometric grains of the matrix and the second phase are spread uniformly (Figure 1.2). What makes a nanocomposite especially interesting is that at least one of its phases has dimensions in the nanometer range (10–100 nm). In this range, chemical and physical interactions have critical length scales and, if the nanoscale building block is made smaller than this one, the corresponding fundamental properties can be to changed. An example of this is light scattering, where by controlling the size of pores and nanocrystals in the range of 8 nm, Nanophase Technologies Corporation was able to produce transparent ceramics in the visible domain. Other enhanced properties have been developed (superplasticity, magnetoresistance, low temperature densification, enhanced and finer homogeneity, etc.). 1.2 Nanostructured ceramic elaboration Two main routes are generally used, depending on the densification rate and the properties to be obtained. 1.2.1 Amorphous ceramics obtained by solid state thermolysis of polymer precursors Pyrolysis, under a controlled atmosphere, of polymers, such as polysilazane, polycarbosilane, etc. is an efficient process for preparation of ceramics. According to Bill et al. (1999) ‘The in-situ crystallisation of these materials permits the preparation of nanocrystalline materials by a completely powder-free process. The structure and composition of the grain boundaries that originate from these crystallisation processes strongly depend on the molecular structure

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Page 3 of the initially used precursors.’ Such processes are particularly interesting in the preparation of fibres. 1.2.2 The powder-making processes To obtain nanocomposites it is necessary to prepare nanopowders so as to increase the reactivity of the different phases during the sintering process. This is due to a high surface/volume ratio (generally diameter < 100 nm). To obtain such powders, which have to present ideal characteristics (nanometric in size, spherical in shape, monodisperse, low agglomerated, high or controlled purity), it is necessary to develop new processes. These new methods must be innovative and need to be adapted to industrial production. Such recently developed processes deal with reactions such as • Solid phase: Evaporation–condensation, laser ablation, ball milling (or mechanical alloying), selfpropagating high-temperature synthesis (SHS); • Liquid phase: Sol–gel, spray-drying of solutions, aerosol pyrolysis; • Gaseous phase: Chemical vapour condensation at low pressure, plasma and laser synthesis. Some nanosized powders are now commercially available in the USA and in Japan: • Nanophase Technologies Corp. has developed a gas condensation process to produce different oxide powders (aluminium, titanium, zirconium oxides). • Nanodyne is focused on the production of cobalt–tungsten carbide, a ceramic–metal nanocomposite used to make tools for cutting and wear-resistant devices. • Ultram International produces ceramic and composite powders by a super-high-frequency plasma chemical process. • MarkeTech International proposes nanosized silicon carbide, silicon nitride, silicon carbonitride, boron carbide and silica powders obtained by a CO2-laser pyrolysis process. • Mitsubishi Mining & Cement Co. Ltd has prepared silicon/carbon/nitrogen nanocomposite powders from pyrolysis in a furnace of an organosilicon compound. A second problem which needs to be solved concerns the sintering for the preparation of complex shaped components so as to keep the initial size of the crystals. Similarly, the identification of the most efficient methods of characterisation, all along the different steps of preparation, is of significant importance. 1.3 Nanostructured Si3N4/SiC materials Different methods have been developed for the preparation of Si3N4/SiC nanocomposites by hot pressing sintering using: • Micrometric silicon nitride, previously recovered with carbon, so as to produce the carboreduction of silicon oxide present at the surface of the initial material: (Ishizaki and Yanai 1995; Watari 1989; Bahloul et al. 1993).

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Page 4 • Micrometric silicon nitride with nanometric Si/C/N or Si/C powders and sintering additives (Al2O3, Y2O3) (Niihara 1990). Nevertheless, the mixture, in a liquid medium, of such different sized powders containing so many different chemical groups on the surface, is difficult to obtain. This produces agglomerates and a heterogeneous distribution of the second phase in the matrix. • Nanometric Si/C/N or Si/C powders with sintering additives (Al2O3, Y 2O3) (Sasaki 1993; Kaiser 1996; Kennedy 1996). 1.4 Analytical methods Different methods can be used for the characterisation of initial to heat-treated and/or pressurised precursors such as: • Chemical analysis (CA) • Electron paramagnetic resonance (EPR) • Extended X-ray absorption fine structure (EXAFS) • Infrared spectroscopy (IR) • Soft X-ray spectroscopy (SXS) • Solid state nuclear magnetic resonance (NMR) • Specific surface area determination (SA) • Transmission electron microscopy and electron diffraction (TEM) • X-ray (XD) and neutron (ND) diffraction • X-ray photoelectron spectroscopy (XPS). Such methods are able to provide information of varied nature, depending on the origins of the physical phenomena involved (Table 1.1). Another important remark concerns the method of preparation of the sample. From the as-formed, heat treated and sintered sample, it is necessary to arrive at a condition adapted to the method of measurement. Crushing the material and filling, dispersing, spreading it in/on the sample holder are able to influence the results. This is why cross measurements are necessary. Table 1.1 Available information depending on the method used Method Surface Bulk Crystalline Amorphous Local or short order Average or long order Sample type CA × × Powder EPR × × Powder EXAFS × × × Powder IR × × × Powder ND × × × Powder NMR × × × × × Powder SXS × × Any TEM × × × Nanopowder XD × × × Powder XPS × × × × Any SA × × Any Note × Indicates the dominant domain of expertise of the method.

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Page 5 Some examples shown in this book demonstrate such phenomena: • X-ray, TEM and NMR. Figure 2.3, section on Si/C-based powders in Chapter 4 and Figure 5.1.6: α-β-SiC and amorphous detected in different proportions; • NMR and TEM. Figures 4.65, 4.66 and 5.1.27: Si3N4 detected in the bulk and not into grains; • XPS and NMR. Figures 5.3.10 and 5.1.10: SiO2 detected on the surface and not in the bulk; where the different phases are not evaluated with the same efficiency. 1.5 Conclusion Although till now many studies on that subject have been published, results are contradictory and fragmentary. No interdisciplinary studies, till now, have been done on the silicon nitride– silicon carbide system. A better knowledge of the behaviour and potential of such materials is needed. Elaboration of high-temperature ductile nanocomposites is one of the objectives of this work. Consequently, systematic studies and analyses of the preparation proceed from step by step. Chemical bonding, nanostructure, microstructure and monolith material properties are determined. Such an approach allows us to obtain original results for each step of the preparation and to establish relations between the synthesis of powders, their sintering aptitude, the development of the microstructure and the properties obtained. The spiral in Figure 1.1 represents this symbolically. References Bahloul, D., Pereira, M. and Goursat, P. (1993) J. Am. Ceram. Soc., 76(5), 1156–68. Bill, J., Wakai, F. and Aldinger, F. (1999) Precursor-Derived Ceramics. Wiley-VCH. Ishizaki, K. and Yanai, T. (1995) Silic. Indus., 7–8, 215. Kaiser, A. (1996) Silic. Indus., 5–6, 215. Kennedy, T. (1996) Silic. Indus., 9–10, 201. Niihara, K. (1990) J. Mater. Sci. Lett., 10, 112. Niihara, K. (1991) The Centennial Memorial Issue of the Ceramic Society of Japan, 99(10), 974. Sajgalik, P., Dusza, J., Hofer, F., Warbichler, P., Reece, M., Boden, G. and Kozankova, J. (1996) J. Mater. Sci. Lett., 15, 72. Sasaki, G. (1993) Mater. Res. Soc. Symp., 287, 335. Watari, K. (1989) Mater. Sci. Eng., A109, 89.

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Page 6 2 Laser synthesis of nanosized powders Michel Cauchetier, Emmanuel Musset, Michel Luce, Nathalie Herlin, Xavier Armand and Martine Mayne 2.1 Generalities on the synthesis of nanosized powders by laser pyrolysis 2.1.1 The laser pyrolysis process Infrared laser pyrolysis (IRLP) is a highly versatile method for the production of a wide range of nanopowders including SiC, SiCN, SiCO, BN and carbon (Haggerty and Cannon 1981; Cauchetier et al. 1988, 1994; Herlin et al. 1996; Boulanger et al. 1995; Ehbrecht et al. 1993; Voicu et al. 1996). A summary of the syntheses performed so far for materials applications is given in Table 2.1 and completes the review appearing in Knudsen (1997). In the IRLP process, the reactants are heated by IR laser radiation and decompose, causing clusters to nucleate and grow rapidly. This process is inherently very clean because homogeneous nucleation occurs in a well-defined reaction zone without interaction with chamber walls. The small reaction volume and the ability to maintain steep temperature gradients (106 °C/s) allow precise control of the nucleation rate, growth rate and residence times. The physical and chemical properties of the particles can be controlled by changing the molecular precursors and the synthesis parameters (laser power, pressure, etc.). The resulting powders are very fine, spherical, extremely pure, more or less agglomerated and nearly monodispersed in size. The mean particle size can be adjusted from 10 to 100 nm. All these characteristics (i.e. nanometric size (


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Page 7 Table 2.1 CO2 laser synthesized nanometric powders Powders Chemical Si SiH4

systems References Haggerty and Cannon 1981; Cauchetier et al. 1987 SiC SiH4 + C2H4 Sumaya et al. 1985; Cauchetier et al. 1987; Förster et al. 1991 SiH4 + C2H2 Cauchetier et al. 1987; Curcio et al. 1989; Fantoni et al. 1990; Tougne et al. 1993 SiH4 + CxHy with x = 4 Cauchetier et al. 1988 SiH2Cl2 + C2H4 Suzuki et al. 1992 (CH3)2Si(C2H5O)2 vapour Li et al. 1994a Si3N4 SiH4 + NH3 Haggerty and Cannon 1981; Kizaki et al. 1985; Symons and Danforth 1987; Buerki et al. 1990 SiH2Cl2 + NH3 Bauer et al. 1989 SiC + Si3N4 or Si/C/N [(CH3)3Si]2NH vapour Rice 1986; Li et al. 1994b SiH4 + CH3NH2 + NH3 Cauchetier et al. 1989; Cauchetier et al. 1991 SiH4 + (CH3)2NH + NH3 Alexandrescu et al. 1991; Borsella et al. 1992 SiH4 + C2H4 + NH3 Suzuki et al. 1993 (CH3SiHNH)x with x = 3 or 4 (aerosol) Gonsalves et al. 1992 [(CH3)3Si]2NH (aerosol) Cauchetier et al. 1994; Herlin et al. 1994; Musset et al. 1997 SiO2 Si(C2H5O)4 (aerosol) Luce et al. 1994 SiO2 + SiC + C or Si/C/ (RO)4−xSi(R′)x with 0 = x = 3, R = Armand et al. 1995; Martinengo et al. O C2H5O, R′ = CH3 and [(CH3)3Si]2O 1996; Herlin et al. 1996; Fusil et al. 1997 (aerosol) BN BCl3 + NH3 Luce et al. 1993; Baraton et al. 1994; Willaime et al. 1995 B4C BCl3 + CH4 + H2 Knudsen 1987a BCl3 + C2H2 + H2 Luce et al. 1993 TiB2 TiCl4 + B2H6 Casey and Haggerty 1987a TiCl4 + BCl3 + H2 Knudsen 1987b ZrB2 Zr(BH4)4 Cauchetier et al. 1987; Rice and Woodin 1988 TiO2 Ti[OCH(CH3)2]4 Casey and Haggerty 1987b; Rice 1987; Curcio et al. 1991 TiC TiCl4 + C2H4 Alexandrescu et al. 1997 Al2O3 Al(CH3)3 + N2O + C2H4 Borsella et al. 1993 WC WF6 + C2H2 + H2 + SF6 Bourgeois et al. 1995 electronic grade remains too expensive (Cauchetier et al. 1994) for use in large quantities (>1 ton/year) even if energy and laser equipment costs are low compared to the raw material costs. In the case of Si3N4 laser synthesis, a study of the economics shows that a price for SiH4 less than one-fourth of its current price is needed to compete with a classic method of synthesis (Schoenung 1991). Alternative methods have been proposed with the substitution

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Page 8 of chlorosilanes (Bauer et al. 1989; Suzuki et al. 1992) or organosilicon compounds (Rice 1986; Gonsalves et al. 1992; Cauchetier et al. 1994) for silane. Technical improvements in two different fields have enhanced the versatility of the method. First, new introduction systems for the silicon (Si) precursors have been tested. Liquid organosilicon compounds are introduced into the laser beam in the aerosol form, fine liquid droplets being obtained by ultrasonic nebulisation (Gonsalves et al. 1992; Cauchetier et al. 1994). Second, the use of high-power tunable continuous wave (c.w.) CO2 lasers, now available, has increased the area of applications of the process (Luce et al. 1994). In this chapter, the synthesis of Si-based powders by two methods, using gaseous or liquid precursors, will be presented: • SiH4 as gaseous precursor for the synthesis of Si, SiC, Si3N4 and SiC–Si3N4 mixture powders and also methylsilane, SiH3CH3, for the synthesis of SiC powder, • hexamethyldisilazane [(CH3)3Si]2NH or HMDS as liquid precursor for the synthesis of Si/C/N (silicon carbonitride) powders; mixture of organosilicon precursors for the synthesis of Si/C/N powders containing the elements of sintering additives. 2.2 Synthesis of Si, SiC, Si3N4 and SiC Si3N4 mixtures with SiH4 and SiH3CH3 as silicon precursors 2.2.1 Experimental setup The experimental device is presented in Figure 2.1 (Cauchetier et al. 1991). The unfocused Gaussian beam (diameter = 12 mm) of a high-power c.w. CO2 laser (CILASCI1000, Marcoussis, France) enters the reaction cell through a KCl window and crosses the path of the gaseous flow of the reactants injected through an inlet capillary (inner diameter = 2mm).

Figure 2.1 Schematic laser irradiation cell for gaseous precursors. (Reprinted from Cauchetier, M., Croix, O., Luce, M., Baraton, M. I., Merle, T. and Quintard, P., (1991) Journal of the European Ceramic Society, Nanometric Si/C/N Composite Powder: Laser Synthesis and IR Characterisation, 8, 215, with permission from Elsevier Science.)

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Page 9 A resonance effect between the emission line of the laser at 10.6 μm and an infrared absorption band of one of the reactant gases (silane, SiH4 or methylsilane, CH3SiH3) causes the reaction to occur. This laser-driven reaction leads to high temperatures (up to 1800°C) and a bright flame. An argon flow prevents powder deposition on the windows and guides the products into the collection chamber. The powders are collected and then stored in a glove-box under argon atmosphere to avoid contamination by air or water vapour. 2.2.2 Synthesis results Table 2.2 summarises the synthesis conditions of the different powder samples used in different parts of this book (Si, SiC, Si3N4 and mixtures of SiC/Si3N4) and also presents first characterisation results. For all the experiments reported here, the cell pressure was kept constant at 1 atmosphere (105 Pa). In Table 2.2, most of the results concern SiC and Si/C/N synthesis, but some results shows the synthesis of Si from pure SiH4 (Si50), Si3N4 from a mixture of SiH4 and NH3 (SiN7). Structural investigations on both these powders are presented in Chapter 5. SiC powders have been obtained from methylsilane (SiC151, SiC152) or from mixtures of SiH4 and C2H2. In the run, SiC171 diborane was added in the silane/acetylene mixture, as a boron source used for sintering aid in densification tests (Croix et al. 1991), and after pressureless sintering at 2000°C, densities up to 0.95 were achieved. The last samples presented in Table 2.2 are Si/C/N samples, obtained from a mixture of SiH4, CH3NH2, and, in some cases, NH3 has been added to the reactive mixture. In two runs (Si50, SiC163), an inert gas (Ar, He) has been added to the reactive gases in order to increase the linear velocity. After comparing the chemical composition of the gas phase with the chemical composition of the powder, the C/Si and C/N ratios were obtained (see Table 2.2). In some cases, the flame temperature was measured using an optical pyrometer. Specific surface area (S)was determined from Brunauer–Emmet– Teller (BET) measurements and the equivalent diameter (D) was calculated from these measurements. The time of residence in the laser beam is calculated from the gas flow rates (2 mm nozzle diameter, 12 mm laser beam diameter) assuming room temperature. From the results presented in Table 2.2, some information can be obtained on the effect of experimental parameters. Laser power Increasing the laser power, and keeping all the other parameters constant, leads to an increase in temperature, which is expected to have an influence on the powders produced. In experiments SiC151 and SiC152, the laser power has been decreased by a factor of three (600/220 W) and the measured temperature decreases from 1530°C to 1150°C. The same crystalline structure is identified by X-ray diffraction (XRD). Therefore, in this temperature range, the flame temperature does not influence the structure of the powder. In contrast, BET measurements indicate a decrease in the powder size as the temperature decreases. Thus, the temperature has a direct influence on the kinetics of the powder growth. Flow rates In runs SiC163–177, the flow rates have been changed keeping the laser power constant. Increasing the flow rates leads to a decrease in the residence time as shown in Table 2.2. Figure 2.2 presents the variation in the particle size of SiC powders with the residence time for samples SiC163–177. Figure 2.2 also presents results (broken line) obtained five years before

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Page 10 Table 2.2 Gas phase synthesis conditions and characterisation results of the different samples Samples Flow rates (cm3/min) Atomic Laser Temperature (° Residence BET XRD ratio (gas power C) time (ms) results phase) (W) SiH4 C2H2 He, CH3NH2NH3 C/N C/Si S D Ar (m2/ (nm) g) Si50 120 0 480 0 0 600 930 4.3 29 89 Si (Ar) SiC151 CH3SiH3 : 600 1530 13.0 34 55 β-SiC 200 SiC152 CH3SiH3 : 220 1150 13.0 52 36 β-SiC 200 SiC163 540 300 1000 0 0 1.1 600 1.2 117 16 (He) SiC171a 1080 600 0 0 0 1.1 600 1.3 135 14 SiC173 200 110 0 0 0 1.1 600 7.3 61 31 SiC174 350 192 0 0 0 1.1 600 4.2 74 25 SiC177 120 66 0 0 0 1.1 600 12.1 36 52 SiC212 596 306 0 0 0 1.0 640 2.5 70 27 SiN7 120 0 0 0 480 600 1250 4.3 71 26 SiCN12 340 0 0 200 0 600 1550 4.8 47 40 α, βSi3N4, β-SiC SiCN27 280 0 0 70 250 600 1615 4.3 26 72 α, βSi3N4 SiCN29 360 0 0 90 320 0.22 0.25 400 2.9 40 49 SiCN35 350 0 0 200 0 1 0.57 600 4.1 48 39 Note a 40 cm3 /min of diborane B2H6 added as boron source which acts as sintering aid for the densification of SiC.

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Figure 2.2 SiC particle size v. the residence time in the laser beam (laser power, 600 W; cell pressure, 1 atm). Similar results have already been observed previously (– – –) with a laser power of 640 W. (Reprinted from Tougne et al. (1993) Diamond and Related Materials, Evolution of the Structure of Ultrafine SiC-Laserformed Powders with Synthesis Conditions, 2, 486–490, with permission from Elsevier Science.) the present results, with a higher laser power (640 W) (Tougne et al. 1993). It must be noted that the value obtained for the SiC212 (640 W) experiment (not plotted) is in good agreement with these early results. Figure 2.2 clearly shows a linear decrease in the size of the particles as the residence time decreases. Figure 2.3 presents the XRD diagrams of the SiC163–177 samples. The lines become sharper and the full width at half maximum (FWHM) decreases as residence time increases, which indicates an increasing size of crystallites, parallel to the increasing particle size (BET measurements). TEM observations also confirm the BET and XRD results. Sample SiC163 (Figure 2.4(a)) presents fine spherical particles with a mean diameter of 20 nm (BET measurement: 16 nm) and with a unimodal distribution. They are not fully crystallised and appear as fine crystallites in an amorphous matrix. It clearly shows that the size of the particles (BET) and the size of the crystallites (XRD) must be distinguished. Sample SiC177 (Figure 2.4(b)) also presents fine spherical particles with an increased mean diameter of 60 nm (BET measurement: 52 nm). They are well crystallised with some black streaks. In correlation with the chemical composition C/Si = 1.1, some fine filaments of turbostratic carbon (C) can be noticed. In conclusion, it must be noted that both laser power and flow rates have a noticeable effect on the size of the particles. A good control of these experimental parameters allows a good control of the size and the structure of powders. Some other points which are relevant for materials applications must be noted: • The amount of synthesised powders is significant. For example, in the run SiC171, a production rate of about 100 g/h indicates that the process can be easily scalable. The

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Figure 2.3 XRD patterns of the as-synthesised SiC powders (arbitrary units (a.u.)). (Reprinted from Tougne et al. (1993) Diamond and Related Materials, Evolution of the Structure of Ultrafine SiC-Laser-formed Powders with Synthesis Conditions, 2, 486–490, with permission from Elsevier Science.) duration of the experiment can be 3 h or more in order to collect powder quantities in the 150–200 g range for the elaboration of materials. • Also, a very good reproducibility of the experiments can be achieved; for example, Table 2.2 shows BET results for the runs SiCN12 and SiCN35 undertaken after a delay of 4 years. • Another interesting point is the good yield of the reaction. For example, in the SiC212 run, and supposing 100% conversion of SiH4 and C2H2 to the solid phase, the maximum calculated production rate is 64 g of SiC per hour and the experimental result is 59 g/h – this reaction yield (almost 100%) is very good. 2.2.3 Chemical analyses of the as-formed powders The chemical composition in stoichiometric compounds (Si3N4, SiC, SiO2, Si, C) has been calculated from elemental analysis results, with the classical assumption that all O is in the form of SiO2, then all N forms Si3N4 and the remaining Si is in the form of SiC. Free C (or Si) is the difference between total C (or Si) and C (or Si) bonded to Si (or C) in SiC. Such a calculation is useful to evaluate the composition of future sintered materials, or to compare several powders but it is obvious that it is only an approximation, especially when the powders are not fully crystallised.

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Figure 2.4 TEMs of two as-formed powders: (a) SiC163; (b) SiC177. SiC powders In order to avoid the presence of free Si in the powders which makes the powder very sensitive to oxygen contamination, most of the SiC samples were synthesised with a gaseous mixture containing excess C (C/Si = 1.1 in the gaseous mixture). In Table 2.3, two examples show that this ratio is also found in the powders produced. In agreement with the chemical composition, XRD patterns (similar to Figure 2.3) correspond to the β-SiC phase. The O and free C contents are low and correspond to the values usually encountered for commercial powders. Table 2.3 also presents the chemical composition calculated in stoichiometric compounds which indicates, in good agreement with the XRD pattern, that the powders are mostly composed of SiC. Si/C/N powders For the synthesis of Si/C/N composite powders, two gaseous mixtures were investigated: a binary mixture SiH4 + CH3NH2 (SiCN35) and a ternary mixture: SiH4 + CH3NH2 + NH3

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< previous page Page 14 Table 2.3 Chemical characterisations of gas Powder Samples Chemical analysis (wt%) Si C N SiC173 68.3 30.2 0 SiC212 68.3 30.3 0 SiCN29 58.3 6.5 34.6 SiCN35 66.9 13.6 18.2

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phase synthesised powders

Atomic ratio Chemical composition (wt%) O C/Si C/N Si3N4 SiO2 SiC 0.9 1.03 0 1.5 96 1.4 1.03 0 2.7 95.7 0.6 0.27 0.22 86.6 1.1 8.3 2.3 0.47 0.87 45.1 4.3 45.1

C 2.0 1.6 4.0 0

Si

0 0 0 5.5

Figure 2.5 XRD patterns of as-formed SiCN nanopowders used in material elaboration. (Reprinted from Mayne 1997.) (SiCN29) (Table 2.2). In the case of SiCN29, the presence of NH3 is known to make the nitriding introduce an increase in the nitrogen (N) content. For SiCN35 a similar content in C and N is expected in the resulting powder due to the C/N initial atomic ratio of 1 in the formula of monomethylamine. Tables 2.2 and 2.3 show a very good correlation between C/Si and C/N ratios in the gas phase and in the powder for SiCN29, but the correlation is not so good for SiCN35 probably due to evolution of carbonaceous gas of C in the gas phase during the synthesis. Chemical composition calculations show that the powder SiCN35, containing free Si is the most sensitive to oxygen contamination. XRD patterns (Mayne 1997) are reported in Figure 2.5 and confirm the calculated compositions. The SiCN29 sample contains mainly α- and β-Si3N4; the quantity of SiC is too small and is not detectable. The SiCN35 sample presents the characteristic lines of Si and β-SiC and a broad peak between 2θ = 30° and 40°, which can be attributed to an amorphous phase of silicon carbonitride. From the results presented in this section, it appears that the decomposition in stoichiometric compounds is a good representation of the powder and also that the chemical composition of the powders produced is most often well controlled by the chemical composition of the gaseous reactants. Together with the possibility of controlling the size and the structure of

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Page 15 the powder, it makes this synthesis technique attractive. But, as explained before, in order to increase the safety and decrease the cost, it has been applied to the liquid precursors containing Si, C and N elements. 2.3 Synthesis of silicon carbonitride (Si/C/N) powders using hexamethyldisilazane as Si precursor 2.3.1 Experimental Choice of the precursor As noted before, investigations based on costs of various silicon precursors for the delivery of large quantities have shown that silane remains too expensive for use in ceramic powder synthesis (Cauchetier et al. 1994). In our case [(CH3)3Si]2NH or HMDS with a Si cost very low compared to the Si cost of SiH4 (43 US$/kgSi in HMDS and 343 US$/kgSi in SiH4)was chosen as a model of the liquid precursor in order to demonstrate the possibility of using liquid wastes of the organosilicon compound chemistry for the synthesis of nanometric Sibased powders. Indeed, HMDS has a strong IR absorption band at 10.6 μm due to the Si N bond, which was, therefore selected. It is a common liquid Si compound (boiling point of 125°C, density of 0.76, and a viscosity of 0.69 cP at 25°C) largely used for the silylation of a wide range of functional groups and it was the precursor used by Niihara and co-workers in the synthesis of Si/C/N powders leading to non-oxide superplastic ceramics (Wakai et al. 1990). The first experiments using HMDS in the vapour phase as Si precursor for Si/C/N powder synthesis by CO2-laser pyrolysis were performed by Rice (1986). A conversion efficiency (liquid → powder) of about 60% has been obtained but with low flow rates of HMDS(7 g/h) and with a low-powered CO2 laser (135 W). Higher production rates have been attained (80–120 g/h) when the laser power has been increased to the 500–1000 W range (Li et al. 1994b). The experimental device The apparatus is shown in Figure 2.6. Liquid HMDS is placed in a special glass jar containing a piezoelectric transducer (Pyrosol 7901 type from RBI, Meylan, France, apparatus developed in collaboration with the CEA (Spitz and Viguié 1970)). The focusing of the intense beam of ultrasonic energy delivered by the transducer near the surface of the liquid yields uniform droplets whose diameter d is given by

where σ and ρ are the surface tension and the density of the liquid, respectively, and f the frequency of the transducer (here 850 kHz). In order to increase the mass of the displaced liquid, the glass jar is heated near 90–100°C with a heating ribbon. The aerosol droplets and vapour are injected into an irradiation cell very similar to those presented in Figure 2.1. The reaction cell is maintained at a regulated pressure of 105 Pa using a flow of argon or an argon–ammonia mixture through a glass inlet tubing (inner diameter of 13 mm). The laser beam diameter is increased from 12 to 24 mm with a beam expander in order to cover the reactant flow entirely. The initial laser power is in the range 480–520 W.

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Figure 2.6 Schematic of the aerosol generator with the irradiation cell and the powder collector. 2.3.2 Synthesis results Table 2.4 compares the influence of two parameters, the nature of the carrier gas (argon and nitrogen) and heating of the liquid precursor (near 100°C), on the synthesis results. It can be seen immediately that the yield of the reaction (expressed as the ratio between the mass of powders produced and the mass of displaced liquid), does not appear to be as good as in the gas phase. Both the amount of displaced liquid and production rate increase when the precursor is heated. The effect of the heating is to combine the enrichment of the gaseous phase in HMDS vapour and decrease the viscosity of the liquid. In order to work with significant amounts of powders, most of the experiments were carried out with liquid heating. The yield seems independent of the heating of the precursor, but is strongly dependent on the nature of the gas. It is 46–49% for Ar carrier gas flow and 38–39% for N2 carrier gas flow.

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Page 17 2.3.3 Effect of experimental parameters Effect of carrier gas (Ar or N2) Table 2.4 shows that the chemical composition of the powder does not depend on the nature of the carrier gas flow (Ar or N2) and the C/N atomic ratio remains stable in the 2.5–2.9 range. The activity of atomic N (N2 ↔ 2N) is low at the measured flame temperature: 1260±20°C. Table 2.4 also shows that all powders are sensitive to pollution by oxygen, but the powders obtained from the heated precursor seem the most sensitive to pollution. It must be noted that the chemical analysis presented in this section is not complete due to the presence of hydrogen in the powder and due to the uncertainty of the measurements. Effect of chemical evolution of the reactive mixture Table 2.5 presents the synthesis parameters of all the samples studied in the different parts of this book. Only samples HMDS40–45 where the effect of chemical composition has been Table 2.4 Effect of the carrier gas nature and precursor heating Run Carrier gas Displaced Powder Yield SBET Chemical analysis (wt C/N flow rate (cm3/ liquid (cm3/h) production rate (wt%) (m2/g) %) (Atomic min) (g/h) ratio) Ar N2 Si C N O HMDS28a 2430 — 108 38 46 88 49.8 30.3 12.3 5.6 2.9 HMDS36b 2500 — 164 61 49 82 45.8 28.0 11.7 9.7 2.8 HMDS38b — 2500 166 48 38 125 47.2 28.4 12.8 11.6 2.6 HMDS39a — 2500 95 28 39 79 49.9 29.7 14.1 5.2 2.5 Notes a Liquid precursor at room temperature. b Liquid precursor heated near 100°C. Table 2.5 Synthesis conditions and chemical analysis of HMDS samples (chemical analyses were always performed after several weeks of exposure to air) Run Carrier gas flow rate Displaced Powder Yield SBET Chemical analysis (wt C/N (cm3/mn) liquid (cm3/ production (wt (m2/ %) (Atomic Ar N2 Total h) rate (g/h) %) g) Si C N O ratio) HMDS40 1935 205 2140 182 59 43 110 46.7 25.3 21.8 6.1 1.35 HMDS41 1770 410 2180 160 61 50 93 46.1 20.5 25.2 6.7 0.95 HMDS42 1570 600 2170 175 75 56 98 44.4 18.5 29.7 7.7 0.73 HMDS43 1370 800 2170 179 74 53 95 48.3 15.5 26.8 8.0 0.67 HMDS44 1150 1040 2190 210 81 51 95 47.1 13.1 26.7 10.9 0.57 HMDS45 830 1346 2180 167 66 52 106 47.6 8.3 28.9 13.4 0.34 HMDS35 1600 1000 2600 87 21 32 144 47.3 6.6 33.4 9.8 0.23 HMDS66 1920 1920 114 32 36 66 51.2 31.9 15.1 1.8 HMDS67a 1440 1920 115 31 35 66 52.0 30.8 14.8 2.4 HMDS73a 862 207 1920 56 14 32 115 51.9 12.6 25.3 10.1 Note a H2 in the carrier gas

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Figure 2.7 C/N atomic ratio in as-formed HMDS powders v. C/N atomic ratio in the precursors (liquid + ammonia). (Reprinted from Musset et al. 1997 with the permission of the Belgian Ceramic Society, Mons, Belgium.) closely studied will be commented on in this section; the main part of the results appear in the PhD thesis of Emmanuel Musset (Musset 1995). For this set of six experiments, the liquid temperature is about 100°C, the total flow rate of the carrier gas mixture remains constant and the corresponding residence time of the precursor in the laser beam is about 0.1 ms. Table 2.5 shows the influence of the addition of ammonia in the argon used as carrier gas, namely, on the chemical composition of the resulting powders (samples HMDS40–45). Figure 2.7 shows a good correlation between the C/N ratio in the as-formed powder and in the reactive mixture; the C/N atomic ratio decreases from 1.35 to 0.36 when the ammonia volume ratio increases from 10% to 60%. This correlation confirms the possibility of controlling the chemical composition of the powders. Table 2.5 shows that the contamination by oxygen becomes more and more important when the N content in the powder increases, because NH bonds are easily hydrolysed. 2.3.4 First characterisations Specific surface area The specific surface area is always high, about 100 m2/g and taking into account the mean value of the density i.e. 2.00 (pycnometry measurement), the corresponding mean value for the diameter is 30 nm, in agreement with TEM observations.

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Figure 2.8 IR spectra of the as-formed HMDS powders. (Reprinted from Musset 1995.) Infrared spectroscopy IR spectra of the samples HMDS40–45 are shown in Figure 2.8. They present a broad absorption band between 750 and 1250 cm−1 with a maximum at 920–940 cm−1 due to a Si/C/N amorphous phase, corresponding to a flat XRD diagram. A weak band at 550 cm−1 is due to the Si N bond. Other bands are related to hydrogenous species: N—H at 1180 cm−1, Si—H at 2050 cm−1 and Si—CH3 at 1250 cm−1, indicating an incomplete pyrolysis of the precursor or absorption of volatile species obtained from the pyrolysis. The presence of hydrogen is in good agreement with the chemical analysis, which is often less than 100% (Table 2.4). The results presented in this section show that it is possible to produce Si/C/N nanopowders with controlled chemical composition from a liquid precursor. The powders are obtained in significant quantities for materials applications. The main differences compared to powders obtained from the gas phase are the amorphous character and high hydrogen content of the powders obtained from the liquid phase. 2.4 Silicon carbonitride (Si/C/N) powders containing in situ the elements of the sintering aids (Al, Y): pre-mixed powders 2.4.1 Introduction: elaboration of SiC–Si3N4nanocomposite materials Materials in the SiC/Si3N4 system are elaborated by heat treatment of powders in the presence of liquidforming sintering aids such as Y2O3, MgO, Al2O3, etc. singly or in combination. Such processes need a previous mixing step of all powders having different natures and sometimes different sizes. Recently, some studies report the incorporation of metal elements

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Page 20 (Al, Y, etc.), necessary for the sintering step, into silicon nitride or carbonitride systems, most of them being focused on chemical processing routes using organometallic precursors (Soraru et al. 1992; Iwamoto et al. 1998; Koyama et al. 1998). A few other publications concern physico-chemical processes: laser irradiation of gaseous mixtures (SiH4, NH3, (CH3)3Al) allows the formation of nanosized powders with Al in a Si3N4 matrix (Borsella et al. 1993b). Aerosol of Si/C/N particles with yttrium and aluminium compounds were prepared under dual irradiation of CO2 and excimer laser beams of a SiH4, C2H4 and NH3 mixture, aluminium being incorporated in the form of alumina powder in the reactant flow and yttrium being introduced as a low volatile organometallic compound in the reaction zone (Yamada et al. 1997). 2.4.2 Synthesis of pre-mixed powders In order to prepare Si/C/N powders containing uniformly dispersed sintering aids investigations were undertaken to set up a laser spray synthesis process of Si/C/N/O/Al(Y) nanopowders using solutions of a liquid organosilicon precursor – HMDS – with solid metallic alcoxides, namely aluminium and yttrium isopropoxides [(CH3)2CHO]3Al and [(CH3)2CHO]3Y. Table 2.6 summarises the synthesis conditions of 14 runs which can be divided into 3 groups: • Group I includes experiments carried out with solutions having different Al (+Y)/Si atomic ratio. • Group II includes experiments carried out with solutions having the same Al (+Y)/Si atomic ratio. All runs are realised with a binary mixture of carrier gases (Ar + NH3). • The third is composed of experiments made with the same solutions as in group II, but the carrier gas is composed of Ar+NH3 +H2 in order to see if the chemical composition of the powders is modified by the presence of H2. Some changes occur compared to the experiments described above in Section 2.3. The initial laser power is in the range 300–330 W, but, here, with an unfocused laser beam (diameter = 13 mm) which covers the aerosol flow sufficiently. For comparison, an experiment with pure HMDS carried out under similar conditions is also reported (HMDS66). 2.4.3 Effect of experimental parameters Effect of the solutions of precursors In Table 2.6, most of the experiments were carried out with metallic isopropoxides in HMDS but, in two cases (HSAl04 and HSAlY05), isopropanol was added to the reactive solution in order to increase the solubility of the alcoxides and to decrease the viscosity of the solutions. In these experiments, more liquid is displaced than in other experiments but the production rate remains at the same level (≈30 g/h). Also, the amount of oxygen incorporated in the as-formed powders is very high (>10%). In conclusion, isopropanol does not improve the production or quality of the powders and this solution was not developed. Precursor solutions are characterised by their atomic ratios: [Al(+Y)]/Si. For the first experiments, the Al/Si ratio in the precursor solution was adjusted to 0.069 (e.g. HSAl03); this value corresponds to 6 wt% Al2O3 content in a final SiC/Si3N4 material according to the different oxide contents reported in the literature from about 10 (Kennedy et al. 1996) to 1 wt%

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Page 21 Table 2.6 Synthesis conditions and results of pre-mixed Si/C/N powders Run Carrier gas flow rates Displaced Powder SBET Chemicalanalysis (wt Atomic (cm3/min) liquid (g/h) production rate (m2/g) %) ratio Ar NH3 H2 (g/h) Si C N O Al Y Al + Y/ Al/ Si Y HSAl03 1900 0 0 87 34 6646.533.912.9 5.72.5 0 HSAl04 1900 0 0 119 32 8143.533.710.713.22.4 0 HSAl07 1900 0 0 88 24 11047.129.112.8 8.62.5 0 0.05 HSAl08 0 0 1900 49 11 28048.726.614.1 8.91.6 0 HSAlY05 1900 0 0 115 30 7943.733.411.211.51.71.4 HSAl09 1140 760 0 71 23 11249.5 6.836.0 6.11.5 0 0.03 HSAl12 1520 380 0 87 28 9647.214.829.0 7.51.5 0 0.03 HSAl13 1710 190 0 87 30 8349.823.019.2 5.92.0 0 0.04 HSAl11 475 570 855 64 27 13449.8 5.837.2 5.71.5 0 HSAl14 855 190 855 74 26 8948.817.324.5 7.71.6 0 HSAl15 665 380 855 75 22 9950.911.726.8 8.81.8 0 HSAlY16 1140 760 0 81 26 10053.0 7.530.3 8.20.40.4 0.01 3.0 HSAlY17 1520 380 0 89 25 9150.716.925.4 5.40.90.7 0.02 4.2 HSAlY18 1710 190 0 79 28 7749.321.922.4 5.10.50.7 0.01 2.5

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Page 22 alumina and yttria (Burger et al. 1997). The first result is that it is possible to incorporate Al and Y additives in the powders, but Table 2.6 clearly shows that the ratio Al/Si, in this range of experimental conditions, does not reach the 0.069 ratio. This is due to the preferential pulverisation of the most volatile compound of the mixture, that is, HMDS. So, in the following experiments, the initial Al/Si and (Al+Y)/Si ratios were decreased to 0.04 (HSAl09, 12, 13) and 0.02 respectively (HSAlY16, 17, 18). The similarity of the Al/Si atomic ratio in the starting solution (0.04) and in the as-formed HSAl powders (0.03–0.04) suggests that the nebulisation of the mixture of precursors is homogeneous. In the case of HSAlY powders, the Al/Y atomic ratios are different in as-formed powders (2.5–4.2) and in the starting solution (1.1), though (Al + Y)/Si atomic ratios remain similar (0.01– 0.02) to the starting solution. This is due to the incomplete dissolution of yttrium isopropoxide in HMDS and probably to an inhomogeneous nebulisation of the mixture of precursors. Effect of carrier gas The nature and composition of the carrier gas mixture play a significant role in the chemical composition of powders. Whatever the powder (HSAl or HSAlY), an increase of ammonia relative flow rate induces a decrease in C content and C/N atomic ratio as obtained for powders without sintering aids (Musset 1995). Table 2.6 also shows that when hydrogen is added partially in the place of argon in argon–ammonia mixtures, the C content decreases. This seems to indicate that more volatile hydrocarbon species are formed during the synthesis. For example, the C content varies from 14.8 to 11.7 wt% in a 80% argon–20% ammonia mixture (HSAl12 and 15 samples) and from 23.0 to 17.3 wt% in a 90% argon–10% ammonia mixture (HSAl13 and 14 samples). It must be noticed that a large quantity of free C (10–14%) remains in powders when the ammonia relative flow rate increases from 0 to 20. 2.4.4 First characterisations Specific surface area Powder production rates in the 22–34 g/h range and specific surface area values in 66–134 m2/g range are very similar to the values obtained for powders without metallic additives. Infrared spectroscopy and X-ray diffraction As for powders without metallic additives, XRD patterns show a large background without any distinct peak indicating an amorphous structure. This is confirmed by infrared spectra showing a large band around 900 cm −1. This one is essentially composed of Si—C bonds when powders are synthesised with argon-rich carrier gas while Si—N bonds in an amorphous environment (large band at 400 cm−1) appear with ammonia-rich carrier gas. Moreover, Si—CH3 bonds are present (1250 cm−1) in all powders. 2.5 Conclusion In this chapter, it has been shown that laser pyrolysis is a versatile method well suited for the production of significant amounts of Si-based nanopowders. Due to their properties (purity

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Page 23 limited only by the purity of the reactants, low size dispersion, etc.), the nanopowders are good candidates for applications in ceramics. Nanopowders can be obtained from gaseous or liquid precursors. The powders obtained from the gas phase are crystallised while the powders obtained from the liquid phase are amorphous. In both cases, the chemical composition of the powders is controlled by the chemical composition of the reactive mixture. Finally, an interesting result is the incorporation of the elements of sintering aids (Al, Y, O) during the synthesis. References Alexandrescu, R., Morjan, I., Borsella, E., Botti, S., Fantoni, R., Dikonimos-Makris, T., Giorgi, R. and Enzo, S. (1991) Journal of Materials Research, 6, 2442. Alexandrescu, R., Borsella, E., Botti, S., Cesile, M. C., Martelli, S., Giorgi, R., Turtu, S. and Zappa, G. (1997) Journal of Materials Science, 32, 5629. Armand, X., Herlin, N., Martinengo, H., Musset, E. and Cauchetier, M. (1995) Fourth Euro-Ceramics, vol. 1 (Faenza Editrice S.p.A., Faenza), pp. 37–44. Baraton, M. I., Boulanger, L., Cauchetier, M., Lorenzelli, V., Luce, M., Merle, T., Quintard, P. and Zhou, Y. H. (1994) Journal of the European Ceramic Society, 13, 371. Bauer, R. A., Smulders, R., Becht, J. M. B., Van der Put, P. J. and Schoonman, J. (1989) Journal of the American Ceramic Society, 72, 1301. Borsella, E., Botti, S., Fantoni, R., Alexandrescu, R., Morjan, I., Popescu, C., Dikonimos-Makris, T., Giorgi, R. and Enzo, S. (1992) Journal of Materials Research, 7, 2257. Borsella, E., Botti, S., Giorgi, R., Martelli, S., Turtu, S. and Zappa, G. (1993a) Applied Physics Letters, 63, 1345. Borsella, E., Botti, S., Alexandrescu, R., Morjan, I., Dikonimos-Makris, T., Giorgi, R. and Martelli, S. (1993b) Materials Science and Engineering A, 168, 177. Bourgeois, L., Barbier, G., Viguié, J. C., Herlin, N. and Cauchetier, M. (1995) Fourth Euro-Ceramics, vol. 1 (Faenza Editrice S.p.A., Faenza), pp. 225–232. Buerki, P. R., Troxler, T. and Leutwyler, S. (1990) High Temperature Science, 27, 323. Burger, P., Duclos, R. and Crampon, J. (1997) Materials Science and Engineering, A222, 175. Casey, J. D. and Haggerty, J. S. (1987a) Journal of Materials Science, 22, 737. Casey, J. D. and Haggerty, J. S. (1987b) Journal of Materials Science, 22, 4307. Cauchetier, M., Croix, O., Luce, M., Michon, M., Paris, J. and Tistchenko, S. (1987) Ceramics International, 13, 13. Cauchetier, M., Croix, O. and Luce, M. (1988) Advanced Ceramic Materials, 3, 548. Cauchetier, M., Croix, O., Robert, C., Lance, M. and Luce, M. (1989) Euro-Ceramics, vol. 1, Elsevier Applied Science, London, pp. 130–134. Cauchetier, M., Croix, O., Luce, M., Baraton, M. I., Merle, T. and Quintard, P. (1991) Journal of the European Ceramic Society, 8, 215. Cauchetier, M., Croix, O., Herlin, N. and Luce, M. (1994) Journal of the American Ceramic Society, 77, 993. Cauchetier, M., Armand, X., Herlin, N., Mayne, M., Fusil, S. and Lefevre, E. (1999) Journal of Materials Science, 34,1. Croix, O., Gounot, M., Bergez, P., Luce, M. and Cauchetier, M. (1991) Ceramics Today – Tomorrow’s Ceramics, Elsevier Science Publishers, Amsterdam, pp. 1447–1455. Curcio, F., Ghiglione, G., Musci, M. and Nanetti, C. (1989) Applied Surface Science, 36, 52. Curcio, F., Musci, M., Notaro, M. and Quattroni, G. (1991) Ceramics Today – Tomorrow’s Ceramics, Elsevier Science Publishers, Amsterdam, pp. 2569–2578. Fantoni, R., Borsella, E., Piccirillo, S., Ceccato, R. and Enzo, S. (1990) Journal of Materials Research, 5, 143. Förster, J., von Hoesslin, M., Schäfer, J. H., Uhlenbussch, J. and Viöl, W. (1991) Proceedings of the 10th International Symposium on Plasma Chemistry (FRG) vol. 1 (eds U. Ehlemann, H. G. Lergon and K. Wiesemann), pp. 1–6.

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Page 24 Fusil, S., Armand, X., Herlin, N. and Cauchetier, M. (1997) Key Engineering Materials, 132–136, 141. Garifo, L. (1995) Laser Focus World, 30. Gonsalves, K. E., Strutt, P. R., Xiao, T. D. and Klemens, P. G. (1992) Journal of Material Science, 27, 3231. Haggerty, J. S. and Cannon, R. W. (1981)Laser-induced Chemical Processing, Plenum Press, New York, pp. 165–241. Herlin, N., Musset, E., Luce, M. and Cauchetier, M. (1994) Journal of the European Ceramic Society, 13, 285. Herlin, N., Armand, X., Musset, E., Martinengo, H., Luce, M. and Cauchetier, M. (1996) Journal of the European Ceramic Society, 16, 1063. Iwamoto, Y., Kikuta, K. and Hirano, S. (1998) Journal of Materials Research, 13, 353. Kennedy, T., O’Neil, J. P., Hampshire, S., Poorteman, M. and Cambier, F. (1996) Silicates Industriels, 9–10, 201. Kizaki, Y., Kandori, T. and Fujitani, Y. (1985) Japan Journal of Applied Physics, 24, 800. Knudsen, A. K. (1987a) Advances in Ceramics, 21, 237. Knudsen, A. K. (1987b) US Patent 4, 689, 129. Knudsen, A. K. (1997) Carbide, Nitride and Boride Materials: Synthesis and Processing, Chapman & Hall, London, pp. 343–358. Koyama, S., Nakashima, H., Sugahara, Y. and Kuroda, K. (1998) Chemistry Letters, 191. Li, Y., Liang, Y., Zheng, F. and Hu, Z. (1994a) Journal of the American Ceramic Society, 77, 1662. Li, Y., Liang, Y., Zheng, F. and Hu, Z. (1994b) Materials Science and Engineering, A174, L23. Luce, M., Croix, O., Zhou, Y. H., Cauchetier, M., Sapin, M. and Boulanger L. (1993) Euro-ceramics II, vol. 1, Deutsche Keramische Gesellschaft, Köln, pp. 233–238. Luce, M., Herlin, N., Musset, E. and Cauchetier, M. (1994) Nanostructured Materials, 4, 403. Martinengo, H., Musset, E., Herlin, N., Armand, X., Luce, M. and Cauchetier, M. (1996) Silicates Industriels, 61,9. Mayne, Martine (May 1997) PhD Thesis, Limoges University, 97LIMO0021; http://thesenet.abes.fr Mayne, M., Armand, X., Cauchetier, M., Doucey, B., Bahloul-Hourlier, D. and Goursat, P. (1999) Ceramics: Getting in the 2000’s, Part B, Techna, Faenza, pp. 211–218. Musset, Emmanuel (Nov. 1995) PhD Thesis, Paris-Sud Orsay University, 95PA112537; http://thesenet.abes.fr Rice, R. W. (1986) Journal of the American Ceramic Society, 69, C183. Rice, R. W. (1987) Journal of the American Ceramic Society, 70, C117. Rice, R. W. and Woodin, R. L. (1988) Journal of the American Ceramic Society, 71, C181. Schoenung, J. M. (1991) American Ceramic Society Bulletin, 70, 112. Soraru, G. D., Ravagni, A., Dal Maschio, R. and Carturan, G. (1992) Journal of Materials Research, 7, 1266. Spitz, J. and Viguié, J. C. (1970) French Patent 2, 110, 622. Suzuki, M., Nakata, Y., Okutani, T. and Kato, A. (1992) Journal of Materials Science, 27, 6091. Suzuki, M., Maniette, Y., Nakata, Y. and Okutani, T. (1993) Journal of the American Ceramic Society, 76, 1195. Sumaya, Y., Marra, R. A., Haggerty, J. S. and Bowen, H. K. (1985) American Ceramic Society Bulletin, 1356. Symons, W. and Danforth, S. C. (1987) Advances in Ceramics, 21, 249. Tougne, P., Hommel, H., Legrand, A. P., Herlin, N., Luce, M. and Cauchetier, M. (1993) Diamond and Related Materials, 2, 486. Wakai, F., Kodama, Y., Sakaguchi, S., Murayama, N., Isaki, K. and Niihara, K. (1990) Nature, 344, 421. Willaime, F., Boulanger, L. and Cauchetier, M. (1995) Materials Research Symposium Proceedings, 359, 53. Yamada, T., Tanaka, Y., Suemasu, T. and Kohtoku, Y. (1997) Nuclear Instruments and Methods in Physics Research, B121, 378.

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Page 25 3 Thermal behaviour of as-formed silicon-based nanopowders General introduction This chapter presents the effects of thermal treatments on Si-based nanopowders obtained as described in the previous chapter from the gas phase or from the liquid phase. It is mainly focused on the behaviour of Si/ C/N and Si/C/N/Al/Y/O powders (0.2 < C/N(at) < 2.6) because such nanopowders may be used to obtain Si3N4/SiC materials after thermal treatment. This chapter is divided into two parts. The first part is devoted to thermal treatment in a high-temperature graphite furnace in order to know the evolution of the physico-chemical characteristics. The second part reports the study of the thermal behaviour using TGA/MS in order to determine the chemical reactions and mechanisms involved during annealing.

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Page 26 3.1 Thermal behaviour in a high-temperature graphite furnace Emmanuel Musset, Martine Mayne, Michel Cauchetier, Xavier Armand, Michel Luce and Nathalie Herlin-Boime 3.1.1 Introduction The aim of the first part of this study is to determine the influence of the annealing treatments performed in a high-temperature graphite furnace both on the weight of samples and on different physico-chemical properties of nanopowders such as the chemical composition, structure, reactivity and size. It is mainly focused on Si/C/N samples obtained from the liquid phase using the aerosol method. The results of annealing treatments and the consequences of these treatments on the chemical composition, size, organisation, etc. of silicon-based powders are presented with the intent of showing the similarities and the differences between the different powders, to correlate these properties to the synthesis conditions, etc., and to begin to compare the properties of these powders with the necessary properties for ceramic applications. 3.1.2 Weight evolution of SiC, Si/C/N and Si/C/N/Al/Y/O nanopowders The as-formed powders are annealed at atmospheric pressure in a high-temperature graphite furnace under argon or nitrogen (N) atmospheres. The temperature generally varies in the range 1000–1700°C, and the dwell time is 1 or 4 h. 3.1.2.1 Weight evolution under inert atmosphere Thermal treatments were performed on SiC or Si/C/N nanopowders in an argon atmosphere for temperatures up to 1600°C. SiC powders exhibit a weight loss whatever the temperature. The maximum weight loss, observed near 1500°C, remains low (usually

Page 27 loss is high (>20%). This also suggests the decomposition of the Si/C/N chemical system, which appears to be more important than in the case of Si/C/N obtained from the gas phase. Detailed mechanisms will be proposed and discussed in the TG–MS section. The decomposition of the Si/C/N chemical system under an inert atmosphere could be partially related to the degradation of the nitrogen-containing species inducing an evolution of nitrogen during the thermal treatment. Usually, the elaboration of silicon nitride based materials is achieved by sintering powders under a nitrogen atmosphere in order to limit the weight loss and the degradation of the silicon nitride based system. That is why a nitrogen atmosphere is used for heat treatment in the following sections. 3.1.2.2 Weight evolution under nitrogen atmosphere The aim of the use of nitrogen is to limit the weight loss by preserving the N content in the as-formed powders in order to further obtain dense silicon nitride based ceramic materials. Heat treatment was performed at 1000°C and then in steps of 100°C between 1300°C and 1700°C with a heating rate of 10°C/min. Most often, the dwell time is 4 h, unless otherwise indicated in the text. Previously, the as-formed Si/C/N powders (about 2 g) are slightly pressed in a Teflon mould to form pellets (diameter = 20 mm; height ≈ 10 mm) which are then placed in graphite crucibles located inside the graphite furnace. For Si/C/N nanopowders synthesised from the gas phase (SiCN29 and SiCN35), a significant mass change (+8 wt%) is obtained for the SiCN35 sample between 1200°C and 1400°C suggesting the nitridation of the free Si present in the powder. Contrary to the SiCN35 sample, the SiCN29 sample exhibits a relative mass stability up to 1600°C (Mayne 1997). This can be related to the high N content and to the absence of Si that cannot enable a nitridation process (Mayne 1997; Mayne et al. 1998). X-ray diffraction (XRD) patterns of the annealed powders show the presence of crystalline phases (α- and β-Si3N4 for SiCN29 sample and a mixture of α- and β-Si3N4 with β-SiC for SiCN35 sample). In contrast to Si/C/N nanopowders obtained from the gas phase, Si/C/N and Si/C/N/Al/Y/O powders obtained from the liquid phase (HMDS or mixtures of HMDS with organic compounds) exhibit weight losses that are of the same order of magnitude (20–30% at 1600°C). Therefore, the weight evolution of Si/C/N/Al/Y/O powders (Cauchetier et al. 1999; Mayne et al. 1999), will not be specifically presented in this chapter. Such weight losses suggest a decomposition of the Si/C/N chemical system even under nitrogen atmosphere. In Table 3.1.1, the weight losses and chemical compositions of six HMDS samples are given. It shows that the weight loss is dependent on the chemical composition of the as-formed powder. After annealing at 1600°C, the weight loss is the highest for the samples with the highest C/N ratio. This difference could be related to different chemical mechanisms involved in the decomposition process and interpretations will be suggested in the TG–MS section. This study on the weight evolution of Si-based samples during heat treatment enables us to show the following: • a strong difference in the behaviour of powders obtained from the gas phase or from the liquid phase using the aerosol method especially under a nitrogen atmosphere. Nanopowders obtained from the gas phase are more stable than those obtained from the liquid phase, which suggests that they are more suitable for ceramic conversion than those obtained from the liquid phase; • that the thermal behaviour of powders obtained from the liquid changes as a function of the chemical composition of the powder, suggesting different mechanisms.

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Page 28 Table 3.1.1 Results of chemical analyses for HMDS40 to HMDS45 Heating treatment (°C) Mass loss (%) HMDS40 AFa 1000 1300 1400 1500 1550 1600 HMDS41 AFa 1000 1300 1400 1500 1550 1600 HMDS42 AFa 1000 1300 1400 1500 1550 1600 1700 HMDS43 AFa 1000 1300 1400 1500 1550 1600 1700 HMDS44 AFa 1000 1300 1400 1500 1550 1600 1700 HMDS45 AFa 1000 1300 1400 1500 1550 1600

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C

N

O

∑b

N/ C/ C/N Si Si

—46.725.021.8 5.948.423.620.9 5.751.323.723.7 6.051.523.722.5 19.459.626.812.6 27.065.728.2 4.7 27.968.629.1 2.9

6.1 99.60.931.24 1.33 5.0 97.90.861.14 1.31 2.2100.90.921.07 1.16 1.5 99.20.871.07 1.23 0.7 99.60.421.05 2.47 0.8 99.40.141.00 7.02 0.5101.10.081.0011.15

—46.120.625.2 4.450.419.323.5 5.851.419.227.3 7.452.518.727.0 17.259.221.019.6 28.661.623.111.8 28.267.425.2 8.5

6.7 98.61.091.04 4.6 97.80.930.89 2.4100.31.060.87 1.7 99.91.030.83 1.1100.90.660.83 1.1 97.60.380.87 1.4102.40.250.87

0.95 0.95 0.82 0.81 1.25 2.28 3.50

—44.418.729.7 4.048.917.927.1 5.951.618.329.5 8.052.618.029.8 14.655.818.422.1 23.465.220.913.1 27.067.222.511.5 28.866.623.3 9.7

7.7100.51.340.98 4.4 98.31.110.85 1.5100.91.140.83 1.6102.01.130.80 0.8 97.10.790.77 1.0100.20.400.75 1.0102.20.340.78 0.3 99.90.290.81

0.73 0.77 0.72 0.70 0.97 1.85 2.27 2.80

—48.315.526.8 4.750.213.830.3 8.053.214.031.5 10.854.314.428.2 18.058.313.026.9 25.358.813.227.3 25.265.416.719.3 28.864.219.417.1

8.0 98.61.110.75 5.9100.21.210.64 3.8102.51.180.61 2.9 99.81.040.62 1.4 99.60.920.52 0.7100.00.930.52 1.0102.40.590.60 0.4101.00.530.71

0.67 0.53 0.52 0.60 0.56 0.56 1.01 1.33

—47.113.026.710.9 97.71.130.64 6.249.811.132.5 7.4100.81.300.52 9.753.311.132.4 4.5101.31.210.48 12.752.610.532.4 3.4 98.91.230.46 19.159.0 9.829.7 2.0100.51.010.39 27.756.910.431.2 0.5 99.01.100.42 24.462.412.123.3 1.1 99.00.750.45 25.762.613.725.6 0.5102.40.820.51

0.57 0.40 0.40 0.38 0.38 0.39 0.61 0.62

—47.6 7.451.8 11.655.4 14.152.9 17.058.5 23.454.7 20.460.3

0.33 0.21 0.20 0.15 0.12 0.05 0.14

8.328.913.4 98.21.210.41 6.134.2 9.7101.81.320.27 5.732.4 5.9 99.41.170.24 4.735.8 4.9 98.31.350.21 3.635.9 0.9 99.01.230.14 1.634.4 0.6 92.31.260.07 4.233.8 1.3 99.61.120.16

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Notes a As-formed powder. b The sum of the different element masses from 100% due to the uncertainty of the chemical analyses and the hydrogen is not taken into account, namely for the as-formed powders.

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Page 29 The as-formed powders obtained from the gas phase are at least partially organised, and the weight changes occurring during thermal treatments are much less noticeable than for the powders obtained by the aerosol method from the liquid phase that are amorphous. 3.1.3 Chemical evolution during annealing in a nitrogen atmosphere As in the previous chapter, the chemical composition in stoichiometric compounds of the different samples annealed under a N2 atmosphere has been calculated from chemical analysis obtained from CNRS analysis laboratory (Vernaison, France). The knowledge of the composition in stoichiometric compounds (Si3N4, SiC, SiO2, Y2O3, Al2O3), especially for annealed powders is important, because it enables us to evaluate roughly the composition of future sintered materials. Moreover, this allows a more precise choice of the adapted powders to obtain the desired material with specific properties. The determination of the chemical composition in stoichiometric compounds for the different samples is as follows: for Si/C/N/Al/Y samples, all Al and Y are assumed to be in the form of Al2O3 and Y2O3, the remaining O being in the form of SiO2, then all N forms Si3N4 and the remaining Si is in the form of SiC. Free C is the difference between total C and C bonded to Si in SiC. The rules are the same for Si/C/N samples apart from Al and Y that are not present. It must be noted that such a calculation is approximate both for asformed and annealed powders. Effectively, in the case of amorphous as-formed powders, compounds such as Si3N4, SiC, Y2O3, Al2O3 do not really exist as crystalline compounds. In annealed Si/C/N powders, crystalline compounds Si3N4 and SiC exist but the crystallisation is not always complete, depending on the annealing temperature (as shown by XRD measurements). In Si/C/N/Al/Y/O powders, Al and Y are not only bonded to O (Y2O3, Al2O3) but may also be combined in SiAlON or YSiAlON amorphous phases (see Chapter 5), as for classical sintering of silicon nitride compounds in the presence of sintering aids (Y2O3, Al2O3) (Ekström and Persson 1990; Hoffmann and Petzow 1993). Note that the calculated Si3N4 and SiC contents should be very similar to that which exists in materials because of the low quantity of Y and Al bonded to N. Furthermore, H present in as-formed powders has not been taken into account for these chemical measurements. The chemical composition in stoichiometric compounds is presented for Si/C/N and Si/C/N/Al/Y powders, both obtained from the liquid phase in Tables 3.1.2 and 3.1.3. Tables 3.1.2 and 3.1.3 also give the C/N ratios in the as-formed powders, because they are often used in the following sections to characterise a sample. 3.1.3.1 Si/C/N samples The samples studied in this section have an initial C/N ratio in the range 0.2–2.6 (HMDS40–45 and HMDS36, 66 and 67). The chemical composition in atomic ratio (C/Si, N/Si and C/N) and its evolution as a function of annealing treatment between 1000°C and 1600°C is presented in Table 3.1.1 for samples HMDS40–45. The results for stoichiometric compounds after annealing at 1600°C, are reported in Table 3.1.2 for HMDS40–45 and HMDS36, 66 and 67. Part of these results appear in the PhD thesis of Emmanuel Musset (Musset 1995). Up to 1400°C, the chemical composition of Si/C/N samples presented in Table 3.1.1 shows no significant changes in the different atomic ratio values. Meanwhile, at higher temperatures

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Page 30 Table 3.1.2 Chemical composition (wt%) of various HMDS samples after annealing 1600°C assuming stoichiometric composition Samples Si3N4 SiC SiO2 C C/N (at) initial ratio HMDS36a 5.3 91.5 1.9 1.3 2.8 HMDS40 7.7 87.6 2.1 2.6 1.3 HMDS41 19.1 75.1 3.3 2.6 0.9 HMDS42 28.6 67.4 3.0 1.0 0.7 HMDS43 56.4 37.7 3.0 2.9 0.7 HMDS44 65.9 29.9 2.8 1.4 0.6 HMDS45 88.9 8.4 2.5 0.3 0.3 HMDS66 6.2 83.0 1.5 9.9 1.8 HMDS67 6.3 82.4 1.5 9.8 2.4 Note a Without ammonia in the carrier gas during synthesis. Table 3.1.3 Chemical composition of Si/C/N/Al/Y/O pre-mixed powders after annealing under N (1 h, 1600°C) assuming stoichiometric composition Powder sample Chemical composition (wt%) (Al + Y)/Si Al/Y C/N initial ratio Y2O3 Al2O3 SiO2 Si3N4 SiC C HSAl03 2.8 0 17.8 70.2 10.4 3.1 HSAl04 7.1 0 18.3 66.7 8.1 3.7 HSAl07 3.1 0 17.6 72.2 5.9 0.04 2.7 HSAl09 3.3 1.7 92.3 0.2 2.5 0.03 0.2 HSAl12 2.6 0.4 52.7 44.0 0.2 0.02 0.6 HSAl13 3.3 0.8 17.0 78.5 0.4 0.03 1.4 HSAl14 3.5 0.5 46.5 48.1 1.4 0.03 0.8 HSAl15 3.9 0 77.5 16.1 2.4 0.04 0.5 HSAlY05 1.3 3.8 0 19.3 64.1 9.0 3.5 HSAlY16 0.1 0.3 1.9 92.4 3.1 2.1 0.01 14.8 0.3 HSAlY17 0.6 0.9 1.4 51.9 43.7 1.5 0.01 3.2 0.8 HSAlY18 1.1 1.1 1.5 14.5 80.9 0.8 0.01 2.3 1.1 different evolutions occur depending on the initial C/N ratio (Tables 3.1.1 and 3.1.2): • For the C-rich samples (HMDS36, 66, 67 and HMDS40), the C/N atomic ratios compared to the as-formed powders increase and the powders are converted essentially towards SiC (between 80% and 90%) during annealing (C/Si atomic ratio near 1, see Table 3.1.1). Free C is always present in these samples after annealing at 1600°C (Table 3.1.2). • For the N-rich samples (HMDS44 and 45), the C/N atomic ratios decrease with a stabilisation of the N content and the powders are converted essentially towards Si3N4 (between 66% and 89%). • For the two samples with intermediate composition (HMDS42 and 43, initial ratio C/N = 0.7), the C/N atomic ratios remain relatively stable with a stabilisation or a slight decrease of the C/Si atomic ratios and a conversion of the powder towards a composite powder containing both SiC and Si3N4.

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Page 31 Whatever the composition of the initial powders, annealing induces a decrease in silica and in free C content because of the carbonitridation of this compound which involves CO evolution. For C-rich powders, the change in composition (Table 3.1.2) to SiC-rich compounds, during annealing at 1600°C under a N2 atmosphere, can be essentially linked to the decomposition of nitrogenous species (Si/C/ N, Si3N4) occurring from 1450°C. Such decomposition is important because of the high C content in asformed powders (C/N > 0.75) (Musset 1995; Musset et al. 1997). Effectively, extended X-ray absorption fine structure (EXAFS) (Chapter 5.4) and nuclear magnetic resonance (NMR) (Chapter 5.1) analysis of C-rich powders synthesised from HMDS have shown that the atomic local arrangement of Si is essentially composed of mixed SiC3N tetrahedra and of SiC4 tetrahedra which are already close to the SiC structure. This suggests the decomposition of slightly nitrogenous compounds (SiC3N) during annealing in order to reach the SiC structure. Moreover, annealing induces a decrease in silica and in free C content because of the carbonitridation of this compound which involves CO evolution. For N-rich powders, EXAFS and NMR analysis have shown that the atomic local arrangement is essentially composed of SiN4 tetrahedra which are the local environment of crystalline Si3N4. Therefore, the evolution towards a major phase of Si3N4 after annealing is not surprising. The drastic change between the evolution of Si/C/N towards SiC or Si3N4 compounds is observed for the initial C/N atomic ratio in the as-formed powders near the 0.75 value and corresponding to a 3SiC–Si3N4 mixture (Figure 3.1.1). 3.1.3.2 Si/C/N/Al/Y/O samples For samples obtained from a mixture of organometallic compounds as precursors, the comments are mainly focused on the aspects specific to the presence of Al and/or Y in the powders. Table 3.1.2 presents the results in stoichiometric compounds after annealing at 1600°C for 1 h under a N2 atmosphere for Si/C/N/Al/Y/ O samples with an initial C/N ratio in the same range as HMDS samples (0.2–2.6). The most important difference between Si/C/N samples obtained from only HMDS and Si/C/N/Al/Y samples obtained from a mixture of organometallic precursors is the incorporation of Al and Y metallic elements into the Si/C/N system, which could modify the evolution of the chemical composition of the Si/C/N system during annealing. That is why a comparison between Si/C/N and Si/C/N/Al/Y samples with a similar C/N ratio is presented below. Tables 2.6 and 3.1.3 clearly show that, in the Si/C/N/Al/(Y) systems, the Al+(Y)/Si ratios remain similar to those of the corresponding as-formed powders. This attests that the thermal treatment under a nitrogen atmosphere (1600°C) does not induce the degradation of all Al and Y containing species. Thus, Al and/or Y additives remain in the powder after annealing, which is important for the future sintering process, in which sintering aids (Y2O3, Al2O3) are necessary in order to obtain dense nanomaterials. By comparing, after heat treatment, HSAl or HSAlY samples (Table 3.1.3) with HMDS samples (Table 3.1.2), one can see that the evolution of the chemical composition in Si/C/N/Al(Y) system as a function of the C/N ratio is comparable to that in HMDS samples previously discussed. In this way, C-rich Si/C/N/Al/Y samples (initial C/N ratio between 1.1 and 2.7) are converted essentially to SiC (content between 64% and 81%, see samples HSAl03, 04, 07 and HSAlY05) while N-rich samples (C/N ratio between 0.2 and 0.3) are

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Figure 3.1.1 Change in phase content (SiC, Si3N4, SiO2 and C) in HMDS powders after 4 h annealing under nitrogen at 1600°C. (Reprinted from Musset 1995.) converted essentially to Si3N4 (content of 92%, see samples HSAl12 and HSAlY17). Samples exhibiting an intermediate composition (C/N ratio between 0.6 and 0.8) are converted to a mixture containing both SiC and Si3N4 (SiC around 44 wt%, Si3N4 between 52% and 53%, see HSAl12 and HSAlY17 samples). Note that Crich samples also contain a noticeable quantity of free C (between 6% and 10%). Meanwhile, even if the evolution of the chemical composition is similar between Si/C/N and Si/C/N/Al(Y)O samples, noticeable differences appear where the C/N values and also the Si3N4 and SiC contents are concerned, especially for Si/ C/N and Si/C/N/Al/Y C-rich samples (C/N = 2.6). Figure 3.1.2 shows the evolution of the C/N ratio after a 1 h annealing treatment at 1600°C under a N2 atmosphere for some C-rich samples (HMDS66, HSAl03, 04 and HSAlY05). In this figure, the C/N ratio is around 15 for HMDS powders while it is only of 4.3–5.1 for HSAl(Y) powders. These results indicate that the loss of N is lower for powders containing Al and Y additives and exhibiting high C/N ratio. The SiC content is then lower and the Si3N4 content is higher in these Si/C/N/Al/(Y) O samples than those in Si/C/N samples (see Tables 3.1.2 and 3.1.3). This suggests that the decomposition of N containing species (Si/C/N, Si3N4) is limited during heat treatment of Si/C/N/Al/Y powders, probably due to the formation of a liquid phase in the

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Figure 3.1.2 Evolution of the C/N atomic ratio with annealing temperature. (Reprinted from Cauchetier et al. 1999 with permission from Kluwer Academic/Plenum Publishers.) quaternary Si–Al–O–N or in the quinary Y–Si–Al–O–N diagram preventing the degradation of nitrogenous compounds. 3.1.4 Structural changes XRD and infrared (IR) measurements were carried out in order to obtain information about the nature of the different phases present and the evolution of structure and of the different chemical bonding after heat treatment. They were carried out on both Si/C/N and Si/C/N/Al/Y samples obtained from the liquid phase. 3.1.4.1 Si/C/N samples For annealing temperatures below 1500°C, the XRD diagrams of Si/C/N nanopowders (HMDS samples) are flat indicating that the powders are still amorphous. The IR spectra are very similar to the spectra of asformed powders (see for example Figure 2.8), apart from the absorption bands at 1170, 2950 and 1255 cm −1 involving hydrogenated bonds that disappeared. At 1500°C, peaks are present on the XRD diagram (Figure 3.1.3), indicating the presence of crystallised β-SiC in HMDS40, 41 and 42 samples and α- and β-Si3N4 phases in HMDS45. For HMDS43 and 44, the peaks are small and broad indicating a lower degree of crystallisation. Figure 3.1.4 presents IR spectra for three samples (C/N = 0.73, 0.67 and 0.57, respectively for HMDS42, 43 and 44) at the same temperature (1500° C). It shows two broad peaks centred near 900 and 500 cm−1 indicating the presence of an amorphous material and corresponding to Si—C—N and Si—N bonds, respectively. These results, together with the XRD result, indicate that crystallisation is not complete at this temperature, especially for samples with an ‘intermediate’ C/N ratio (HMDS43 and 44). This late crystallisation will be discussed in more detail in Chapter 5.5. Between 1500°C and 1600°C, noticeable changes occur both in the XRD diagram and in the IR spectra (Figures 3.1.4 and 3.1.5). Figure 3.1.5 presents the XRD diagrams of samples annealed at 1600°C. C-rich powders (HMDS41 sample) lead preferentially to the β-SiC phase, as was already noticed at 1500°C. N-rich powders (HMDS45 sample) mainly show

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Figure 3.1.3 XRD patterns of HMDS powders after annealing at 1500°C. (Reprinted from Musset 1995.) the presence of α- and β-Si3N4 phases and very small β-SiC peaks. Samples exhibiting intermediate composition (HMDS43 and 44) lead to a mixture of SiC and Si3N4. In the IR spectra (Figure 3.1.4), the broad peak centred near 900 cm−1 becomes sharper between 1500°C and 1600°C. The shape becomes similar to the signature of crystallised SiC nanopowders obtained from the gas phase. The fine absorption bands present in the 800–400 cm−1 range indicate the crystallisation of N-rich powders only present in HMDS43 and 44 samples (α- and β-Si3N4 phases, Luongo 1983). These results (XRD and IR) are in good agreement with the calculated chemical compositions given in Table 3.1.2.

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Figure 3.1.4 Change in IR spectra during annealing under nitrogen between 1500°C and 1700°C. (Reprinted from Musset 1995.) 3.1.4.2 Si/C/N/Al/Y/O samples As mentioned above for Si/C/N samples, annealing treatment under a nitrogen atmosphere also induces a drastic change in the structure of Si/C/N/Al/Y powders. From amorphous, in their as-formed state, they become crystallised or at least partially crystallised after a 1 h treatment at 1600°C under nitrogen atmosphere. Figure 3.1.6 shows XRD diagrams of Si/C/N samples (0.2 < C/N < 1.7) annealed under nitrogen at 1600°C for 1 h. The nature and the crystallographic variety of the phases that crystallise during heat treatment are very similar to those obtained from Si/C/N powders (HMDS samples). Thus, for a C-rich Si/C/N/ Al powder (HSAl13), the XRD diagram shows an evolution towards α- and β-SiC compounds while for a N-rich Si/C/N/Al powder (HSAl09), α- and β-Si3N4 are formed with a higher quantity of the α phase compared to the β phase. For samples having intermediate composition (HSAl12), α-and β-Si3N4 and β-SiC phases are formed. The IR spectra of the same samples (Figure 3.1.7) are in good agreement with the XRD diagram, that is to say that SiC bonds (around 900 cm−1) are detectable in a C-rich system, while Si N bonds in a Si3N4 structure are detectable in a N-rich system. The sample with intermediate composition contains the Si— N bonds in the Si3N4 structure and also Si—C bonds. IR and XRD results are very consistent with the chemical composition mentioned above. For C-rich samples (C/N = 2.6, HSAlY05 and HMDS66 samples), a comparison between the XRD diagrams of Si/C/N and Si/C/N/Al/Y powders annealed at 1600°C is presented (Figure 3.1.8). Moreover, the evolution of the IR spectra as a function of the temperature is presented for the same samples (Figure 3.1.9). XRD patterns of HSAlY05 and HMDS66 annealed at 1600°C are shown in Figure 3.1.8. The Si/C/N/Al/Y powder annealed for 1 h (HSAlY05 sample) presents broad diffraction

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Figure 3.1.5 XRD patterns of HMDS powders after annealing at 1600°C. (Reprinted from Musset 1995.) lines of β-SiC that are indicative of a low degree of crystallisation (Figure 3.1.8(a)). There is also evidence for the presence of the α-SiC phase (2H polytype). When annealing time is increased, well-crystallised α-SiC (2H polytype) and β-SiC are present (Figure 3.1.8(b)). In contrast, for the Si/C/N powder (HMDS66 sample) there is evidence of α- and β-SiC even after only 1 h of annealing and no difference can be noticed in the diffractogram for an annealing duration of 1 or 4 h (Figure 3.1.8(c) and (d)). These results suggest that for Crich powders, the presence of Al and Y either limits the crystallisation or enables the formation of an amorphous phase initiated in the Y–Si–Al–O–N quinary system.


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Figure 3.1.6 XRD analysis of annealed HSAl powder. (Reprinted from Mayne et al. 1999.)

Figure 3.1.7 IR spectra of annealed HSAl powder. (Reprinted from Mayne et al. 1999.) Figure 3.1.9 presents examples of IR spectra of the as-formed and annealed Si/C/N and Si/C/N/Al/Y/O powders (high C/N ratio). Up to 1400°C, the spectra of Si/C/N/Al/Y/O powder are very similar to the spectra obtained for Si/C/N powder. The only difference seems to be a flat base line for the powders containing Al (+Y) additives whereas a pronounced dip appears near 1000 cm−1 in the IR spectrum of powders without additives

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Figure 3.1.8 XRD patterns of samples (a) and (b) HSAlY05 and (c) and (d) HMDS66 after 1 and 4 h annealing under N at 1600°C. (Reprinted from Cauchetier et al. 1999 with permission Kluwer Academic/Plenum Publishers.)

Figure 3.1.9 Changes in IR spectra with annealing temperature. (Reprinted from Cauchetier et al. 1999 with permission Kluwer Academic/Plenum Publishers.) indicating crystallisation or grain growth as observed previously in the annealing of SiC powders obtained by a laser gas-phase driven reaction (Tougne et al. 1993) or from pyrolysis of a polycarbosilane (Sasaki et al. 1989). No significant change is observed when the dwell time increases from 1 to 4 h and no signature characteristic of Al or Y is observed. This study of chemical and structural changes during annealing of Si/C/N and Si/C/N/Al/Y/O samples under nitrogen atmosphere has shown a comparable evolution of Si/C/N and Si/C/N/Al/Y/O samples with comparable C/N ratio. A more detailed comparison seems to indicate that the presence of elements of sintering aids limits the degradation of the nitride phase during heat treatment of Si/C/N/Al/Y/O samples. Another point is that the elements Al, Y and O aids remain present in the powder after thermal treatment, which is encouraging for future elaboration of materials.

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Page 39 3.1.5 Information about changes in the grain sizes BET measurements were carried out in order to measure the specific surface area, which enables determination of the changes in reactivity and grain size after heat treatment. These measurements should be related to transmission electron microscopy (TEM) analysis (see Chapter 4). 3.1.5.1 Si/C/N samples Figure 3.1.10 shows the variation of only two Si/C/N samples (HMDS41 and HMDS44). Below 1000°C, the SBET decrease can be explained by the smoothing of the particles due to the reticulation reactions, which occur during the ceramisation of the powders. Then the SBET value remains constant up to 1400–1500°C and the decrease at 1600°C is due to the crystallisation and growth of the particles, as confirmed by TEM (Chapter 4). 3.1.5.2 Si/C/N/Al/Y/O samples Annealing (nitrogen atmosphere, 1600°C, 1 h) of Si/C/N/Al/Y/O powders induces a decrease of the specific surface area compared with that of initial powders attesting to a structural and morphological change during heat treatment such as crystallisation, structural modification of free C and particle growth. The difference between the specific surface areas of as-formed and annealed powders increase when the C/N atomic ratio decreases. Thus for C-rich Si/C/N/Al/O powders (HSAl07 sample), the specific surface area varies from 110 m2/g (as-formed) to 54 m2/g (annealed) while for N-rich Si/C/N/Al/O powders (HSAl09 sample), the decrease of the specific surface area is more important (112–12 m2/g). A similar evolution is also true for the series of Si/C/N/Al/Y/O samples. This phenomenon could suggest a different grain growth process in C-rich powders compared to N-rich powders, that is to say that the grain growth could be more pronounced in Nrich samples. Apart from a low grain growth,

Figure 3.1.10 Change in specific surface area during annealing under Nfora C-rich HMDS powder (HMDS41 sample) and a N-rich HMDS powder (HMDS44 sample). (Reprinted from Musset 1995.)

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Page 40 this could also be related to the low crystallisation degree (see XRD patterns above) and to a structural modification of free C giving rise to some C ribbons as was already observed (Musset 1995) and as is usually formed from the heat treatment of amorphous C between 1500°C and 2000°C (Dresselhauss et al. 1988). These interpretations must be checked by TEM observations in order to determine the main phenomenon responsible for the surface area results. 3.1.6 Conclusion From the first section of this chapter devoted to thermal behaviour, several interesting points can be noted: • The thermal behaviour of Si/C/N powders can be related to synthesis conditions; powders obtained from the gas phase appear more stable than powders obtained from the liquid phase using the aerosol method. • Heat treatment under nitrogen allows one to obtain crystallised powders with a chemical composition varying from a major phase SiC to a major phase Si3N4 with intermediate mixtures of SiC + Si3N4. • The Al, Y and O elements of sintering aids remain present in the Si/C/N/Al/Y/O powders after thermal treatment under N atmosphere. The behaviour of these powders is qualitatively comparable to the behaviour of Si/C/N powders with comparable C/N ratio. The presence of these in situ sintering aids can be very helpful to obtain homogeneous dense materials – their influence is presented in Chapter 7. References Cauchetier, M., Croix, O., Herlin, N. and Luce, M. (1994) Journal of the American Ceramic Society, 77, 993. Cauchetier, M., Armand, X., Herlin, N., Mayne, M., Fusil, S. and Lefevre, E. (1999) Journal of Materials Science, 34,1. Croix, O., Gounot, M., Bergez, P., Luce, M. and Cauchetier, M. (1991) Ceramics Today: Tomorrow’s Ceramics, Elsevier Science Publishers, Amsterdam, p. 1447. Dresselhaus, M. S., Dresselhaus, G., Sugihara, K., Spain, I. L., Golbert, H. A., Cardona, M. (1988) Graphite Fibers and Filaments, Springer Verlag. Ekström, T. and Persson, J. (1990) Journal of the American Ceramic Society, 73, 2834. Hoffmann, M. J. and Petzow, G. (1993) Materials Research Symposium Proceedings, 287, 3–14 Luongo, J. P. (1983) Journal of the Electrochemical Society, 130, 1560. Mayne, Martine (May 1997) Ph.D. Thesis, Limoges University, 97LIMO0021; http://thesenet.abes.fr Mayne, M., Balhoul-Hourlier, D., Doucey, B., Goursat, P., Cauchetier, M. and Herlin, N. (1998) J. Eur. Ceram. Soc., 18, 1187. Mayne, M., Armand, X., Cauchetier, M., Doucey, B., Bahloul-Hourlier, D. and Goursat, P. (1999) Ceramics: Getting in the 2000’s, Part B (Ed. P. Vincezini) Techna, Faenza, pp. 211–218 Musset, Emmanuel (Nov. 1995) Ph.D. Thesis, Paris-Sud Orsay University, 95PA112537; http://thesenet.abes. fr Musset, E., Herlin, N., Cauchetier, M. and Luce, M. (1997) Silicates Industriels, 62, 111. Sasaki, Y., Nishima, Y., Sato, M. and Okamura, K. (1989) Physical Review B: Condensed Materials, 40, 1762. Tougne, P., Hommel, H., Legrand, A. P., Herlin, N., Luce, M. and Cauchetier, M. (1993) Diamond and Related Materials, 2, 486.

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Page 41 3.2 Thermal reactivity of silicon-based nanopowders Djamila Bahloul-Hourlier, Benoît Doucey, Jean-Louis Besson and Paul Goursat 3.2.1 Introduction Si/C/N nanopowders may be converted into Si3N4/SiC bulk materials by consolidation and hot pressing. Several studies have been devoted to the synthesis of Si/C/N nanopowders and their sintering. However, only a few studies on the thermal decomposition of these preceramic nanopowders have been reported (Giorgi et al. 1996; Suzuki et al. 1995; Borsella et al. 1997; Mayne et al. 1998; Pan et al. 1999), in contrast to the situation with organometallic polymer precursors of Si/C/N ceramics (Burns et al. 1987; Blum et al. 1989; Han et al. 1992; Corriu et al. 1992; Schmidt et al. 1992; Lavedrine et al. 1991; Bahloul et al. 1993, 1997). It is well known that ultrafine powders can undergo structural changes when they are heated at specific temperatures. This is due to their high surface energy and non-equilibrium phases. A good knowledge of adsorbed species on the surface, the nature of chemical groups present in nanopowders and their reactivity during annealing may be used to further improve the colloidal processing, shaping and forming techniques and also to optimise sintering conditions and thus mechanical properties. The aim of the present work is to further clarify the effects synthesis conditions and chemical composition of the preceramic Si/C/N nanopowders have on their thermal stability while under different atmospheres. The synthesis conditions and chemical composition of the studied powders as well as their corresponding designation are reported in Table 3.2.1. The studied samples have a C/N ratio varying from 0.2 to 2.6. Various experimental techniques: thermogravimetric analysis coupled with mass spectrometry (TGA/MS), elemental analysis and X-ray diffraction (XRD) were used to follow the evolution of annealing in preceramic nanopowders. Table 3.2.1 Nanopowders chemical composition Powder % NH3 (cm3/min) Chemical analysis (wt%) Atomic ratio C/N Chemical composition (at %) Si C N O Al Al2O3 SiO2 Si3N4 SiC Si C N SiCN29 40 58 6 35 1 — 0.2 —1 87 8 — 4 — SiCN35 0 67 13 18 2 — 0.8 —4 45 45 6 — — HMDS35 40 51 6 38 5 — 0.2 —9 77 — — 6 8 HMDS66 0 51 32 15 2 — 2.4 —5 37 39 — 19 — HMDS73 10 52 13 25 10 — 0.6 —19 63 7 — 11 — HSAl07 0 47 29 13 9 2 2.6 412 32 33 — 19 — HSAl09 40 49 7 36 4 2 0.2 49 76 — — 7 6

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Page 42 3.2.2 Results and discussions 3.2.2.1 TG/MS analysis Figure 3.2.1 shows a comparison of the measured weight losses during the decomposition of the as-received Si/C/N preceramic nanopowders in flowing helium (He). At first sight, the results indicate a direct dependence of the ceramic yields on the synthesis conditions. Obviously, the overall weight loss is greater for the family of nanopowders produced from liquid precursor (HMDS) than for that obtained from gas reactant SiH4 (gas). The major weight loss variations basically evolve at high temperature (T > 1300°C), where extensive degradation of the preceramic nanopowders occurs (Figure 3.2.1). A combined TGA/MS study was then used to determine the identity of the evolved species at each step. On account of the large number of powders produced by particular synthesis conditions, it is not possible to illustrate here all the mass spectra. Only the data corresponding to the sample that exhibited the highest weight loss (HMDS35) will be given in detail. However, when differences exist, they will be compared with the other samples. The TG/MS analysis of HMDS35 under He is given in Figure 3.2.2. On heating from room temperature to 300°C, the weight loss of 1% corresponds to the simultaneous escape of water (M/z = 18, 17) and ammonia (M/z = 17, 16, 14). The intensity of water progressively diminishes with increasing temperature up to approximately 300°C, while ammonia seems to stabilise between 400°C and 600°C and then diminishes with further heating. The appearance of two shoulders at M/z = 17 is evidence for the occurrence of at least two different reactions.

Figure 3.2.1 TG profiles of different nanopowders pyrolysed under He flow.

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Figure 3.2.2 TG–MS analyses of HMDS35 nanopowder pyrolysed under He flow. The escape of H2O and NH3 at low temperatures (less than 300°C) may be a consequence of the higher latent reactivity induced by the presence of the condensable ≡Si—NH2 and Si OH groups on the preceramic nanopowders (Equations 3.2.1, 3.2.2a, b). (3.2.1) (3.2.2a)

(3.2.2b) While the nitrogen (N) bridging may take place at low temperatures, the reaction between silanol surface groups to form oxygen bridges is known to take place above 350°C (Equation 3.2.3). (3.2.3) Nakamatsu et al. (1998) proposed the following reaction (Equation 3.2.4) to explain the simultaneous desorption of both H2O and NH3 at low temperatures (around 200°C). (3.2.4) Besides NH3 and H2O, some other fragment ions at M/z = 44, 30 are also detected below 400°C. The intensities of these peaks are very low and thus prevent a detailed characterisation of the origin structures. Between 300°C and 900°C, a second weight loss (3.8%) is attributed to the second escape of ammonia (M/z = 17, 16), which indicates the occurrence of the

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Page 44 transamination/condensation reactions (Equation 3.2.5) which are known to occur during the thermal conversion of polysilazanes (Bahloul 1993).

(3.2.5) Concomitant with the evolution of NH3 is the loss of other gases with M/z varying from 42 to 12 atomic mass units (AMU) (Figure 3.2.3). It is noteworthy that compounds with molecular weights higher than M/z = 42 were not detected. Looking at these peak fragments, one would attempt to assign them to light hydrocarbons containing up to three carbon (C) atoms at most, like alcyne, olefin and alkane compounds. However, it should be remembered that the synthesis of this powder was made with NH3 as a reactant gas, being very reactive over Si—H, Si—C and C—C bonds. Consequently, the probability of formation of hydrocarboned species having two or three carbons in their structure seems very weak (Van Dijen and Pluijmakers 1989). Moreover, it is seen from the comparison of the intensities of the obtained peaks that the most intense peak at M/z = 41 is followed, by descending order in intensity, by the peak at M/z = 40, then M/ z = 39, and finally M/z = 42. No hydrocarbon combination with two or three carbons leads to such an evolution. It is, therefore, believed that these fragments may be assigned to R—C≡N species such as CH3 CN (M/z = 41, 40, 39, 38), HCN (M/z = 26, 27) and also to less protonated species such as carbo-imines like CH3 —CH=NH (M/z = 42, 28). Further data that confirms the assignment of such species will be presented during the analysis of the exhaust gas at high temperatures (T > 900°C).

Figure 3.2.3 MS analysis of HMDS40 nanopowder pyrolysed under He flow.

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Page 45 The maximum desorption of these species is situated at 600°C, whereas in the temperature range (600–900° C) the major volatile products detected are methane (M/z = 16, 15, 14) and H2 (M/z = 2). Methane continues to evolve up to 1000°C, whereas hydrogen evolution occurs in a wider temperature range up to 1500°C with a maximum loss at about 900°C. Some of these gaseous emissions are classical during the thermal conversion of polycarbosilane and polycarbosilazane polymers and were described by several authors whereas others are more specific to this type of laser-driven powders. Considering this temperature range, the reactions that describe the thermal decomposition of these materials are those that most probably involve radical mechanisms. There are several potential sources of radical species that can explain the above exhausted molecules (Equation 3.2.6–3.2.12). (3.2.6) (3.2.7) (3.2.8) (3.2.9) (3.2.10) (3.2.11) (3.2.12) For example, the formation of CH3—C≡N arises from radical cleavage of Si—CH2—C≡N leading to the •CH2— C≡N radical, followed by H abstraction and/or radical combination reactions. The formation of the other species can also be explained using this approach. The C—N≡groups do not appear in the initial structure of HMDS precursors, and seem to have been formed during the synthesis process. This is confirmed by Rice’s studies (1986) on the laser induced dissociation of the molecule of HMDS in the gas phase under argon. Analysis of gases of the processing chamber indicated the presence of, amongst other species, HCN and hydrocarbons, in particular. Thus, it is not surprising to think that, in our case, as the synthesis is being carried out in the presence of NH3, the concentration of C—N≡or amine groups giving HCN and RCN species, should increase at the expense of C—C groups yielding to light hydrocarbon species. Indeed, the reaction of C or hydrocarbon compounds with ammonia has been well described elsewhere and the resulting product is HCN while no hydrocarbons are formed (Corriu et al. 1992; Bahloul et al. 1993). Above 900°C, a significant weight loss is observed. This is accompanied by the increased evolution of the shape of the TGA curve. Between 900°C and 1400°C, besides H2, the major gaseous decomposition product is unambiguously ascribed to HCN at M/z = 26, 27. This exhaust gas gives credence to the description given above concerning the attribution to HCN and amine species of the signals at M/z between 42 and 12 AMU. This result is in good agreement with the chemical analysis of the as-received HMDS35 (Table 3.2.1) from which an excess of N and C are obtained assuming that only the equilibrium phases SiO2, SiC, Si3N4 and C, are present in the as-prepared powders (rule of mixture calculations). Furthermore, the short-range atomic structure description of the same Si—C—N nanopowders by XAS reveals the presence of C—N bonds in the amorphous network (Ténégal et al. 1996). A structural model has been proposed (Ténégal et al. 1996), where mixed tetrahedral SiCxN4−x are randomly linked through C—N bonds. The release of HCN at high temperature is further

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Page 46 proof that CN groups are not only pendant groups X—CN as observed for temperature T < 900°C, or radical scavenger/quencher groups under the experimental conditions employed in the synthesis process, but the branching groups to probably a mixed tetrahedra SiCxNy. Homolytic cleavages of C—N bonds between two mixed tetrahedra of silicon need more energy in comparison with those at chain-ends. Parallel to the evolution of HCN gas is the loss of N2 at M/z = 28, 14 up to the temperature limits of the experiment (1500°C) whereas the CO loss at M/z = 28, 12 begins at 1450°Cin minor amounts (comparison of ions at M/z = 12 for CO and M/z = 14 for N2). At 1500°C, despite the important weight loss, the intensity of N2 diminishes rapidly, some traces of CO are still detected by MS during the isothermal holding. Interestingly, at 1500°C, during the holding time, we observe that the mass continues to decrease with increasing time, whereas any species cannot be detected by MSin large enough quantities to account for this continued weight loss. It is believed that this weight loss in the TGA curve is due to compounds such as SiO which condense rapidly in cold parts of the apparatus and may not be detected by MS. Evidence for this proposal comes from the deposition of SiO as a wool-like product on many parts of the apparatus, which consequently leads to the breaking of the thermocouple based on platinum. The most probable chemical reactions in the Si/C/N/O system, that involve the evolution of volatile species and which produce a significant weight loss, can be described in two steps. Primarily, decomposition of the quaternary SiCxNyOz phase leads to the formation of an intermediate ternary oxynitride solid phase SixNyOz and the simultaneous release of SiO and CO, N2, in which N2 is the gas detected in the greatest quantity by MS. Then, in the second step, which occurs during holding at 1500°C, the oxynitride phase SixNyOz undergoes a decomposition in which gaseous N2, Si(l, g), SiO and a solid phase are produced according to the following reactions:

(3.2.13) The most known ternary oxynitride phase is Si2N2O. Although several investigators have reported on the thermal decomposition of Si2N2O, no mechanism has yet been clearly established. This is due to some of the following experimental difficulties: (a) the reproducibility of experiments, (b) the error in the determination of the weight loss, which is sometimes not representative of the real value of mass loss due to the condensation of the SiO species on the rod attaching the crucible to the balance, (c) progressive clogging of the aperture of the tube furnace and capillary for admission of gas to the mass spectrometer, by the deposit of SiO, and (d) the rupture of the thermocouple, stopping subsequent treatment. The majority of decomposition reactions of Si2N2O suppose the simultaneous escape of volatile species SiO(g) and N2(g). Our analyses revealed that, after 10 min at 1500°C, no gas species are detected by MS, despite the continued weight loss as observed by TGA. It may be assumed that as the powder HMDS35 has been synthesised with ammonia as a reactant, a noticeable number of SiN4 sites are present and the amount of C in this powder is low. Taking into account our previous observations, it is thus hypothesised that the oxynitride SiNxOy resulting from the decomposition of oxycarbonitride phase SiNxOyCz coexists with SiN4 sites initially present in the material. Hence when the duration of heating is increased, the oxynitride phase degrades to give SiO evolution in greater quantity, while the two other species

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Figure 3.2.4 XRD patterns of nanopowders pyrolysed under a reactive (20%N2/He) and inert (He) atmosphere. Si(l, g) and N2 interact on the SiN4 sites to give better organised Si3N4 as supported by XRD illustrated in Figure 3.2.4. Rocabois et al. (1996) used Knudsen effusion MSto study the thermal degradation of several systems based on Si like binary, tertiary and even quaternary ones containing Si/C/N/O (Rocabois 1993). They stressed that the thermal degradation of Si-based materials is much less pronounced than what is predicted by thermodynamics and that their vaporisation is hindered and related to the different evaporation coefficients of each species. Moreover, the grain size, chemical composition, surface of contacts or interfaces during vaporisation would play an important role in the degradation of various phases. In our case, the nanopowder is initially amorphous as shown by XRD (Figure 3.2.4) and we can postulate that the number of surface contacts is high and the degradation of material is more homogeneous. By admitting, as suggested by Rocabois et al. (1996), that the evaporation coefficient of N2 is low in comparison with that of SiO, the two concomitant reactions, decomposition of oxynitride phase with release of SiO, Si(l, v), and N2, and the reaction of the two latter gaseous species on SiN4 sites initially present leading to new Si3N4 by epitaxial growth, is expected. 3.2.2.2 Major differences in thermal behaviour of the preceramic nanopowders The results pointed out major differences between the nanopowders synthesised from the same HMDS precursor with NH3 (HMDS35 and HMDS73) or without (HMDS66). Elemental

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Figure 3.2.5 MS analysis of HMDS nanopowders pyrolysed under He flow. analysis of these powders confirmed that when NH3 is used during synthesis, the C/N ratio decreases. This diminution of C also depends on the concentration of NH3 in the reactant gas mixture. Moreover, during the heat treatment of these powders, the MS analysis also confirmed that the powders synthesised in the presence of NH3 contain more N—X bonds (X = H, C, Si). Indeed, the evolution of NH3 at M/z = 17 (Figure 3.2.5) occurs over a wide temperature range and is characterised by two waves for HMDS35 and HMDS73, whereas for HMDS66, it exhibits only one narrow wave up to 300°C. In the case of HMDS66 between 400°C and 700°C, the MS spectra (Figure 3.2.6) reveal the loss of Sicontaining species at M/z = 73, 59, 45, 44, 30, which suggests the formation of (CH3)x≤4Si, whereas for HMDS35 and HMDS73, species with M/z greater than 42 AMU were not detected. In view of the abundance of such organosilicon species, it seems that under the experimental conditions used during synthesis, the laser interaction with HMDS does not induce the total atomisation (bursting) of the molecular structure since NH3 is more reactive toward Si—C, which implies the amination of Si and C sites. This is the first, chief origin of the formation of N-rich powders. Above 800°C, it can be seen that for preceramic powders prepared in the presence of NH3 during the laser synthesis process, HCN evolves in larger quantities compared to the powders prepared without NH3.

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Figure 3.2.6 MS analysis of HMDS nanopowders pyrolysed under a He flow. It can also be observed, that as the temperature increases (T > 1200°C) the release of N2 and CO illustrated in Figure 3.2.6 for the C-rich powder HMDS66 and the N-rich powder HMDS35 is different. Indeed, the release of N2 and CO takes place during the entire thermal cycle for HMDS66, whereas no gas is detected after 10 min at 1500°C for HMDS35 N-rich powder. This suggests that the thermal degradation of these two nanopowders proceeds according to different reaction paths and depends on the chemical composition as well as on the microstructure of the materials. The CO response for C-rich powder HMDS66 was not monomodal indicating that the process involves more than one reaction. The oxynitride intermediate phase suggested for a N-rich powder in reaction (3.2.13) would not be formed, since an excess of C is present. Regarding previously discussed thermal decomposition of silicon carbonitride material containing free-carbon phase obtained from polysilazane precursors, degradation of C-rich nanopowders may be related to the following reaction (Bahloul 1993):

(3.2.14) The loss of CO starts at about 1150°C, whereas the escape of N2 is observed from 1300°C. This can be explained in analogy to the reaction (Equation 3.2.15) observed in the solid phase

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Page 50 by the carboreduction of Si O bonds by excess C. (3.2.15) In an analogous fashion, the large amount of N2 evolved during the holding time at 1500°C can be explained by the reaction of Si—N and C: (3.2.16) Several authors (Rocabois et al. 1996; Lee and Jacobson 1992) have suggested that this reaction is inhibited when a SiC layer is formed on the surface of the C. This causes the system to shift to the equivalent SiC/ Si3N4 system and the excess loss of N2 would be suspected from the direct decomposition of Si3N4 as shown in the following equation: (3.2.17) It was supposed that, as the experimental procedure of heat treatment is the same for all powders, we should expect that N-enriched powders generate more N2 than C-enriched powders, since they contain more Si3N4 phase. This was not the case. Consequently, the decomposition of Si3N4 is not able to offer an explanation for the evolution of N observed. Furthermore, powder samples after heat treatment showed no evidence for the presence of metallic silicon when examined by XRD (Figure 3.2.4). SiC is one of the main compounds in the samples obtained from C-rich powder HMDS66 heated at 1500°C under He. The low emission of CO during holding time can be attributed to the following reactions: (3.2.18) (3.2.19) We, therefore, suggest that in our system the SiC obtained from the reaction of Si3N4 with C (3.2.16) may be consumed by SiO2 following reaction (3.2.18), which in turn releases C particles and consequently allows reaction (3.2.16) to continue producing gaseous N2. 3.2.2.3 Thermal behaviour of nanopowders under a nitrogen atmosphere The reason for studying the heat treatment under N was to answer the question: could the decomposition of silicon carbonitride, occurring under He, be stopped by using N environmental gas? The structural evolution of various nanopowders was thus followed as a function of heat treatment temperature and time whilst under a N atmosphere. The resulting curves in Figure 3.2.7, which represent the sole nanopowders that have the same ratio C/N = 0.2, show that below 1300°C the weight losses are nearly identical regardless of the nature of the atmosphere (nitrogen or He for the same kind of powder). The main gases detected by MS are attributed to the same species as those detected under He. In contrast, above 1400°C, comparison of the variation of mass is noticeably lower under N than He, as is shown for the HMDS35 sample. These differences are more pronounced for SiCN nanopowders obtained from only gaseous reactants (SiH4 + CH3NH2 + NH3 + Ar), as shown in the data in Table 3.2.2. The volatile species which evolve are similar to those obtained under He with a majority of CO but of a lesser degree. Furthermore, it can be seen from Figure 3.2.8 that the variation

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Figure 3.2.7 MS analysis of HMDS nanopowders pyrolysed under He flow. Table 3.2.2 Nanopowders weight loss changes according to atmosphere heat treatment Atmosphere heat treatment Weight loss ΔM/Mo (%) SiCN35 SiCN29 HMDS66 HMDS35 HSAl07 HSAl09 He −7.7 −6.8 −16.3 −19.3 −20.6 −19.8 20%N2/He +6.1 +1.3 −6.2 −15.5 −14.6 −13.1 of mass is remarkably decreased with increasing partial pressure of N. These results suggest that carbothermal nitridation reactions take place at high temperatures under N according to the following reactions (Mayne et al. 1998): (3.2.20)

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Figure 3.2.8 TG profiles of different nanopowders pyrolysed under a 20%N2/He flow. For C-enriched powders that contain more SiC phase, the nitriding process is also possible via the following reaction: (3.2.21) The weight gain observed, particularly for SiCN synthesised from gaseous reactants, can be attributed to the nitridation reaction of excess metallic silicon detected by XRD (Mayne et al. 1998). According to the α- or βSi3N4 content calculated from XRD patterns of heat treated powders, the nitridation reaction leads to the preferential formation of the α phase. This could be explained by an oversaturation in nitrogen coming from the nitrogen atmosphere, which results in the crystallisation of the less ordered α phase. Meanwhile, the β silicon nitride phase that is more ordered and stable, is preferentially formed during annealing under the inert He atmosphere. 3.2.2.4 Influence of the nature of the precursors’ source of silicon on the chemical composition and thermal stability It is observed that the aluminium doped powders (HSAl), synthesised under the same conditions as that of HMDS powders, except that only aluminium isopropoxide precursor is added, exhibit similar behaviour. Likewise, the differences observed between HMDS nanopowders and their conditions of synthesis, in particular the NH3 pressure in the reactive mixture of synthesis, seem to be reproduced for the HSAl powders. The role of aluminium isopropoxide introduced during the synthesis on the thermal behaviour is being studied. The use of a gas source of Si such as SiH4 has led to nanopowders of a more stable nature in terms of temperature, without doubt, due to their crystallisation state as revealed by XRD.

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Page 53 The formation of CN bonds does not seem to be favoured when SiH4 is used. However, contrary to the HMDS and HSAl powders, the use of silane (SiH4) generates an excess of metallic silicon within the powders, which will be used profitably during sintering under N. 3.2.3 Conclusion These studies indicate that the relative atomic concentrations of Si, C, N and H are primarily influenced by the types of carrier gases used for the synthesis laser process. Residual functionalities C—H, N—H and Si—H bonds will still be present as supported by the identification of volatile species by MS. The study of heat treatment of such systems based on Si indicates that thermal degradation is not an individual reaction but the net result of many reactions. The difference in thermal behaviour is attributed to variations in the preceramic structure. Taking into consideration the gases detected during our investigation and comparing them with similar results found by other research groups (9, 7), we are able to put forward the suggestion that by analysing preceramic nanopowders by thermal decomposition, it is possible to identify the processes by which they were synthesised. A potential advantage of laser-driven synthesis of refractory materials over gaseous reactants is the ability to include additives and co-reactants (e.g. sintering aids) (Cauchetier et al. 1999) in a truly homogeneous fashion. References Bahloul, D., Pereira, M. and Goursat P. (1993) J. Am. Ceram. Soc., 76(5), 1156–68. Bahloul, D., Pereira, M. and Gerardin, C. (1997) J. Mater. Chem., 7(1),109. Blum, Y. D., Schwartz, K. B. and Laine, R. M. (1989) J. Mater. Sci., 24, 1707. Borsella, E., Botti, S., Martelli, S., Alexandrescu, R., Cesile, M. C., Nesterenko, A., Giorgi, R., Turtu, S. and Zappa, G. (1997) Silic. Indus., 1–2,3. Burns, G. T., Angelotti, T. P., Hannemen, L. F., Chandra, G. and Moore, J. A. (1987) J. Mater. Sci., 27, 4651. Cauchetier, M., Armand, X., Herlin, N., Mayne, M., Fusil, S. and Lefevre, E. (1999) J. Mater. Sci., 34(21), 5257. Corriu, R. J. P., Leclercq, D., Mutin, P. H. and Vioux, A. (1992) Chem. Mater., 4, 711. Giorgi, R., Turtu, S., Zappa, G., Borsella, E., Botti, S., Cesile, M. C. and Martelli, S. (1996) Appl. Surf. Sci., 93, 101. Han, H. H., Lindquist, D. A., Haggerty, J. S. and Seyferth, D. (1992) Chem. Mater., 4, 705. Lavedrine, A., Bahloul, D., Goursat, P., Choong Kwet Yive, N. S., Corriu, R. J. P., Leclercq, D., Mutin, P. H. and Vioux, A. (1991) J. Eur. Ceram. Soc., 8, 221. Lee, K. N. and Jacobson, N. S. (1992) Ceram. Eng. Sci. Proc., 16th Ann. Conf. ou Comp. and Adv. Ceram. Mater., Cocoa-Beach. Mayne, M., Bahloul-Houlier, D., Doucey, B., Goursat, P., Cauchetier, M. and Herlin, N. (1998) J Eur Cer Soc, 18(9), 1187. Nakamatsu, T., Saito, N., Ishizaki, C. and Ishizaki, K. (1998) J. Eur. Ceram. Soc., 18, 1273. Pan, Z. W., Xie, S., Wang, G., Li, H. L. and Zhang, L. T. (1999) J. Mater. Sci., 34, 3047. Rocabois, P., Chatillon, C. and Bernard, C. (1996) J. Am. Ceram. Soc., 79(5), 1351. Rocabois, Philippe (1993) PhD, I.N.P.G., 93INPG; httl://thesenet.fr Rice, G. W. (1986) J. Am. Ceram. Soc., 69(8), C183. Schmidt, W. R., Marchetti, P. S., Interrante, L. V., Hurley, W. J., Lewis, R. H., Doremus R. H. and Maciel, G. E. (1992) Chem. Mater., 4, 937. Suzuki, M., Hasegawa, Y., Aizawa, M., Nakata, Y. and Okutani, T. (1995) J. Am. Ceram. Soc., 78(1), 83. Ténégal, F., Flank, A.-M. and Herlin, N. (1996) Phys. Rev. B, 54(17), 12029. Van Dijen, F. K. and Pluijmakers, J. (1989) J. Eur. Ceram. Soc., 5, 385.

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Page 54 4 Structure of some nanopowders and micro/nano composites by transmission electron microscopy Marc Monthioux, Catherine Béraud and Nahida El Horr 4.1 As-prepared and heat-treated nanopowders Some of both the raw and heat-treated powders were investigated by transmission electron microscopy (TEM). Sample preparation prior to TEM investigations and TEM operating conditions are reported in Section 4.3, together with a brief reminder of the principles for the various TEM modes used in the study. 4.1.1 Synthesis from gaseous precursors Starting from a mixture of silane (SiH4) and ethylene (C2H2) as gaseous precursors, a serie of SiC-based powders was obtained. Likewise, still using silane as Si source but replacing ethylene by ammonia (NH3) and/ or CH3NH2 allowed Si/C/N-based powders to be obtained. Si/C-based powders Materials were investigated either as-prepared or after various treatments. Related preparation conditions and elemental analysis of the as-prepared materials are reported in Tables 2.2 and 2.3, respectively. Some of the samples reported in Table 2.2 were investigated after various treatments, including SiC173, SiC177 and SiC212. Various annealing in argon atmosphere were carried out, in order to check structural and textural changes in both the SiC and C phase. Also, SiC173 and SiC212/1200 were investigated after thermal treatment at 450°C performed in air (1–3 h isothermal), in order to eliminate the carbon (C) phase when present. Finally, further heat-treatment experiments in inert atmosphere were also performed at 1800°C on sample SiC173 oxidised. Special precautions were taken for the 1800°C heat-treatment experiments (Laanani 1997) to prevent any possible oxidation reaction due to oxygen traces in the gas phase, that is, a non-dynamic, high purity argon atmosphere (N56 grade from Air Liquide) and an all-graphite furnace were used. SIC173, AS-FORMED Daylight colour of the as-prepared powder is dark grey, somewhat related to the polyaromatic C content, as already expected from the elemental analysis (see Section 3.1.1) and as supported in the following by the direct evidence from TEM investigation. Grains exhibit round and isometric shapes (Figure 4.1).

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Figure 4.1 TEM bright-field image from SiC173. Arrows show stripes on SiC crystals indicating polytypes as stacking faults.

Figure 4.2 TEM lattice fringe image from SiC173. Note how two adjacent SiC crystals are able to accommodate a slight mismatch in orientation of their respective lattices, making both the 111 lattice fringes connected to each other without interruption. Grain sizes range according to a two-mode distribution. Large grains (10–60 nm) are prevalent, each being a single crystal, sometimes faulted (Figure 4.1, arrows). Grains are associated into aggregates, most often chain-like, due to both weak forces (electrostatic) or stronger ones, that is, structural. Figure 4.2 illustrates the latter, showing two grains whose lattices are connected despite a misorientation of 10°, obviously inducing a strong (covalent) bonding between grains. They are SiC crystals, β structure mainly, but some stacking faults

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Page 56 (indicating the local occurrence of polytypes) are revealed in electron diffraction patterns as additional aligned spots (Figure 4.3), as a probable consequence of the rapid growth. Grains are coated with a thin (~1 nm) layer of an amorphous phase, assumed to be a silicon oxide, as supported by X-photon energy dispersive spectroscopy (X-EDS) measurement, whose spectra show the presence of the element O, in addition to Si and C as expected (Figure 4.4). 4nm–1

Figure 4.3 Electron diffraction of SiC crystals from SiC173. The structure is mainly β-SiC (see Figure 4.7 for diffraction ring identification). However, arrows show some aligned spots indicating stacking fault sequences. The occurrence of an amorphous coating at the grain surface (thickness ~1 nm) is probable, though confusion with overfocusing Fresnel fringes is possible.

Figure 4.4 X-EDS spectrum of SiC crystals from SiC173. The presence of O attests the occurrence of an amorphous silica layer associated to the SiC grains.

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Page 57 The minor occurrence of a silica phase is consistent with the results from elemental analysis (Table 2.3). On the other hand, small grains are a very minor phase, whose sizes (3–10 nm) are difficult to estimate from regular bright-field images. Dark-field images help in more easily appreciating the occurrence and size distribution of it, as compared to the major phase (Figure 4.5(c), arrows). They are not Si3N4 but possibly SiC crystals too, since they appear in Figure 4.5(c) while not in Figures 4.5(a) or (b). However, other small, minor crystals are also found, despite that the amount of it is too small to appear in the diffraction pattern, from which high-resolution images reveal fringes whose spacing is too large to match the SiC lattice distances (not illustrated). Rarity and concentration of this minor phase are such that identification was not possible. A third phase is found associated with the major (SiC) phase, more or less coating the grain surface, and identified as polyaromatic C. Its low proportion is such as it does not appear in the electron diffraction patterns when exposure times are appropriately chosen to reveal the SiC phase (Figure 4.3). However, its structure is most probably turbostratic (i.e. not having the three-dimensional structure of graphite), as assumed from its morphology and nanotextural characteristics. Lattice fringe imaging mode illustrates the latter (Figure 4.2). The C phase appears as thin, somewhat buckling and polygonal, ribbon-like graphene stacks. Using the parameters N (number of graphenes per stack) and L2 (length of the continuous, though distorted, d002 lattice fringe) to characterise the nanotexture of the C phase, as proposed by Oberlin (1989), the latter exhibits a rather low N values (N ~ 2–5) with respect to

Figure 4.5 TEM dark-field image sequence from SiC173. (a) and (b) correspond to ‘polyaromatic Cspecific’ (DF1) dark-field imaging conditions for two selected orientations. The selected orientation direction is indicated by the ‘=’ sign, that is, only graphene stacks displayed parallel to the ‘=’ sign are imaged. Graphene stacks appear as small bright rims between SiC crystals. (c) corresponds to ‘SiC-specific’ (DF2) dark-field imaging conditions. Arrows show the occurrence of the minor phase (probably SiC) as small size grains among the larger SiC crystals.

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Page 58 a rather high L2 value (up to 20–30 nm). As far as the existence of free edges of graphenes is unlikely, since inducing improbable dangling bonds, and since the so-called ribbons never appear twisted as genuine ribbons might be, the graphene stacks are believed to be walls of three-dimensional, somewhat polyhedral, polyaromatic C shells. These shells are likely to form under mechanisms related to the thermal and/or catalytic decomposition of a carbide phase for instance (Badami 1965; Oberlin and Rouchy 1971; Audier et al. 1981). Similar textures, whose nanotexture parameters are fairly consistent with those measured here, were also found in oxygen-free SiC-based materials obtained from heat-treated polycarbosilane (Delverdier et al. 1991; Laanani 1997). Low magnification, specific (i.e. using an objective aperture selecting some diffracted polyaromatic 002 beams) dark-field images (Figures 4.5(a) and (b)) illustrate the proportion and location of the C phase relative to the SiC grains, as the latter appear as round-shaped, grey to pale white objects while the former appears as bright, elongated, tiny spots (arrows) in which graphenes are oriented parallel to the elongation direction. Though not systematic, graphene stacks are often found oriented parallel to the grain surface. SIC177, 1200°C ANNEALED The daylight colour of the starting powder is dark grey. The material is very similar to SiC173/as-prepared. Slight discrepancies concern the distribution of grain size (one-mode instead of two-mode), and the size range (slightly larger: 10–100 instead of 10–60 nm). Another difference lies in the probable absence of the non-identified, crystallised, minor phase of SiC173. Instead, another minor phase is found, mainly amorphous since specific darkfield images show that it does not contain any crystallised phases, neither Si3N4 nor SiC (not illustrated). Considering the high diffusion contrast of this phase in dark-field images when the beam tilt angle corresponds to the Bragg angle for the 002 reflection of polyaromatic C, it is believed that this amorphous phase is enriched in poorly organised, turbostratic C, as expected from the transformation of a formerly amorphous C phase due to the annealing temperature (1200°C is actually far above the minimal carbonisation temperature – ~500–600°C – i.e. the temperature sufficient to transform any kind of nongaseous C precursor into polyaromatic solids). This phase is distinct from the better organised, polyaromatic, turbostratic C phase associated with the SiC phase (Figure 4.6). Figure 4.6 exhibits a C shell still partially filled with a SiC crystal (arrow), which strongly supports that the porous C phase originates from a carbide decomposition process. SIC212, 1200°C ANNEALED The daylight colour of the starting powder is dark grey. The sample is mainly made of agglomerated SiC grains with round shapes (see Figure 4.1 for reference). Grain size range is 3–40 nm, that SiC-specific darkfield images show with a probable two-mode distribution similar to that of SiC177/as-formed (see Figure 4.5 (c) for reference). The structure is strictly β since only diffraction rings corresponding to 111, 200, 220 and 311 reflections of β-SiC appear (Figure 4.7). SiC crystals are coated with amorphous silicon oxide, ~1.5 nm thick (Figure 4.8). The coating may help the SiC crystals to link to each other (Figure 4.8). However, although confusion with an amorphous or poorly organised C coating was unlikely due to the annealing temperature, an oxidation experiment was performed to ascertain the oxide nature of the coating (see below). Unlike SiC173/as-formed or SiC177/1200°C, the relatively well-organised, turbostratic, polyaromatic C phase associated to the SiC phase (see Figures 4.2 or 4.6 for reference) is not

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Figure 4.6 TEM lattice fringe image from SiC177/1200. Arrow shows a small SiC crystal filling incompletely a polyaromatic C shell, supporting the hypothesis that the polyaromatic C phase between SiC grains originates from the early decomposition of some primary SiC crystals.

Figure 4.7 Electron diffraction of SiC crystals from SiC212/1200. The structure is exclusively β-SiC. found. Instead, a poorly organised, turbostratic, polyaromatic C phase is found, concentrated into a minor phase consequently exhibiting a nanoporous texture (Figure 4.9), assumed to be similar to that found in sample SiC177/1200. SIC177, 1500°C ANNEALED The thermal treatment did not induce any change but the average SiC grain size, which has increased mainly due to a coalescence mechanism by impingement according to

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Figure 4.8 TEM lattice fringe image of SiC crystals from SiC212/1200. The amorphous, about 2 nm wide silica coating is imaged, somewhat bonding crystals together.

Figure 4.9 TEM bright-field image from SiC212/1200. The left part of the image illustrates the minor, poorly organised, polyaromatic C phase. Le Coustumer et al. (1993), Monthioux and Delverdier (1996) and Laanani (1997). The smallest grains are now about 30–40 nm, many are in the range of ~100 nm, and the largest diameters are difficult to measure by TEM due to the chances of overlapping for crystals whose large dimensions make them no longer electron transparent. In addition, an amorphous oxide layer about 2 nm thick at the grain surface is still found, quite similar to that of Figure 4.8. Nevertheless, the C phase associated to SiC crystals is also still found, though less abundant, with about the same textural and structural features than that found in Figures 4.2 or 4.6. Electron diffraction patterns show that β is the most frequent SiC structure, but the occurrence of

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Page 61 diffraction spots located besides the regular 111, 200, 220 and 311 diffraction rings for the 3C polytype indicates that other polytypes (α structures) are present (Figure 4.10). SIC173, 1800°C ANNEALED The main effect of such a high heat-treatment temperature is the development of the polyaromatic C phase, either as SiC-free porous material or shells around SiC crystals (Figure 4.11). Though still turbostratic as ascertained by electron diffraction patterns, since only 00l reflections and hk bands are present (Figure 4.12 (a)), the nanotextural state of the C phase is very good, that is, exhibiting perfect graphene stacks at long range (Figure 4.13), which is consistent with a probable origin from the partial SiC decomposition (Oberlin et al. 1976; Iijima

Figure 4.10 Electron diffraction of the SiC phase from SiC177/1500. The 3C (β structure) polytype is prevalent, but some other polytypes are found.

Figure 4.11 TEM bright-field image from SiC173/1800. The partial decomposition of SiC crystals induces the secondary formation of empty or crystal-filled C shells.

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Figure 4.12 Electron diffraction patterns from SiC173/1800. (a) On the C phase: the structure is turbostratic, since only hk reflections are found in addition to the 002 reflection. (b) On SiC crystals (+C). The 3C polytype of SiC is no longer prevalent.

Figure 4.13 TEM lattice fringe image from SiC173/1800. Walls of the C shells seen edge-on. The nanotexture (i.e. graphene perfection) is very good, though the structure is not graphitic but turbostratic according to electron diffraction patterns. 1982). Indeed, though the decomposition of SiC might sound doubtful since 1800°C is still early regarding the theoretical sublimation temperature of SiC (2700°C), the SiC crystals are the only C source available to explain the increase in the C phase amount. Beside, the average SiC grain size increased dramatically (more than two to four times larger than in the starting material), which is consistent with the similar effect observed in the SiC177/1500°C sample (but in a lesser extent). Meanwhile, the diffraction patterns now exhibit many diffraction spots other than that of the 3C polytype, indicating the occurrence of other polytypes (Figure 4.12(b) as compared to Figure 4.3).

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Page 63 SIC173, 450°C OXIDISED As a result of the oxidising treatment, the daylight colour of the powder has turned from dark grey to greengrey. The treatment was successful in removing the primary polyaromatic C phase without altering the SiC grains (Figure 4.14), as compared to the non-oxidised, as-prepared powder (Figure 4.15). Every other feature was kept unchanged. SIC173, 450°C OXIDISED, 1800°C ANNEALED A secondary polyaromatic C phase has appeared again, abundant, as C shells systematically surrounding the SiC grains (see Figure 4.11 for reference). This experiment definitely ascertains that SiC crystals start to decompose at temperatures as low as 1800°C (but beyond

Figure 4.14 TEM bright-field image from SiC173/450—O2. The oxidation treatment was successful in removing the C phase without oxidising the SiC crystals.

Figure 4.15 TEM bright-field image from SiC173. Arrows show the polyaromatic C phase associated to the SiC crystals. To be compared with Figure 4.16.

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Page 64 1500°C, as shown by the 1500°C annealing experiment of sample SiC177, where only the primary C phase was found), although the average grain diameter increases. The fact that this temperature is far below the theoretical sublimation temperature for SiC (2700°C) may be explained by the small sizes of the SiC grains. Generally speaking, physical constants of a given phase are likely to be lowered when finely partitioned into nanocrystals as compared to macrocrystals, due to a higher contribution of surface atoms relative to core atoms (Buffat and Borel 1976). SIC212, 1200°C ANNEALED, 450°C OXIDISED Like for sample SiC173, the daylight colour of the sample has turned from dark grey to green-grey. This is the only change observable. Specifically, the amorphous coating on the SiC grains is still found (not illustrated), ascertaining that the coating is silicon oxide, since amorphous C should have burnt into CO2 during the oxidation experiment. Si/C/N-based powders Two materials were investigated as-prepared, labelled SiCN29 and SiCN35. Related preparation conditions and elemental analysis of the as-prepared materials are reported in Tables 2.2 and 2.3, respectively. Though the amount of silane is similar, they mainly differ from each other both by the amount of nitrogen (N) supplied (twice for the former), and the N-containing gaseous precursors used (CH3NH2 only for the latter, as compared to CH3NH2 + NH3 for the former). This induces a higher supply in C for SiCN35. SICN29, AS-FORMED Daylight colour of the material is light grey-green. As often, the sample exhibits two kinds of morphologies. The major phase is made of large grains, in the range of 20–150 nm, with frequent round shapes (Figure 4.16), sometimes polyhedral (Figure 4.17). According to the

Figure 4.16 TEM bright-field image from SiCN29.

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Figure 4.17 Low magnification TEM lattice fringe image of some crystals from SiCN29. dark-field images, each grain is a single crystal (Figure 4.18(b)). Grains are associated to each other into aggregates or chains both by covalent bonding, since lattice fringe imaging suggests structural continuity between adjacent crystals through grain boundaries (Figure 4.19), or by weak forces. They are identified as α-Si3N4 crystals by electron diffraction (Figure 4.20), as supported by the occurrence of the specific reflection 101 at 0.432 nm. However, the concomitant occurrence of the β structure is possible. Indeed, from the point of view of selected area electron diffraction mode, diffraction patterns from α- and β-Si3N4 structures are difficult to distinguish from each other, due to the high amount of reflections many of which are present in both structures on one hand, and due to the frequent absence of either complete diffraction patterns or complete diffraction rings on the other hand. Although lattice fringe images reveal some distorted fringes at the Si3N4 crystal surface looking like turbostratic C stacks (Figure 4.19), it is believed that, unlike in the SiC-based powders, they do not correspond to a polyaromatic C phase. Indeed, the X-EDS spectra (Figure 4.21) do not show any presence of C (either sp2 or sp3), neither associated to the Si3N4 major phase nor to the SiOxNy minor phase (see below). Also, considering the chemical balance for a SiC-based sample like SiC173 (Table 2.3) shows that 2 wt% only of calculated excess C are enough to provide obvious evidence of polyaromatic C stacks in the lattice fringe or dark-field images (Figures 4.2 or 4.5(a) and (b), respectively). On the contrary, twice a calculated amount of excess C in SiCN29 would barely correspond to the few distorted stacks (in lattice fringe mode, Figure 4.19) or bright rims (in dark-field mode, Figure 4.18(a), open arrow) shown in the TEM images. The absence of a polyaromatic C phase is therefore suggested, and the phase found at the crystal surface is more likely a silicon oxinitride phase. Actually, previous studies combining both TEM and XPS investigations and performed on various Si3N4 powders have shown that silicon oxinitride is able to exhibit the same kind of distorted fringes as turbostratic C (Madigou et al. 1990a, b; Madigou 1991). However, since the chemical balance indicates the presence of C atoms, they could possibly be incorporated into the amorphous silicon oxinitride phase with such a low proportion that would make C non-detectable. Such amorphous silicon oxicarbonitride phases do actually exist, as synthesised from the thermal curing under gaseous oxygen or air of polycarbosilazane polymers (Okamura et al. 1984; Mocaer et al. 1993).

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Figure 4.18 TEM dark-field image sequence from SiCN29. (a) ‘C-specific’ (DF1) imaging conditions. Actually, crystals appearing bright are not C but Si3N4 crystals. They are imaged in the selected imaging conditions because the crystals exhibit lattice distances close to that of C002. Likewise, the white arrow shows bright rims which are more likely an organised Si—O—N phase rather than a polyaromatic phase, according to XEDS spectra (see text). The black arrow shows a polyaromatic C-rich phase. (b) ‘SiC-specific’ (DF2) imaging conditions. Crystals appearing bright are actually not SiC but Si3N4. They are imaged in the selected imaging conditions because the crystals exhibit lattice distances close to that of SiC111. The fact that the crystals are thus able to be imaged in both dark-field imaging conditions allow them to be identified as Si3N4 instead of C or SiC.


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Figure 4.19 TEM lattice fringe image on Si3N4 crystals from SiCN29. Evidence of a coating layer, ~1 nm thick, at the crystal surface, whose distorted fringes are oriented parallel to the grain surface.

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Figure 4.20 Electron diffraction from SiCN29. Si3N4 crystals typically exhibit a lot of lattice distances at small scattering angles.

Figure 4.21 X-EDS spectrum on some Si3N4 crystals from SiCN29. The occurrence of O and the absence of C suggests that the phase coating the crystals is actually Si—O—N-type. The minor phase is made of grains whose size range is often lower than for the major phase, with round shapes. Since always faintly scattering the electron beam in dark-field imaging mode whatever the position of the objective aperture (Figures 4.18(a) and (b)) – though decreasingly as the tilt angle increases – the structure of this phase is amorphous. The chemical composition is qualitatively given by X-EDS measurements (not illustrated), which indicate an enrichment in oxygen and a depletion in N relative to the Si3N4 phase (Figure 4.21), in addition to the presence of Si. It is therefore believed that this minor phase is amorphous silicon oxinitride or oxicarbonitride.

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Page 68 SICN35, AS-FORMED Daylight colour of the material is dark grey-green. As sample SiCN29, two morphologies are again observed. The major phase exhibits large grains with irregular and polyhedral shapes associated into compact aggregates (Figure 4.22). Size range cannot be ascertained due to the superimposition effects which make the grain contours difficult to determine. Dark-field imaging reveals that each grain is itself an aggregate containing crystals with various sizes up to about 40 nm (Figure 4.23(c)). These crystals were identified as pure β-SiC (=3C polytype only) by electron diffraction (Figure 4.24). Both dark-field (Figure 4.23(a)) and lattice fringe (Figure 4.25) images reveal that the SiC crystals are associated with an amorphous phase somewhat embedding them. According to the X-EDS analysis which shows the presence of N and O in addition to Si and C (not illustrated), this amorphous phase is probably a silicon oxinitride or oxicarbonitride. No turbostratic C phase was revealed. The minor phase exhibits much smaller grain sizes, in the range of 10–30 nm, with round shapes, associated into elongated chains (Figure 4.22, arrow). Dark-field imaging (Figure 4.23(b), arrow) supported by lattice fringe imaging (Figure 4.26) both reveal that each grain is itself an aggregate of tiny crystals, few nanometres large. They were identified as Si crystals (CFC structure) by means of electron diffraction (Figure 4.27). Correspondingly, lattice fringe images show lattice distances consistent with the (111) planes of Si (Figure 4.26). These images also reveal that Si crystals are associated with an amorphous phase, located at their outer surface as a coating, about 2 nm thick (Figure 4.26). According to the X-EDS analysis which barely shows another light element than O (in addition to the prevalent Si peak – not illustrated), this amorphous phase is probably a silicon oxide phase, which is likely considering the sensitivity of pure Si to oxidation.

Figure 4.22 TEM bright-field image from SiCN35. Arrows show chain-like aggregate of small grains.

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Figure 4.23 TEM dark-field image sequence from SiCN35. (a) ‘C-specific’ (DF1) imaging conditions. No crystals appear, indicating that no Si3N4 crystals are present (see Figure 4.20 caption for further explanation). (b) Dark-field imaging conditions selected are intermediate between (a) and (c), therefore admitting both C002 and SiC111 distances. Both SiC crystals and another phase (arrowed) are imaged. (c) ‘SiC-specific’ (DF2) imaging conditions. The azimuthal angle is same as for (b). Correspondingly, the same SiC crystals as in (b) are imaged, but the crystals arrowed in (b) no longer appear, attesting that they are not SiC, but Si crystals (see Figures 4.26 and 4.27).

Figure 4.24 Electron diffraction on a SiC-rich area from SiCN35. Crystals are pure β-SiC crystals.

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Figure 4.25 TEM lattice fringe image from the SiC phase from SiCN35. SiC crystals are embedded in an amorphous phase.

Figure 4.26 TEM lattice fringe image from the Si phase from SiCN35. Si crystals are coated with an ~1–2 nm thick amorphous phase, very probably silica. Conclusion regarding ceramic nanopowders from gaseous precursors SI/C-BASED MATERIALS Though SiC173 was the only material investigated as-prepared, it is believed that comparing SiC173 with SiC177/1200 and SiC212/1200 is possible since the annealing temperature is rather low. Since the ethylene/ silane ratio is about the same (~0.5) for all the samples, the main difference in preparation conditions lies in the total flow rate. It increases from SiC177 to SiC173 to SiC212 for a similar C/Si ratio in the precursor gases, inducing a decrease in the residence time from 12.1 to 7.3 then 2.5 ms, respectively (Table 2.2). Correspondingly,

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Figure 4.27 Electron diffraction on a Si phase-rich area from SiCN35. Crystals are pure Si, CFC structure. the grain size range decreases, and SiC polytypes other than the primary form 3C (β structure), though present in the two former samples, do not have the time to develop in the latter. Likewise, the organisation of the polyaromatic C phase tends to decrease. Conditions for the preparation of SiC212 prevent the occurrence of a primary polyaromatic C phase. One explanation could be that long residence times used for SiC173 and SiC177 allow a Si—C—(H) solid solution enriched in C (relative to the 1:1 stoichiometry of SiC) to form in the reaction zone. The subsequent crystallisation of the SiC phase makes the excess C to be expelled as a polyaromatic C phase associated to the SiC crystals. Heat treatments in inert atmosphere provoke the increase in the average SiC grain size due to a coalescence mechanism (i.e. growth from coalescence) by impingements, while the primary C phase, when any, remains unchanged up to 1500°C. Beyond 1500°C (at least from 1800°C), the latter mechanism competes with an early decomposition mechanism, due to the nanometric size of the SiC phase. The partial decomposition of SiC induces the formation of a secondary polyaromatic C phase. Meanwhile, the SiC structure tends to increasingly change from a majority of 3C (cubic) polytype to a mixture of various polytypes (except 2H which is irreversibly unstable beyond ~1600°C – Inoue and Kourachi 1983), probably mostly as stacking fault sequences within the same grains. An explanation is that the coalescence mechanism involves solid crystals, that does not allow single crystals with defect-free, homogeneous structure to be obtained. SI/C/N-BASED MATERIALS Correspondingly to the higher N content in the precursor gas phase for SiCN29 (twice that of SiCN35), Si3N4 crystals only develop in the former. A SiC phase develops in SiCN35 instead, further enhanced by a higher amount of C supply. The fact that no Si3N4 phase is able to form in SiCN35 is amazing since, consistently with the theoretical phase content calculated from the elemental composition (Table 2.3), the N supply should be sufficient. This is not merely due to a lower formation kinetic for Si3N4 relative to SiC, since Si3N4 is able to form in SiCN29 while the residence time is lower. The main parameter is therefore believed to be the nature of the primary gases, and possibly the presence of C specifically.

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Page 72 Thus, when combined with silane and ethylene, a C-free precursor like NH3 is more likely to induce the Si3N4 formation than a C-containing precursor like CH3NH2. The reason is probably that the former precursor gas combination only provides C—H, Si—H and N—H bonds which are easier to thermally break than the preexisting C—N bonds in the latter. Si N bonds are therefore allowed to form more easily and early in the former than in the latter. 4.1.2 Synthesis from liquid precursors Nebulisation of liquid precursors such as hexamethyldisilazane (Si2C6NH18) produced other Si/C/N-based powders. Adding an increasing amount of NH3 allowed the C/N atomic ratio of the as-prepared material to be controlled. The various experiments result in a whole series of samples listed in Table 2.5, and characterised by a continuous decrease in their C/N atomic ratio. On the other hand, the same principle of using an increasing NH3/Ar ratio for a constant total flow rate was applied to a nebulised liquid precursor made of a mixture of HMDS and aluminium isopropoxide (see Section 2.4.2) in order to prepare Si/Al/C/N-based powders. Materials obtained and relative preparation conditions are reported in Table 2.6. Si/C/N-based powders Among the whole series of samples listed in Table 2.5, HMDS42 and HMDS44 were investigated by TEM. They correspond to a slope change in the variation trend of the C/N ratio value measured in the starting materials v. the SiC/Si3N4 content in the 1600°C-annealed materials (see Figure 3.9). HMDS42 and HMDS44 exhibit C/N ratios of 0.73 and 0.56, respectively. Since amorphous at the as-prepared state, they were investigated after increasing annealing heat treatments under pure N atmosphere at temperatures ranging from 1000°C to 1600°C. HMDS42, 1000°C ANNEALED The bulk structural evolution of the sample with increasing annealing temperature from 1000°C and beyond is illustrated by electron diffraction, whose patterns are gathered on Figure 4.28. After 1000°C annealing, the daylight colour of the material is black. Though all the grains exhibit a rather round and isometric morphology (Figure 4.29(a)), a two-mode distribution of grain size is observed, with small grains (20–60 nm) being prevalent, and the subsidiary occurrence of large grains (80–60 nm). Small grains exhibit a nanometric porosity revealed by a Fresnel contrast effect when unfocused. Grains are associated into aggregates (often chain-like) either by strong, structural bonds (illustrated by the chain-like aggregate made of three large grains at the centre of Figure 4.29(a)), or by weak (electrostatic) forces (as illustrated by the small grain aggregate at the top of Figure 4.29(a)). It is worth noting that isolated grains look isometric, but aggregates are not. Considering one morphology (isometric) or the other (chain-like for instance) will provide very different shape factors, which is likely to explain possible mismatches between actual and calculated surface area of the powder. Electron diffraction patterns indicate an overall amorphous structural state, both for large and small grains (Figures 4.28(a) and (b)). However, patterns from small grains exhibit an additional modulation (white arrow) whose location suggests the early occurrence of

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Figure 4.28 Electron diffraction from HMDS42 on large grain and small grain phases respectively after increasing annealing treatment. From (a) to (g) the arrows show the occurrence of a faint C002 reflection.

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Figure 4.29 TEM bright-field image and dark-field image sequence from HMDS42/1000. (a) illustrates the differentiation into two phases with different grain size range. (b) ‘C-specific’ dark-field imaging conditions. The polyaromatic C phase appears as tiny bright dots representing primary graphene stack entities ~1 nm large. (c) ‘SiC-specific’ dark-field conditions. No phase has crystallised yet. very small (


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Page 75 surface, where it now appears as bright rims in dark-field images (Figure 4.30(a)), and as small fringe stacks in high-resolution images, mostly oriented parallel to the grain surface, either at small grain surface (Figure 4.31) or large grain surface (though C is systematically less abundant than for small grains, not illustrated). Such a preferred occurrence of the polyaromatic stacks at the grain surface in C-rich amorphous ceramics with increasing annealing temperature is a common observation (Monthioux et al. 1990; Delverdier et al. 1992a,b, 1993, 1994). It is probably due to easier atomic movements at the grain surface, allowing C segregation on one hand and a graphene display more or less parallel to the grain surface on the other hand. The proportion of the C phase is obviously much higher in small grains than in large grains. It is still poorly organised, since it is made of distorted graphenes about 1–2 nm wide (L2) only, stacked by N = 2–4 only. Also, the occurrence of some brighter dots irregularly disseminated in SiC-specific dark-field images (Figure 4.30(b)) indicates the very beginning of crystallised SiC-phase segregation as nanometric crystals, possibly clusters. Such an early crystallisation step was hardly guessed from the related diffraction patterns (Figures 4.28(c) and (d)) due, on one hand to the relative amount of scattering volumes for the amorphous and the crystallising parts of the material respectively, on the other hand to the overlapping of the SiC reflections, which are very faint and broad as a consequence of the tiny sizes of the SiC crystals involved.

Figure 4.30 TEM dark-field image sequence from HMDS42/1300. (a) ‘C-specific’ (DF1) imaging conditions. Graphene stack entities parallel to the ‘=’ sign are selected, showing that they tend to align parallel to the grain surface. (b) ‘SiC-specific’ (DF2) imaging conditions. Nanometric dots appear, indicating the early stage of overall SiC nucleation.

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Figure 4.31 TEM lattice fringe image from HMDS42/1300 (small grains). Poorly organised graphene stacks at the grain surface are imaged. HMDS42, 1400°C ANNEALED Same features as above (Figure 4.28(e)), but the occurrence of the C phase becomes higher and higher, while its organisation improves, mainly by increasing L2 more than increasing N (not illustrated). HMDS42, 1500°C ANNEALED In both small and large grains, the overall crystallisation of a β-SiC phase is now obvious, as revealed by the evidence of an intense, though broad, 111 reflection (black arrow in diffraction patterns, Figures 4.28(f) and (g) for small and large grains, respectively). Crystal sizes are still very small, nanometric, inducing wide reflections so that 220 and 311 overlap. This is verified in SiC-specific dark-field images, where crystals as large as ~5 nm are very seldom, most of them being in the 1 nm range (Figure 4.32(b)). It is worth noting that the 002 reflection of the turbostratic C phase, now close to 0.344 nm, still appears for small grains only (Figure 4.28(f)). The C phase is now obviously present at the small grain surface and possibly within them (Figure 4.32(a)). It is responsible for an apparent nanometric porosity (Figure 4.33). High-resolution images illustrate the improvement in the abundance and organisation of the polyaromatic C phase associated with small (Figure 4.34) and large (see Figure 4.19 as reference for the abundance and organisation of the surface C exclusively) grains, respectively. Compared to the chemical composition of the grains in the material annealed at 1000°Cas seen by X-EDS(see Figure 4.21 for reference), the chemical composition given by X-EDS analysis of the 1500°C-annealed material reveals a depletion in oxygen content indicated by a lower O/N ratio in both small and large grains (Figure 4.35). This suggests that oxygen atoms were initially not involved in SiO4 tetrahedra from silica since the latter is stable at 1500°C, but possibly in SiOxCyNz tetrahedra similar to that formed in polymer-derived Si/C/N-based ceramics (Delverdier et al. 1994). This is consistent with the fact that no other crystallised phase but polyaromatic C and SiC is evidenced, indicating the necessary occurrence of an amorphous silicon oxi(carbo)nitride phase as an intergranular phase between

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Figure 4.32 TEM dark-field image sequence from HMDS42/1500 (small grains). (a) ‘C-specific’ (DF1) imaging conditions. As the contrast between the bright dots and the background increases, the organisation of the graphene stacks improves. (b) ‘SiC-specific’ (DF2) imaging conditions. SiC growth has started.

Figure 4.33 TEM bright-field image from HMDS42/1500 (small grains). The development of the C phase provides a textured aspect to the grains. Superimposition effects of the projected image do not allow to ascertain whether most of the C is inside or outside the grains.

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Figure 4.34 TEM lattice fringe image from HMDS42/1500 (small grains), illustrating the improvement of the polyaromatic C phase organisation and development with temperature (compare with Figure 4.31).

Figure 4.35 X-EDS spectrum from HMDS42/1500. Valid for large or small grain phases. SiC grains. The partial removal of oxygen atoms from this amorphous phase due to the annealing temperature should probably be related to the crystallisation of the SiC phase, as a cause or a consequence of it for thermodynamic reasons. HMDS44, 1400°C ANNEALED As determined by means of electron diffraction, dark-field, bright-field, and lattice fringe imaging modes, its features are very similar to that of sample HMDS42 annealed at the same

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Figure 4.36 TEM bright-field image from HMDS44/1400. The extensive development of a polyaromatic C phase at the grain surface is obvious. temperature, and will therefore not be extensively illustrated again. However, one discrepancy is a slightly smaller grain size range, that is, 20–50 nm. Another is that crystallisation of SiC phase has not even started yet. The average polyaromatic C content is difficult to estimate since the material appears again heterogeneous, with both grains associated with an abundant C phase (Figure 4.36), and grains associated with few C phase (see Figure 4.19 as reference for the abundance and organisation of the surface C) or none. At least, there is no obvious contradiction with the result from the bulk elemental analysis, which reveals a continuous decrease in the C content from sample HMDS40 to HMDS45 (Table 2.5). Qualitatively, it is worth noting that the occurrence of polyaromatic C is also ascertained by other investigation methods like electron spin resonance (ESR) and X-photon spectroscopy (XPS) (Chapters 5.2 and 5.3). HMDS44, 1500°C ANNEALED Features are identical to that of the 1400°C-annealed sample described above. The interesting observation is that, contrarily to what happened in sample HMDS42 in which a widespread β-SiC phase is found after annealing at this temperature, no crystallisation of a Si-based phase has occurred yet. The bulk material is still amorphous, and still contains Si, C and N, as identified by electron energy loss spectroscopy (EELS) (Figure 4.37). It is worth noting that no oxygen is detected. An extended EELS study in a VG STEM microscope was specifically performed by N. Brun (Laboratoire de Physique du Solide, Paris, France) on some grains from this sample in order to tentatively characterise further the C content, using a very fine electron probe combined with scanning facility. Quantitative exploitation of the EELS spectra led to a bulk

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Figure 4.37 EELS spectrum (large probe) from HMDS44/1500. Si, C and N are revealed. atomic C content estimated at ~16–20%, most of it being segregated at the grain surface (Figure 4.38(a)) as sp2 C atoms, that is, as a polyaromatic C phase (Figure 4.38(b)). On the contrary, the grain core was found to contain ~4% of sp2-type C atoms (Figure 4.38(b)), giving an inner C/N atomic ratio of ~0.2 (together with an inner N/Si atomic ratio of ~1.3). All these values strongly support the previous observations on the C phase occurrence, type and location, and are fairly consistent with the quantitative data obtained otherwise. In addition, the Si amount in the bulk is more that twice that of the C amount corresponding to the sum of the inner and outer C (Figure 4.38(a)). This and the occurrence of an abundant N content (~45%) show that an Si—(C)—N amorphous phase is also present in the grain in addition to the SiC crystals. HMDS44, 1550°C ANNEALED The SiC phase crystallises first, exhibiting peculiar polyhedral morphologies, most often triangular (sometimes rectangular) corresponding to the primary crystalline form of the cubic (hemihedral) structure of SiC (Figure 4.39). HMDS44, 1600°C ANNEALED Overall crystallisation of both Si3N4 and SiC has occurred. It is likely that the location where Si3N4 and SiC phases crystallise is correlated to the previous occurrence of polyaromatic-carbon-depleted and polyaromaticcarbon-rich grains, respectively. Grain sizes are in the 50–200 nm range (Figure 4.39). Electron diffraction shows that the cubic (3C) structure is probably no longer prevalent, and a lot of other polytypes are found since additional diffraction spots other than 111, 200, 220 and 311 of β-SiC are present, often as stacking faults as indicated by the specific occurrence of aligned spots (Figure 4.40, arrow). The triangular shape of the SiC crystal projection suggests that the growth has occurred in conditions of saturation (i.e. no limitation in the C and Si atom supply) or, at least, in conditions such that feeding atoms have equal access opportunity to the various crystal faces, since the relative growth speeds of the various faces were respected. In such conditions, ‘slow’ faces (like (111)) are actually favoured relative to ‘fast’ faces (like (100)) since the latter require less atoms than the former to be completed.

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Figure 4.38 Data obtained from scanning nanoprobe EELS performed on single grains from HMDS44/1500. (a) EELS quantification of Si, C and N (in at% along y axis) within half a grain (grain radius along x axis) showing that most of the C content is at the grain surface. (b) sp2 (solid symbols) and sp3 (open symbols) C atom content quantification within a grain about 50 nm large, indicating that the surface C is exclusively polyaromatic (sp2). Data from N. Brun, Laboratoire de Physique du Solide, Paris, France.

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Figure 4.39 TEM bright-field image from HMDS42/1600 (SiC phase). Most of the SiC crystals exhibit a triangular projection.

Figure 4.40 Electron diffraction from HMDS42/1600 (SiC phase). A lot of polytypes other than 3C are present, including as stacking faults (arrows). On the other hand, the Si3N4 phase was not investigated, due to the very large size range of the crystals (several micrometers) making them not transparent to the TEM electron probe. However, crystallised Si3N4 was otherwise undoubtedly revealed by X-ray diffraction (XRD).


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Page 83 Finally, it is worth noting that the polyaromatic C phase is no longer found, suggesting an oxidation into CO (rather than CO2, because of the high temperature) due to the release of oxygen consequently to the crystallisation of SiC and Si3N4. Si/C/N/Al-based powders Among the various samples prepared following this route, three were specifically investigated by TEM, asprepared first, then after a 1600°C (1 h isothermal) annealing under N atmosphere. Samples labelled HSAl07, HSAl12 and HSAl09 thus respectively correspond to 0/100, 20/80 and 40/60 NH3/Ar proportions in the gas phase. HSAL07, AS-FORMED The colour of the macroscopic powder at natural light is grey. In TEM, electron diffraction shows that the bulk sample is amorphous (Figure 4.41). Grain size range from ~20 to ~130 nm, with a tendency to a twomode distribution (Figure 4.42). The overall grain morphology is round, although the smallest grains may tend to be more or less slightly polyhedral. Another difference between small (i.e. in the range ~20–60 nm) and large grains (i.e. in the range 50–130 nm) is that the former exhibit some nanometric porosity (Figure 4.43), suggesting some primary difference in chemical composition. No polyaromatic C phase was revealed at the grain surface either by means of dark-field (not illustrated) or lattice fringe imaging (Figure 4.44). However, some primary, nanometric graphene stack entities are possibly imaged (Figure 4.44, arrows). Also, some poorly organised polyaromatic C particle (~15–30 nm) aggregates similar to C blacks (not illustrated) were found as a seldom phase apart from the amorphous phase. HSAL12, AS-FORMED The natural light colour of the macroscopic powder is black. The material differs from HSAl07 by the grain size range, which is more obviously shared into two sub-size ranges

Figure 4.41 Electron diffraction from HSAl07. The whole material is amorphous.

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Figure 4.42 TEM bright-field image from HSAl07. The two-mode grain size distribution is shown.

Figure 4.43 High magnification TEM bright-field image from HSAl07. Small grains (bottom) exhibit a nanotexture while large grains (top) do not. (see Figure 4.42 for reference). The most prevalent one is ~10–30 nm, whose grains are more or less polyhedral and exhibit a nanometric porosity (see Figure 4.43 for reference). The minor one is ~50–100 nm, with grains having round contours and no obvious porosity. The main difference with HSAl07 is that the early crystallisation of a widespread, poorly organised, turbostratic, polyaromatic phase is suggested by a faint C002 halo in electron diffraction patterns (Figure 4.45, arrow), further supported by the occurrence of tiny brighter

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Figure 4.44 TEM lattice fringe image from HSAl07. Primary nanometric graphene stacks are probably showed.

Figure 4.45 Electron diffraction from HSAl12. The whole material seems amorphous. spots into the grains and some bright rims at the grain surface in the related dark-field images (Figure 4.46, arrow). Likewise, the early crystallisation of SiC has also started in the small grains, as revealed by the SiC-specific dark-field images (left side in Figure 4.47, arrows) while large grains remain amorphous (right side in Figure 4.47). Indeed, the fact that no such crystals appear in ‘C-specific’ dark-field images (Figure 4.46) ascertain that crystals are actually SiC, not Si3N4 (see Section 4.3). It is worth noting that, unlike what was found for SiCxxx- or HMDSxx-type materials, the SiC crystallisation does not start according to an overall, widespread germination mechanism resulting in a homogeneous dissemination of nanometric crystals in

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Figure 4.46 TEM dark-field image from HSAl12. ‘C-specific’ (DF1) imaging conditions. Bright dots within small grains and bright rims at grain surface appear, indicating the occurrence of a poorly organised, polyaromatic C phase. the material, but rather corresponds to the local occurrence of the appropriate nucleation-then-growth conditions, resulting in the random but seldom occurrence of SiC crystals. The amount of crystals is not enough to appear in the diffraction patterns (Figure 4.45), but the β structure can be ascertained from the dark-field data. HSAL09, AS-FORMED The natural light colour of the macroscopic powder is heterogeneous with grey to light brown areas. The overall TEM features of the sample are very similar to HSAl12 regarding the two-class grain size distribution (see Figures 4.42 or 4.47 for reference), the early occurrence of seldom SiC crystals in the small grains, the absence of Si3N4, and the possible occurrence of a widespread, poorly organised, polyaromatic phase. However, the latter is less ascertained since no C rims at the grain surface were found. Some C black type aggregates are also found (not illustrated).

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Figure 4.47 TEM dark-field image from HSAl12. ‘SiC-specific’ (DF2) imaging conditions. SiC crystals are found locally in small grains, already several nanometres large. HSAL07, 1600°C ANNEALED The macroscopic natural light colour of the sample has turned from grey to black. The material is quite homogeneous, and a single morphology of round or slightly polyhedral grains was found in TEM images (see Figure 4.22 for close enough reference). The grain size range is ~30 to ~60 nm, therefore, significantly narrower than in the as-prepared materials, which has to be related to some weight loss during the ceramisation process. Appropriate darkfield imaging (not illustrated) shows that each grain is actually an aggregate of round single crystals (~2–25 nm) with a non-symmetrical grain size distribution since a majority of them are less than ~10 nm. Electron diffraction identifies the crystals as SiC with the prevalent β (3C) structure since reflections 111, 220 and 311 are the only ones being revealed as diffraction rings, although some diffraction spots beside these rings appear, showing the seldom occurrence of other polytypes (see Figure 4.8 for reference). Correspondingly, some crystals exhibit stripes in dark-field images, revealing polytype sequence as stacking faults. The association of the SiC crystals into larger grains implies the occurrence of a binder, that is, an intergranular amorphous phase to ensure the grain cohesion, obviously containing N and Al, in addition to some Si and possibly C. A polyaromatic C phase is now revealed at the grain surface (though not systematically), either by dark-field (not illustrated) or lattice fringe imaging (see Figure 4.37 for reference), which is responsible for most of the polyhedral contours of the grains. The cohesion of the aggregates is ensured by an intergranular amorphous phase, that X-ED Sreveals to contain at least Al and O, in addition to probably Si and possibly C. Also, traces of K, Ca and Na are found (Figure 4.48).

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Figure 4.48 X-EDS spectrum of the SiC-containing grains from HSAl07/1600. The occurrence of Al and O in addition to Si and C suggests that these elements are present in the composition of an intergranular amorphous phase. No N-containing amorphous or crystallised phases are found, while N is more abundant than O in the material according to the bulk chemical composition (Table 2.6). This emphasises the fact that the qualitative comments about the respective proportions of the various phases are only valid for the material once prepared for TEM investigation. Actually, when a phase is present as very large grains (i.e. micrometric) together with another phase with smaller grains (sub-micrometric), the suspension method used to prepare an electron transparent aliquot of the materials favours the prevalence of the latter relative to the former, with no respect to the original proportions. It is therefore possible that large grains of Si3N4 (or SiAlON) were present in the material, though not found in the TEM preparation. HSAL12, 1600°C ANNEALED The natural light colour of the macroscopic powder has turned from dark to light grey-green. The whole material is crystallised. The SiC phase is the most frequent (in the TEM preparation of the material, see the end remark of previous section) and appears as grains ~20–100 nm large, exhibiting systematic triangular projections (see Figure 4.39 for reference). Electron diffraction shows that the cubic (3C) structure is not prevalent since the related 220 and 311 reflections barely form complete rings in the diffraction patterns (see Figure 4.40 for reference). Many other diffraction spots corresponding to other lattice distances are found instead, therefore, revealing the occurrence of other polytypes. Crystals are frequently coated with an amorphous phase, few nanometres thick (not shown), believed to be different from mere silica since containing at least Al and O (in addition to possibly Si and C) according to X-EDS spectra (see Figure 4.48 for reference), which helps in bonding the crystals to each other (not shown, similar to Figure 4.8). Large, somewhat elongated, micrometric crystals, therefore, not electron transparent, were also found (Figure 4.49). They are free of any amorphous coating, and provide only Si and N with a small contribution of Al and traces of O in X-EDS spectra (Figure 4.50), suggesting a SiAlON-type composition. Other crystals exhibit a needle-like morphology (Figure 4.51). Some observations suggest that they somehow originate from cleavages of the larger, massive

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Figure 4.49 TEM bright-field image from HSAl12/1600. Large, dark grains are SiAlON crystals, while needlelike grains are Si3N4 crystals according to X-EDS spectra. Arrow shows that the latter probably originates from the former by cleavages.

Figure 4.50 X-EDS spectrum from large, dark crystals from HSAl12/1600. Al is shown, probably together with O as a shoulder at the right side of the N peak, suggesting a SiAlON composition. crystals (Figure 4.49, arrow). Correspondingly, X-EDS spectra similarly exhibit Si and N, however with the systematic absence of Al and the probable absence of O (Figure 4.52) suggesting that they are genuine Si3N4 crystals, while the needle-like morphology would indicate a structural transformation from α to β. Quantitative mismatches with the phase proportions estimated from XRD (Table 4.1) originate from the segregation effect mentioned at the end of the previous section.

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Figure 4.51 TEM bright-field image from HSAl12/1600. Arrow shows that even needle-like crystals are cleaved into smaller needle-like crystals.

Figure 4.52 X-EDS spectrum from needle-like crystals from HSAl12/1600. The absence of Al (and O) indicates a Si3N4 composition. Again, like for material HMDS44/1600, evidence for a polyaromatic C phase was no longer found, suggesting an oxidation into CO (rather than CO2, because of the high temperature) due to the release of oxygen consequently to the crystallisation from the primary O-containing amorphous phase.

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Page 91 Table 4.1 Phase proportions as determined from XRD spectra in 1600°C annealed HSAlxx-type samples Sample reference Results from XRD HSAl07/1600°C β-SiC + α-SiC(ε) HSAl09/1600°C α-Si3N4 + β-Si3N4(ε) HSAl11/1600°C α-Si3N4 + β-Si3N4 HSAl12/1600°C α-Si3N4 + β-Si3N4 + β-SiC HSAl13/1600°C β-SiC + α-SiC HSAl14/1600°C α-Si3N4 + β-Si3N4 + β-SiC HSAl15/1600°C α-Si3N4 + β-Si3N4 + β-SiC HSAl16/1600°C α-Si3N4 + β-Si3N4 HSAl17/1600°C α-Si3N4 + β-SiC + β-Si3N4(ε) HSAl18/1600°C β-SiC + α-SiC Note Underlined = most abundant. (ε) = less than ~5%. HSAL09, 1600°C ANNEALED The natural light colour of the macroscopic powder has turned from light grey-brown to light brown. The same phases (i.e. SiC, and Si3N4 or SiAlON) as in HSAl12/1600 are found, however with different proportions. In addition, some of the primary amorphous phase(s) remain(s), as minor small and large round grains (see Figure 4.42 for reference), whose chemical composition contains Si, N (and O) according to XEDS spectra (not illustrated). The most prevalent morphology is made of aggregates of polyhedral crystals about 20–60 nm large, whose projection are often triangular (Figure 4.53), identified as β-SiC by electron diffraction (Figure 4.54). Some but few polytypes are also present. This SiC phase is associated with an amorphous phase coating the crystals and somewhat acting as an intergranular phase and ensuring the aggregate cohesion. As for the other HSAlxx materials investigated though with various element proportions, the chemical composition of the amorphous phase at least includes Al and O, in addition to possibly Si and C according to X-EDS spectra (see Figure 4.48 for reference). Massive, large, therefore, non-transparent crystals are also found as in material HSAl12/1600 (see Figure 4.49 for reference). A difference is that the morphological transformation into needle-like crystals has barely occurred yet (Figure 4.55) and the latter are seldom, possibly indicating that the α to β transformation has just started. Another difference is that either massive or needle-like crystals exhibit Al and O in addition to Si and N according to X-EDS spectra (Figure 4.56) suggesting a SiAlON-type composition. Mismatches with the phase composition estimated from X-rays originate from the sample preparation for TEM investigation. Conclusion regarding ceramic nanopowders from liquid precursors SI/C/N-BASED MATERIALS Due to their high hydrogen content which makes the materials behave like Si-based polymer-derived ceramics, the as-prepared materials remain mainly amorphous up to annealing temperatures as high as ~1500°C. Such a thermal stability is induced by the complexity

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Figure 4.53 TEM dark-field image (‘SiC-specific’ (DF2) imaging conditions) from HSAl09/1600. Many of the crystals exhibit a triangular projection.

Figure 4.54 Electron diffraction from the SiC phase from HSAl09/1600. The 3C polytype is highly prevalent.

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Figure 4.55 TEM bright-field image from HSAl09/1600. Large, dark grains are SIALON crystals according to XEDS spectra.

Figure 4.56 X-EDS spectrum from large crystals from HSAl09/1600 (also valid for the needle-like crystals). The occurrence of Al and O indicate a SiAlON composition. of the chemical system (Si—C—N—O—(H)) which tends to unfavour the atomic rearrangements at the solid state. Indeed, the previous formation and segregation of tetrahedra with appropriate compositions (like SiC4) is necessary to subsequently allow the various phases (like SiC) to crystallise.


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Page 94 The C/N atomic ratio of the as-prepared materials appears to be the leading parameter to determine whether SiC and/or Si3N4 will finally crystallise at some point of the annealing treatment. Starting from a nearly similar Si content, a high C/N content favours the early formation of a SiC phase within a surrounding amorphous matrix, in which heteroatoms (O, N) are subsequently concentrated as an amorphous Si (carbo) oxinitride phase. On the contrary, a low C/N content helps in maintaining the amorphous structure of the bulk at a higher annealing temperature, lately inducing a dramatic nucleation-then-growth of both SiC and Si3N4 phases (Si3N4 only for the lowest C/N ratios). Because nucleation occurs at higher temperatures (>1500°C) with low or no temperature gradient, growth mechanisms (and kinetics) are different, possibly due to a higher mobility of atoms allowing crystals to grow freely, therefore, giving rise to primary crystal morphologies instead of the round, isometric morphology (like that resulting from the flash-temperature nucleation-then-growth mechanism occurring for the gaseous-precursor-based nanopowders). Whatever, polyaromatic C is always the first phase to crystallise, though with a low-dimensionality – turbostratic – structure, according to a common behaviour of Si-based polymer-derived ceramics (Monthioux and Delverdier 1996). Consequently, the amount of it is directly related to the C content of the as-prepared material. SI/C/N/AL-BASED MATERIALS Whether NH3 is present or not in the gas phase, all the as-prepared materials investigated are nearly all amorphous but tend to a two-morphology differentiation with small, nanoporous grains, and larger, nonporous grains, which could be related to the subsequent SiC/Si3N4 phase differentiation. Effect of the 1600°C annealing • Average grain sizes decrease, revealing a weight loss subsequent to the ceramisation process. The crystallised phases mainly depend on the N content of the chemical system. • Not mentioning the polyaromatic C phase, which is not systematic, the SiC phase crystallises prior to the Si3N4/SiAlON phase. • SiC crystals are always coated with an amorphous (Si?)—Al—O phase which is an interesting feature for subsequent attempts of high-temperature sintering. • Correspondingly to the high annealing temperature, the SiC structure is always a mixture of various polytypes (though with various proportions from a material to the other), present as stacking faults rather than as single crystals. Effect of the presence of NH3in the precursor gas phase • After annealing, N-depleted chemical system, induced by the absence of NH3 in the precursor gas phase, on one hand allows the occurrence of a polyaromatic C phase and on the other hand allows the formation of a crystallised SiC phase only (i.e. no Si3N4). • Adding NH3 in the precursor gas phase induces the early crystallisation of SiC. However, the nucleation step does not occur as a widespread germination but occurs locally. Such a heterogeneous behaviour has to be related to a heterogeneity in chemical composition, possibly originating in the non-homogeneous incorporation of N and/or H into the material. On the contrary, the absence of NH3 provides a quite amorphous material.

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Page 95 • Adding NH3 in the precursor gas phase allows the required stoichiometry to be reached, resulting in the rapid crystallisation of a SiAlON-type phase after annealing. Meanwhile, the high temperature at which the overall SiC crystallisation occurs induce such an atom mobility than the respective intrinsic growth rates of the various crystal faces are respected. Effect of varying the NH3concentration in the precursor gas phase • For the lowest and the highest supply in NH3, the prevalent polytype of the resulting crystallised SiC phase after annealing is 3C, corresponding to the β-cubic structure. An intermediate NH3 supply results in a more rapid growth of the SiC phase, for which the other polytypes are correspondingly prevalent. Such a nonmonotonous effect of the increasing N content of the chemical system regarding the behaviour of the material during annealing was already observed for Si-based polymer-derived ceramics and originates from the competition between antagonist mechanisms (Delverdier et al. 1992a). • Increasing the N proportion slows down the kinetics of structural rearrangement; inducing the highest amount of remnant amorphous phases in the N-richest system after annealing. Another reason is also the early crystallisation of SiC prior to Si3N4, which induces a local excess of N atoms relative to the Si atoms left with respect to the required Si3N4 stoichiometry, which then gather into an amorphous phase. • The crystallisation of a nitride phase occurs for a 20% and 40% input of NH3 in the precursor gas phase. In both materials, the high annealing temperature tends to induce the overall crystallisation of massive SiAlON-type crystals rather than Si3N4. In addition, seldom SiAlON needle-like crystals are found in the latter sample (40% NH3), while numerous Si3N4 needle-like crystals are found in the former (20% NH3). Assuming that the massive and needle-like morphologies correspond to α and β structures respectively, these data suggest that the structure change might be somehow controlled by the N content of the chemical system, independently from the extent of Al—Si and O—N substitution. 4.2 As-prepared sintered composites Some of the as-prepared, sintered, micro–nano-type composites, whose detailed processing conditions and mechanical properties are described in Chapter 7, were investigated by TEM to investigate further the structural and textural relationships between the composite components. Related sample preparation prior to TEM investigations and TEM operating conditions are reported in Section 4.3. Sintering temperature was 1650°C for all the composites (1600°C for the bulk Si3N4 material). All the materials (i.e. bulk Si3N4 and composites) are added with the regular sintering aids, Al2O3 and Y2O3, added in the proportions of 3% and 6%, respectively. For technical reasons, all the materials but the bulk Si3N4 were sectioned, thinned and observed in a direction parallel to the hot pressing direction. In the following, after describing the monolithic Si3N4 material, composites will be presented according to the decreasing amount of theoretical SiC content (Table 7.1). 4.2.1 Bulk Si3N4 Monolithic sintered Si3N4 – that is, not added with any nanopowder – was studied as a reference material. The sintering temperature at 1600°C provides the highest densification rate

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Figure 4.57 TEM bright-field image from the SiAlON monolithic material. (see Figure 7.6(b)). The sample was investigated in a plane perpendicular to the hot-pressing direction. This plane also contains the direction of mechanical testing. The material appears highly densified (Figure 4.57). Due to the high sintering temperature and the presence of oxides, a dramatic chemical and structural transformation has occurred, inducing a high β-SiAlON (Si6 −zAlzOzN8−z) crystal content (see Section 7.2), about 70% according to calculations based on XRD data (Rossignol 1995), as compared to 100% of α structure in the starting powder (see Section 7.4.1). Correspondingly, a majority of the grains are more or less elongated, some exhibiting an obvious needle-like morphology. This is consistent with previous TEM observations which demonstrated that oriented pressure sintering is able to statistically induce a preferred orientation perpendicular to the pressing direction as far as non-isometric crystals – like high-temperature β-SiAlON or β-Si3N4 – are concerned (Champion et al. 1992). Of course, aspect ratios revealed by TEM are just apparent, since depending on the crystal orientation relative to the sectioning direction for the non-isometric grains. Actually, many of the nearly isometric crystal morphologies are likely to also derive from β-type crystals, that is, about 250–300 nm, while the maximal value observed for the elongated crystal dimension is about 2 μm. Crystals most often exhibit a uniform grey contrast, indicating that they are mainly free of stacking faults. 4.2.2 Si3N4/SiC173 composites The composites investigated were initially prepared with 20% and 40% of SiC173 nanopowder, respectively (Table 7.1). Considering that the starting SiC173 nanopowder is made of quasi pure SiC (with some content of polyaromatic C, see section on Si/C-based powders), such proportions also correspond to the theoretical proportions of the SiC phase in the final composites (Table 7.1).

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Page 97 No obvious discrepancy was found between composite Si3N4/SiC(20) and composite Si3N4/SiC(40) as observed in TEM. The following comments are therefore valid for both materials. Low magnification TEM images are consistent with the theoretical SiC phase proportions (Figure 4.58), and emphasise the size difference between the SiAlON and SiC phases. SiAlON crystals appear as large, polyhedral grains, quite similar to that of the monolithic material (Section 4.2.1). Although diffraction patterns indicate a majority of β structure consistent with the high-temperature sintering (Figure 4.59(a)), a reflection specific of the α phase (at 0.432 nm) is still found sometimes (Figure 4.59(b)), indicating that the structural α to β transformation rate is not 100%, which is consistent with the theoretical calculations (Table 7.3). However, considering that XRD was not able to reveal any α structure, the actual amount of α crystals is surely less than the 11% predicted from calculations (and probably less than 7%, the lowest percentage determined by X-ray in composite Si3N4/SiCN35(22.2)). Most of the crystal facets appear as straight lines, as for single crystals when allowed to grow freely (Figure 4.58). However, crystal facets of SiAlON crystals at the contact of SiC grains often appear notched (see Figure 4.66 for reference) due to the presence of SiC grains at the interface, confirming that SiC grains were not allowed to melt or soften during the sintering process, unlike the growing SiAlON crystals. It is worth noting that only few elongated crystals are found. Calculating an average apparent crystal size is therefore possible: ~470 nm

Figure 4.58 TEM bright-field image from composite Si3N4/SiC(20). The SiC crystals (small, round grains) are gathered into aggregates.

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Figure 4.59 Electron diffraction from composite Si3N4/SiC(20). (a) Most of the diffraction spots are characteristic from the β-SiAlON phase. (b) However, some specific diffraction spots (arrow) reveal the minor occurrence of the α structure as well. for composite Si3N4/SiC(40). The slightly lower value for Si3N4/SiC(20) is not significant. Comparison with the monolithic Si3N4 material regarding the evolution of the grain size is limited since it was not investigated along the same section plane. SiC crystals appear as small, round crystals, most often gathered into agglomerates instead of being randomly dispersed all around the SiAlON crystals (Figures 4.58). Contrarily to the latter, SiC crystals exhibit numerous, non-periodical, dark contrasts revealing the occurrence of polytypes as stacking faults. Indeed, the β to α (i.e. polytypes other than β) SiC transformation has slightly progressed, as ascertained by the more frequent occurrence of specific features in the diffraction patterns, like the occurrence of diffraction spots aside the regular diffraction rings for the β structure, and of some aligned spots corresponding to the presence of stacking faults (Figure 4.60 to be compared with Figure 4.3). On the contrary, the average SiC grain size has not progressed significantly compared to the starting powder (see section on Si/C-based powders). This indicates that the growth mechanism by coalescence, though an usual process at such temperatures (Monthioux and Delverdier 1996), was prevented, probably because of the presence of the intergranular phase (Figure 4.62), as opposed to the starting nanopowder (see Section 4.1.1). Some SiC grains are also found possibly enclosed into SiAlON crystals (Figure 4.61). However, high magnification images are not sufficient to ascertain if the SiC crystals are actually enclosed into the SiAlON crystals or if they merely are superimposed to them in the projected image as allowed by the thickness value of the TEM preparation (~100 nm). Proofs of enclosing of SiC crystals into SiAlON crystals will be brought using darkfield imaging (see Figure 4.66 to come). The polyaromatic C phase associated with SiC grains revealed in the starting powder is no longer found, suggesting that it might have burnt into CO during the sintering process. An amorphous, intergranular phase is found, gathering the whole yttrium content of the composite (in addition to some content in the other elements present, Si, Al, O, N) as revealed

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Figure 4.60 Electron diffraction from a SiC grain agglomerate in composite Si3N4/SiC(40). Though reflections from the β structure are still prevalent, many other polytypes are present.

Figure 4.61 High magnification TEM image from composite Si3N4/SiC(20). A SiC grain (faulted) seems to be enclosed in the SiAlON crystal. by the darkness of the grey contrast associated to this phase, originating from the intense electron inelastic scattering induced by the high Z value for yttrium (Z = 39). However, most of the intergranular phase is found within SiC grain agglomerates either filling the voids between grains (Figure 4.62) or at the grain interfaces (Figure 4.62, arrows). On the contrary, the intergranular phase is absent between SiAlON crystals (Figure 4.63) or might eventually

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Figure 4.62 High magnification TEM image from composite Si3N4/SiC(20). Spaces between SiC grains or at the SiC/SiAlON interface appear dark due to the presence of the yttrium-containing amorphous intergranular phase. be present as some very thin –


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Figure 4.63 TEM lattice fringe image from composite Si3N4/SiC(20). The interface between SiAlON crystals is free of interfacial phase. Si3N4/SiCN35(44.4) composite Many of the morphological, textural and structural features described above for the Si3N4/SiC173 composites are valid for this composite. One discrepancy, however, concerns the SiC phase, which now exhibits two kinds of morphology. In addition to the regular, round, major morphology previously described (with about the same grain size range, though slightly lower, 10–60 nm), polyhedral crystals whose sizes range from ~50 to ~150 nm are found as a minor phase (Figure 4.64). Correspondingly to this larger average size, the β to α transformation is more pronounced for this second morphology (see Figure 4.40 for reference, to be compared with Figure 4.60 for the major SiC phase morphology, and with Figure 4.24 for the structural state of the SiC phase in the as-prepared SiCN35 nanopowder). This morphological differentiation indicates at least two distinct sources for the SiC phase, one originating from the pre-existing SiC crystals, the other possibly originating from the crystallisation of the pre-existing Si—C—O—N amorphous phase and/or the reaction of the pre-existing Si phase with any dispersed C atoms. Another discrepancy concerns the SiAlON phase, whose average apparent crystal size is ~360 nm, that is, smaller than for composite Si3N4/SiC173(40). Si3N4/SiCN35(11.1) composite A major difference with the three composites described so far is brought by the SiC phase, which is now more often located between the SiAlON crystals, that is, clearly located in intergranular position (Figure 4.65). Beside, in this sample as in all the other SiC-containing composites described in this chapter, SiC crystals structurally enclosed into SiAlON crystals are found, as clearly demonstrated by the example of the bright-field/dark-field comparison of Figures 4.66(a) and (b). Two SiC crystals supposedly enclosed into a SiAlON crystal are

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Figure 4.64 TEM dark-field image (DF2 position) from composite Si3N4/SiCN35(44.4). Example of the minor SiC morphology, exhibiting larger sizes and polyhedral contours, found in addition to the major, regular, round morphology. arrowed in the bright-field image of Figure 4.66(a), to be compared with the specific dark-field image of Figure 4.66(b), whose objective aperture position is adequately chosen so that none of the possible diffracted beams from crystallised SiC are allowed to participate in the image, while the diffracted beam from the SiAlON crystal selected is allowed (see Section 4.3). On Figure 4.66(b), the SiAlON crystal appears dark at the very place of one of the SiC crystal (solid arrow on Figure 4.66(a)), while it appears bright at the very place of the other SiC crystal (open arrow on Figure 4.66(a)) making the latter invisible. The former SiC crystal (solid arrow) is therefore really enclosed into the SiAlON crystal, while the latter SiC crystal (open arrow) is merely superimposed to it in the thickness of the TEM preparation. Contacts between enclosed SiC and surrounding SiAlON crystals were always found direct, that is, with no interfacial phase. Though enclosing of SiC crystals into SiAlON crystals is thus ascertained and supported by the literature (Niihara et al. 1990; Sasaki et al. 1993), examples are actually few in our materials, making doubtful that such nanometric SiC crystals are able to play a substantial role in the germination or phase transformation of the SiAlON crystals in the composites investigated. Meanwhile, Figure 4.66(b) also illustrates how the development of the crystal faces is hindered by the presence of SiC grains in interfacial position, inducing a highly indented topology of the related crystal facets. Decreasing the SiC content in the Si3N4/SiCN35 composites help the SiAlON crystals to develop further, since the average apparent SiAlON crystal size is ~395 nm (as opposed to ~360 nm in composite Si3N4/SiCN35 (44.4)).

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Figure 4.65 TEM bright-field image from composite Si3N4/SiCN35(11.1). Light grey contrast-free areas are SiAlON crystals. SiC grains are in interfacial position between SiAlON crystals.

Figure 4.66 Composite Si3N4/SiCN35(11.1). (a) TEM bright-field image. SiC grains are gathered into aggregates. Arrows show examples of SiC grains possibly enclosed in a SiAlON crystal. (b) TEM dark-field image. Imaging conditions (DF1) are adequately chosen to image SiAlON crystals (bright area), while SiC is excluded. One of the SiC grain (solid arrow) is still visible and is therefore really enclosed into the SiAlON crystal, while two other SiC grains (open arrows) have vanished and are therefore merely superimposed to the SiAlON crystal in the projected image. Note how the SiAlON crystal facets are notched due to the contact with SiC grains which have hindered the normal growth. The amorphous intergranular phase is revealed between SiC grains as a slightly scattering phase.

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Figure 4.67 TEM bright-field image from composite Si3N4/SiCN29(20). Many elongated crystals are found. 4.2.4 Si3N4/SiCN29(20) composite This composite was prepared with 20% of SiCN29 nanopowder, providing a theoretical SiC phase content of 1.7% only. The textural description of the composite is basically similar to that of composite Si3N4/SiCN35 (11.1) described above. SiC grains are again often disseminated in intergranular position between SiAlON crystals. However, SiC crystals are found to exhibit a single morphology, which is consistent with the fact that the SiC source is unique, since no pre-existing SiC nor Si phases were present in the starting nanopowder (see Section 4.1.2). Also, SiC crystals are less numerous, consistently with the calculation, allowing SiAlON crystals to grow and develop more freely than in other composites. Correspondingly, SiAlON crystals very frequently exhibit elongated sections, as opposed to the other composites (Figure 4.67), making an average apparent crystal size difficult to calculate. Grain section range is wide, up to several micrometers for maximal lengths, and 200–400 nm for maximal widths. Although the direct comparison is not allowed with the monolithic Si3N4 material due to the fact that TEM preparations were not performed following the same direction (i.e. perpendicular and parallel to the hot-pressing direction for the monolithic and the composite materials, respectively), the overall impression is that the average crystal size is smaller for the composite (Figure 4.67) than for the monolithic material (Figure 4.57), which is consistent with results from SEM investigations (see Figures 7.8(b) and 7.10, respectively for comparison). 4.2.5 Conclusion regarding as-prepared sintered materials An overall statement is the elimination of the pre-existing polyaromatic C phase – if any – during the sintering process, due the carbo-thermal reduction reaction with the oxide phases added (or pre-existing). This is likely to induce some gas release (CO) possibly responsible for some of the remaining porosity.

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Page 105 Pre-existing SiC grains are stable in the material being sintered, more or less preventing the SiAlON crystals to develop freely. Specifically, SiAlON crystal facets appear flat at the SiAlON/SiAlON contacts, while they appear notched at the SiAlON/SiC contact. When in high proportions (>5% according to the theoretical calculations), the SiC phase tends to gather into agglomerates apart of the SiAlON phase, instead of being dispersed in intergranular position. Agglomeration could be caused by the occurrence of the surrounding minor phase (polyaromatic C in case of SiC173 nanopowder, and an amorphous silicon oxycarbide phase in case of SiCN35 nanopowder), preventing a convenient dispersion during the mixing step prior to making the green bodies (see Chapter 7). Such agglomerates, because they are made of small, round grains having limited surfaces of contact, exhibit a high intergranular porosity allowing to somehow trap most of the yttrium-containing amorphous phase, which is then no longer available to form interfacial films between SiAlON crystal facets. A poor superplastic behaviour is therefore expected for these composites. On the contrary, a low content in SiC nanograins is favourable to a better dispersion in intergranular position between SiAlON crystals. However, SiAlON crystal faces are then prevented to grow freely, the hindering presence of the SiC grains at the SiAlON crystal interfaces while developing from the liquid phase inducing indentations in the related crystal faces. SiC grains thus act as ‘keystones’ between SiAlON grains, which is also detrimental to a superplastic behaviour. Nano/nano composites are therefore expected to be preferable to micro/nano composites regarding the ability to creep without developing cracks. Further and extended comments will be found in Chapter 7. 4.3 TEM sample preparation prior to investigations and operating conditions 4.3.1 Preparation of samples In order to be electron transparent at 120 kV acceleration voltage, materials investigated should exhibit a thickness ideally lower than 100 nm in the direction parallel to the electron beam. Ceramic nanopowders therefore require limited efforts for TEM preparation, considering the convenient initial grain sizes. On the contrary, composites require a specific thinning process. Powders Aliquots of each ceramic nanopowder were dispersed into ethanol under sonication for several minutes in order to make a suspension from which a drop was then deposited on a regular (3 mm) 200 mesh copper microgrid previously coated with a holey or a lacey amorphous C film. For samples suspected to be possibly hydrolysable, dry preparation – that is, direct deposition of the nanopowder onto the microgrid – were sometimes also made for double checking. When large powder agglomerates were formed unable to be dismantled by the wet sonication step, gentle grinding using a boron carbide mortar was previously made (use of a silica mortar should be avoided, due to its lower hardness relative to SiC, which is likely to provoke some input of silica into the sample being prepared). The main limitation of such a preparation procedure is the possibility to induce a gravimetric discrimination – the phases exhibiting much higher density, or much larger grains being likely to sink to the bottom of the flask.

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Page 106 Sintered materials A slice about 100 μm thick of each sintered materials was obtained using a diamond-coated slow circular-saw (preferably to a diamond-coated tungsten-wire saw, whose efficiency was limited in this specific case). The slice dimensions was then adjusted to fit in a 3-mm wide copper aperture, then glued on a stainless steel support using a thermofusible resin. Then the thickness of the slice was decreased down to ~50 μm using a mechanical polishing procedure involving submicrometric diamond powders. The slice was then reversed upside down in order to complete the polishing/thinning process through the other side of the slice, hopefully down to ~20 μm. The slice was then glued onto a 3-mm copper aperture, then ion-milled both sides for several hours. Acceleration voltage for the argon ions was ~4kV, and ion gun direction angles with respect to the sample surface were 12–15°. A hole was created, whose edges were suitable for the TEM investigation. All the micro/nano composites investigated (specifically those with a high SiC content) were more or less found to specifically exhibit a very high hardness making them unusually difficult to slice during the first step of the procedure (diamond-enhanced sawing), as opposed to the behaviour of regular micro/micro-type composites (such as in Champion et al. 1992). On the contrary, the micro/nano composites were found to dismantle more easily than regular micro/micro composites at the end of the polishing/thinning procedure. A possible explanation is that the primary specifically high hardness is due to the lack of amorphous intergranular phase at the SiAlON crystal interfaces, making the crystals collectively supporting the shearing stresses developed during the sawing procedure as a monolithic single crystal would do. On the other hand, the removal of SiC grain agglomerates is easier during the next polishing/thinning step due to their high content in amorphous phase. This is likely to promote the earlier dismantling of the composite with respect to micro/micro composites, as soon as a critical limit (~50 μm) for the slice thickness is reached. 4.3.2 TEM operating procedure We do not intend to describe all about the principle of electron imaging. The goal of this section is merely to provide to the non-microscopist reader the keys needed to understand the comments reported in this chapter. Basics and more about the principles of TEM can be found in Williams and Carter (1996). About dark-field mode imaging Contrasted bright field and lattice fringe are two regular imaging modes involving the use of apertures with suitable opening diameters placed in the back focal plane of the objective lens. For the non-microscopist reader, considering the electron microscope as a powerful, high-resolution magnifying lens able to image down to the lattice planes of crystals is enough to understand the kind of comments made in this chapter. Likewise for the selected area electron diffraction, whose principles are similar to that of XRD. However, further explanation might be useful to fully understand the data brought by the dark-field imaging mode. Although a TEM may contain seven or more electromagnetic lenses, its main quality and mode abilities depend on the objective lens. As for regular, optical glass lenses, electromagnetic lenses are characterised by three planes: the focalisation plane (where the specimen should be located for its image to be focused on the image plane), the focal plane (specifically the back focal plane, which contains the electron diffraction patterns in the case of the objective lens), and the image plane, where the enlarged image of the specimen is formed

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Figure 4.68 Sketch of the electron beam paths in the objective lens of a microscope. (a) Bright-field image, (b) dark-field image.

Figure 4.69 Examples of various positions of the objective aperture relative to the diffraction pattern of a multiphase material (β-SiC + turbostratic C) used to perform specific dark-field imaging. (Figure 4.68(a)). Placing an aperture (so-called objective aperture) in the back focal plane of the objective, therefore, allows specific electron beams to be selected, that is, the image is built with some electron beams only (Figure 4.68(b)). For instance, selecting a specific diffracted beam with the objective aperture results in an image whose bright parts indicate the regions of the specimen from which the selected diffracted beam is emitted, while everywhere else, in or around, the specimen appears dark (or nearly). It is then already quite clear that, by successively selecting adequately various positions for the objective aperture, sequential dark-field images can be formed which can be specific of different orientations for a single crystallised phase, or of different phases within a multiphase material, or both. This principle is illustrated in Figure 4.69, with the example of a material

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Page 108 containing both β-SiC and polyaromatic, turbostratic C. The related theoretical diffraction patterns are both represented. Positioning the aperture in DF1 allows part of the C002 diffracted beams to go through, and to build the image of polyaromatic entities – if any – whose (001) lattice planes are oriented both parallel to the electron beam (i.e. graphenes are seen edge-on) and parallel to the E–W direction (or nearly, since the width of the aperture is such that it admits an arc instead of a single dot). Then, positioning the aperture at 90° along the same diffraction ring (position DF1⊥) will image other polyaromatic entities – if any – oriented S— N, that is, perpendicularly to the former. The related two images are exclusive from SiC crystals and are therefore specific of the C phase. SiC crystals will in turn be able to be imaged by using position DF2. In that case, rotating the aperture relative to the diffraction pattern all along the SiC111 diffraction ring (meanwhile also rotating along the SiC200 ring, which is very close) will successively reveal any SiC crystal oriented so that they present (111) or (100) reticular planes parallel to the electron beam. It is worth noting that the DF2 position is not really specific of a SiC phase, strictly speaking, since the aperture width is such that the 10 diffraction band of turbostratic C is also admitted, allowing the related electrons to participate in building the image (using smaller apertures would be too much detrimental to the resolution). Graphene stack entities oriented perpendicularly to the electron beam (i.e. lying flat on the specimen support grid) are thus supposed to be imaged together with the SiC crystals. Practically, due to the intensity of the SiC111 reflection, by far much higher than for the C10, DF2 images are considered specific of a SiC phase in most cases. Because the objective aperture is placed in various positions relative to the diffraction pattern, such dark-field imaging sequences are said to correspond to a radial and azimuthal exploration of the reciprocal space. For this reason, the objective aperture apparent widths are given in nm−1. TEM operating conditions Bright-field, dark-field, lattice fringe and selected area diffraction mode images were obtained using a Philips EM400 microscope operating at 120 kV and equipped with a high-resolution stage providing a point resolution better than 3 Å, with a tungsten wire as electron source. The width of the aperture used for the contrasted bright-field and lattice fringe imaging was 8nm−1, while that used for dark-field imaging was 2 nm −1. According to the Abbe relation where λ is the wavelength and a the half-width of the aperture (in radians) using an objective aperture as small as 2 nm−1 reduces the resolution ρ to ~6 Å. Dark-field imaging was carried out most often using DF1 (aperture centred at 2.2nm−1) and DF2 (aperture centred at 4.3nm−1) positions described in the previous section. Considering the width of the aperture opening, the radial tolerance regarding the frequency range admitted for both positions are 1.3–3.3nm−1 and 3.3–5.3nm−1, respectively. Likewise, the azimuthal deviations relative to the centring position due to the aperture opening are ±24° and ±15°, respectively. Attention has to be paid that, depending on the phases present in the material investigated, the specificity of the positions may vary. For instance, the powder diffraction pattern for a Si3N4 phase exhibits numerous diffraction rings (Figure 4.20 for instance), making Si3N4 crystals likely to be imaged in any possible position for the objective aperture, including DF1 and DF2. Also, other aperture positions were sometimes used, for specific purposes. Finally, though they do not provide diffracted beams strictly speaking, but are anyway able to scatter electrons in some specific directions because of short-range structural periodicities (often around

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Page 109 the most intense diffracted beam for the related crystallised material), so-called amorphous materials are likely to also appear in dark-field images. For instance, amorphous silica may be revealed in DF1 position, providing an intense halo at 0.44 nm−1 (Chaudhari et al. 1972), while the most intense reflection for any form of crystallised silica is located at ~0.41 nm−1. X-EDS spectroscopy was performed using a Philips CM20 microscope equipped with a high-resolution goniometric stage (super-twin stage) and an EDAX detector (Li-doped silicon diode), protected by an ultrathin window allowing the detection of chemical elements including light elements down to C. References Audier, M., Oberlin, A., Oberlin, M., Coulon, M. and Bonnetain, L. (1981) Carbon, 19, 217. Badami, D. V. (1965) Carbon, 3, 53. Buffat, P. and Borel, J.-P. (1976) Phys. Rev. A, 13, 2287. Champion, E., Goursat, P., Besson, J. L., Madigou, V., Monthioux, M. and Lespade, P. (1992) Ceram. Eng. Sci. Proc., American Ceramic Society, USA, p. 732. Chaudhari, P., Graczyk, J. F. and Herd, R. (1972) Phys. Stat. Sol., 51, 801. Delverdier, O., Monthioux, M., Mocaer, D. and Pailler, R. (1992a) Proc. 5th Eur. Conf. Compos. Mater., (eds A. R. Bunsell, J. F. Jamet and A. Massiah), Bordeaux, France, p. 691. Delverdier, O., Monthioux, M., Oberlin, A., Lavedrine, A., Bahloul, D. and Goursat, P. (1992b) High Temp. Chem. Proc., 1, 139. Delverdier, O., Monthioux, M., Mocaer, D. and Pailler, R. (1993) J. Eur. Ceram. Soc., 12, 27. Delverdier, O., Monthioux, M., Mocaer, D. and Pailler, R. (1994) J. Eur. Ceram. Soc., 14, 313. Iijima, S. (1982) Solid State Chem., 42, 101. Inoue, Z. and Kourachi, Y. (1983) Proc. Int. Symp. Ceram. Components Engines, Japan, p. 519. Laanani, Fadéla (1997) Thèse de Doctorat, no 2707, Université Paul Sabatier, Toulouse, France. Le Coustumer, P., Monthioux, M. and Oberlin, A. (1993) J. Eur. Ceram. Soc., 11, 95. Madigou, Véronique, Monthioux, M. and Oberlin, A. (1990a) Proceedings of Matériaux Composites pour Applications à Hautes Températures (eds R. Naslain et al.), AMAC-CODEMAC, Bordeaux, France, p. 221. Madigou, Véronique, Monthioux, M. and Guimon, C. (1990b) Proc. XX Int. Conf. Carbon, GFEC, Paris, France, p. 220. Madigou, Véronique (1991) Thèse de Doctorat,n°116, Université de Pau et des Pays de l’Adour, France. Mocaer, D., Pailler, R., Naslain, R., Richard, C., Pillot, J.-P., Dunoguès, J., Delverdier, O. and Monthioux, M. (1993) J. Mater. Sci., 28, 2639. Monthioux, M. and Delverdier, O. (1996) J. Eur. Ceram. Soc., 16, 721. Monthioux, M., Oberlin, A. and Bouillon, E. (1990) Compos. Sci. Technol., 37, 21. Niihara, K., Suganuma, K., Nakahira, A. and Izaki, K. (1990) J. Mater. Sci. Lett., 9, 598. Oberlin, A. (1989) Chemistry and Physics of Carbon, vol. 22 (ed. P. A. Thrower), Marcel Dekker, New York, p. 1. Oberlin, A. and Rouchy, J.-P. (1971) Carbon, 9, 39. Oberlin, A., Oberlin, M. and Comte-Trotet, J. R. (1976) J. Microsc. Spectrosc. Electron., 1, 391. Oberlin, A., Boulmier, J. L. and Villey, M. (1980) Kerogen (ed. B. Durand), Technip, Paris, p. 191. Oberlin, A., Bonnamy, S., Bourrat, X., Monthioux, M. and Rouzaud, J.-N. (1986) Petroleum Derived Carbons (eds J. D. Bacha, J. W. Newman and J. L. White), ACS Symposium Series 303, USA, p. 85. Okamura, K., Sato, M., Hasegawa, Y. and Amano, T. (1984) Chem. Lett., 2059. Rossignol (1995) Thèse de Doctorat, Université de Limoges, France. Sasaki, G., Suganuma, K., Fujita, T., Hiraga, K. and Niihara, K. (1993) Mater. Res. Soc. Symp. Proc., 287, 335. Williams, D. B. and Carter, C. B. (1996) Transmission Electron Microscopy. A Textbook for Materials Science, Plenum Press, New York.

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Page 110 5 From short- to long-range order Structural organisation of silicon-based nanopowders

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Page 111 5.1 Chemical order studied by solid-state nuclear magnetic resonance André Pierre Legrand, Jean-Baptiste d’Espinose de la Caillerie and Youssef El Kortobi Native or heat-treated ceramics contain different atoms, the nuclei of which can be used, if provided with a magnetic moment, as NMR probes. Among them 1H,13C,27Al and 29Si are of particular interest. Different information can be obtained, depending on the method of observation and on the nature of the nucleus under analysis. Each proton has a magnetic moment (1H) and is relatively easy to observe. Nevertheless, it physically interacts through dipole–dipole interactions. This induces a line broadening of the spectrum, which can hide the chemical information concerning the different chemical species on, and inside ceramics (hydroxyls, water, etc.). So different approaches are possible by taking advantage of this dipolar interaction or on the contrary by suppressing it. Silicon nucleus (29Si), on the contrary to 1H, is an isotope with a low natural abundance of 4.7%. The same is true for 13C (1.1%). For as-formed ceramics, these diluted species are in dipolar interaction with the abundant proton (1H) species. To observe the nature of the different chemical species in the material, it is necessary to suppress this dipolar interaction. This is obtained through specific techniques. Aluminium nucleus (27Al) has a natural abundance of 100%. Contrarily to the previous nuclei, 27Al is not a Its quantum number is 5/2. This has positive and negative effects. A positive one is that the nucleus spin is sensitive to the local electric field gradient and consequently this gives an important information on the symmetry of the aluminium sites. The negative one is that if the electric field gradient from one site to another site is different, for example, tetrahedral to octahedral, the quantitative determination of the relative proportions could be questionable. 5.1.1 Nuclear magnetic methods As mentioned above, the as-formed amorphous ceramics may contain hydrogen atoms belonging to different chemical species such as hydroxyls and adsorbed water molecules. They constitute a set of 1H spins which, by dipolar interaction, produce a line broadening hiding the chemical information. This harmful effect disappears by heat treatment, with the expulsion of protonated species. 5.1.1.1 Rare spin half spectroscopy (29Si, 13C) Considering the low abundance of 29Si, 13C isotopes, the dipolar interaction responsible for line broadening is mainly due to neighbouring protons, for as-formed ceramics. To determine

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Page 112 the nature of the different functionalities of such nuclei in ceramics, it is necessary to use specific techniques. High power decoupling (HPD) To suppress the unwanted proton-rare spin dipolar interaction, the following method can be used. It requires application of a strong radio frequency field at the Larmor frequency of the abundant spins. Through this operation, the populations of spins in the low and high Zeeman energy levels are made equal and the resulting magnetisation falls to zero, leaving only the chemical shift interaction to be studied (Bloch 1958). Averaging the chemical shift anisotropy (CSA) – Magic angle spinning (MAS) The method involves putting the sample into a rotor, the axis of rotation of which is at an angle of rotation of 54.73° with the external magnetic field. This method is efficient not only to average the dipolar interaction, when MAS speed is high enough, but also the CSA (Engelhardt and Michel 1987). The chemical shift has a theoretical expression that is mathematically similar to that of the dipolar interaction. It is composed of an isotropic value, similar to that observed in liquid state, and an anisotropic part. In consequence, for a powder, this term produces a line broadening, due to the superposition of all possible orientations of the chemical species studied. This requires that the speed of rotation be high enough relatively to the CSA. If not, a more complex spectrum appears which is composed of the mere spectrum, flanked on either side with more or less intense replicas of this one. These satellites are separated by multiples of the rotation frequency. The different isotropic lines are characterised by a position in frequency with respect to a given reference, tetramethylsilane (TMS) in this book, and consequently can be identified. Numerous experiments in solids have demonstrated that the observed chemical shifts are near what is observed in liquids, if the environment is the same. Data already collected for species in solution can therefore be used to analyse solidstate spectra (Lippmaa et al. 1982; Lippmaa and Samoson 1982; Engelhardt and Michel 1987; Fyfe 1983; Marsmann 1981). As the technology has improved, allowing in particular to increase the MAS speed, it is possible to achieve such measurements using higher magnetic fields than previously, which enables increasing resolution and sensitivity. Cross-polarisation (CP) For as-formed ceramics containing an important amount of hydrogen, this method is jointly used with MAS to enhance the signal of carbon (C) or silicon (Si) bonded to protons. First, it involves cooling down the proton reservoir through a sequence shown in Figure 5.1.1. Second is the time to transfer the low temperature of the proton reservoir to that of rare spins. Last is to observe the signal of the rare spin while at the same time using the HPD. The method is consequently denominated CP/MAS. The complete sequence is represented in Figure 5.1.2. Model compound example Role and interest of the different methods mentioned above are summarised in Figure 5.1.3. 29Si NMR spectra of a powdered sample of pure cubic octamer silicic trimethylsilyl ester

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Figure 5.1.1 Setting of a spin temperature for the abundant spins (observation in a plane perpendicular to the director field B0): (a) the magnetisation is brought in the perpendicular plane after a 90° pulse; (b) the radio frequency field B1 is phase shifted by a 90° angle, bringing it parallel to the magnetisation; (c) after a certain time, the magnetisation decreases slowly in the B1 field and is measured after switching this field off, on the free induction decay. The upper part of the figure represents schematically the application of radio frequency pulses along the O x and O y axis in the rotating frame. (Reproduced with permission of John Wiley & Sons from Legrand et al. 1998.)

Figure 5.1.2 Complete sequence of CP for abundant spins (e.g. protons 1H) in interaction with less abundant spins (e.g. silicon 29Si). (Reproduced with permission of John Wiley & Sons from Legrand et al. 1998.)

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Figure 5.1.3 Rare spin 29Si solid state NMR spectra at 59.6 MHz of a powdered pure cubic octamer silicic trimethylsilyl ester (Q8M8) as model compound, showing the influence of different types of NMR measurements: (a) static or broadband spectrum with no high power decoupling of the protons. Only a broad line is observed: no chemical information is available; (b) static or broadband spectrum with high power decoupling of protons. Two bands, due to chemical anisotropy, corresponding to two Si are observed: no clear chemical identification available; (c) HPD with CP (contact time 1 ms) and MAS with a speed or rotation not high enough (567 Hz). Chemical information partially available; (d) CP/MAS + HPD. Chemical information available. (Courtesy of P. Tougne.) (Q8M8) present the following aspects: 1 With static or broadband method, only broad lines mainly due to the dipolar interaction 1H—29Si are observed. No chemical information is available. 2 As above, with HPD of the protons, CSA of two Si types are observed. 3 CP/MAS method, but with MAS speed of 560 Hz. Side bands with the mere spectrum are observable. 4 CP/MAS method, with MAS speed high enough. Side bands are suppressed and the spectrum shows a triplet and a quadruplet.

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Page 115 5.1.1.2 Aluminium 27Al NMR spectroscopy Unlike 29Si or 13C, 27Al has an I = 5/2 spin and therefore a nuclear quadrupole moment. This implies that not only this nucleus is sensitive to the chemical shift and dipolar interactions but also to the quadrupolar one. The six energy levels (2I + 1) of the 27Al due to the Zeeman effect are modified, assuming a first-order perturbation effect, which results in five different transition frequencies. The corresponding Hamiltonian is:

(5.1) where eQ is the nuclear quadrupolar moment and Vzz = eq, the main electric field gradient (|Vzz| > |Vyy| > | Vxx|); η = (Vyy − Vxx)/Vzz is the asymmetrical parameter. Oxyz are principal axes of the tensor that describe the electric field gradient. All transitions except the central one are subject to a firstorder quadrupolar interaction depending on the orientation of the electric field gradient tensor with respect to the external magnetic field. Since in powder samples there is a random distribution of such orientations, it results in a large spreading of the line which prevents its observation. On the contrary, the central line is only sensitive to the second-order perturbation effect, which causes a line spreading proportional to (e2Qq/h) 2/νL = νQ, where νL is the Larmor frequency of the 27Al nucleus. Contrarily to spin half, the observation is not obtained with a 90° pulse, although this one is obtained with a solution of aluminium salt, in which the Brownian motion averages the quadrupolar effect. Samoson and Lippmaa (1983) proved that the signal intensities are independent of the magnitude of νQ only if the pulse length is short enough. For an error less than 5% in the signal intensity for the centre band of the central transition the spin flip angle ω1τ must satisfy the condition ω1τ(I + 1/2) ≤ π/6, with I = 5/2 for 27Al. |Vzz| > |Vyy| > |Vxx|); η = (Vyy − Vxx)/Vzz is the asymmetrical parameter. Oxyz are principal axes of the tensor that describe the electric field gradient. All transitions except the central one are subject to a first-order quadrupolar interaction depending on the orientation of the electric field gradient tensor with respect to the external magnetic field. Since in powder samples there is a random distribution of such orientations, it results in a large spreading of the line which prevents its observation. On the contrary, the central line is only sensitive to the second-order perturbation effect, which causes a line spreading proportional to (e2Qq/h)2/νL = νQ, where νL is the Larmor frequency of the 27Al nucleus. Contrarily to spin half, the observation is not obtained with a 90° pulse, although this one is obtained with a solution of aluminium salt, in which the Brownian motion averages the quadrupolar effect. Samoson and Lippmaa (1983) proved that the signal intensities are independent of the magnitude of νQ only if the pulse length is short enough. For an error less than 5% in the signal intensity for the centre band of the central transition the spin flip angle ω1τ must satisfy the condition ω1τ(I + 1/2) ≤ π/6, with I = 5/2 for 27Al. An example of the influence of the pulse duration on the relative intensities of the lines is given in Figure 5.1.4 for 23Na (spin 3/2). Such quadrupolar and line shifts render more difficult the interpretation of the line structure but are able to provide additional information of the charge distribution (Figure 5.1.5). 5.1.1.3 Interest of the method The chemical shift of the nuclear resonance provides an incredibly delicate fingerprint of the chemical functionality of an element in a compound. NMR has provided chemists with one of the most powerful qualitative and quantitative analytical tools available to the chemist, material scientist and medical technician. The advent of high-resolution techniques appropriate to nuclei in solids supplied further dimensions to do qualitative analysis of chemical functionalities. Moreover, some information concerning structure and local organisation of clusters can be obtained. Such a method had been largely used for ceramic characterisation (Müller 1999; Templin et al. 1999). 5.1.2 SiC powders analysis In Figure 5.1.6, the 29Si spectrum for sample as-formed SiC163 is shown. Like previously published spectra of silicon carbide powders (Table 5.1.1), it consists of a superposition of two kinds of line: a single line at −16.0 ppm can be attributed to the β-SiC type, whereas the triplet consisting of lines at −15.6, −20.5 and −24.9 ppm can be attributed to α-SiC polytypes.

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Figure 5.1.4 2D solid state 23Na MAS–NMR spectra at 79.39 MHz of sodium chlorine and nitrite salts. (a) Influence of the pulse length (from 0.25 to 5 μs, by step of 0.25 μs) on the relative intensities of the two sodium sites; (b) Relative proportion has to be evaluated at short pulse length (here 1:6). (Courtesy of P. Tougne.) As the experimental high-resolution solid-state NMR technique involves only MAS, the spectra can be decomposed and the amount of each type, α-SiC or β-SiC, is proportional to the respective areas of the


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different lines, if the repetition time for accumulation is long enough relatively to the spin-lattice relaxation times. The 29Si spectrum of SiC174 is very similar to the previous spectrum except that the amount of α-SiC has decreased a little.

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Figure 5.1.5 Solid state 27Al MAS–NMR spectrum at 130.32 MHz of gibbsite and saponite. MAS speed: 12 kHz. Clear distinction between four- and sixfold coordinated aluminium in AlO4 and AlO6 units is observed. 27Al chemical shifts are between +50 and +80 ppm for AlO4 and about −10 to +20 ppm for AlO6.

Figure 5.1.6 Solid state 29Si MAS–NMR spectra at 59.6 MHz of as-formed silicon carbide (duty cycle, 500 s; 90° pulse, 5 μs; number of scans, 161, 100, 164 and 60). Spectra are classified as function of decreasing residence time into the laser irradiation cell (Table 2.1). (Reprinted from Tougne et al. (1993) Diamond Relat. Mater., Evolution of the structure of ultrafine SiC-laser-formed powders with synthesis conditions, 2, 486– 490, with permission from Elsevier Science.)

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Page 118 Table 5.1.1 29Si chemical shift in ppm/TMS for silicon carbides Authors Polytype 3C (C) 6H (A B C) 15R (A B C) 4H (B C) 2H (D) Finlay et al. 1985 −18.3 −13.9 −20.2 −24.2 Guth et al. 1987 −18.4 −16.1 −14.3 −20.4 −24.9 −14.6 −20.5 −24.1 Hartman et al. 1987 −18.3 −13.9 −2.02 −24.5 −14.9 −20.8 −24.4 Wagner et al. 1989 −18.4 −16.1a Dando and Tadayyoui 1990 −16 −14.4 −20.5 −25 −15.2 −20.2 −24 Apperley et al. 1991 −18.4 −14.7 −20.9 −25.4 −19.7 −22.5 −20 Schmidt et al. 1991–2 −18 −14 −20 −25 Richarson et al. 1992 −14.5 −20.4 −25.1 Dybowski et al. 1996 −16 −22 −27 Tateyama et al. 1997 −18.4 −14.7 −19.7 −25.4 −20.9 −22.5 −20 Note a Depends on the crystalline structure of the sample (−16.1 for small single crystals). No quantitative accurate decomposition has been performed but it is qualitatively apparent from the spectra. Moreover, on comparing the figure which shows the spectrum for sample as-formed SiC173 with the figure which shows the spectrum for sample SiC177 the trend follows the same line. Increasing the residence time into the hot region of the cell where the synthesis takes place results in a decreasing amount of α-SiC phase. To assess this evolution, samples as-formed SiC163 and SiC177 were annealed in a furnace at 1600°C for 24 h. Figure 5.1.7 shows the spectra of the annealed samples SiC163 and SiC177. One single line at −16.0 ppm now appears which can unambiguously be attributed to a pure β-SiC phase, it appears therefore rather surprisingly that the β-SiC phase is the high temperature stable phase. By comparing the evolution of samples SiC163 and SiC171, one can study the effect of heat treatment. Table 5.1.2 shows that the grain growth occurs at a higher temperature for sample SiC171 than for sample SiC163. At 2000°C, the densification of sample SiC171 is 94% but the grain size has become comparable to the grain size of a commercial powder. The problem is then to achieve good densification and conserving a fine structure. A good knowledge of the local structure evolution would be very helpful. Figure 5.1.8 presents the NMR spectra of samples SiC163 and SiC171 after heat treatment between 1600°C and 2050°C. One can see that the α-phase is always present in sample SiC171, while it almost disappears in sample SiC163 between 1700°C and 1900°C. At 2050°C, β-phase is transformed into α polytypes. 6H, 15R and 4H polytypes are unambiguously identified in sample SiC171 while only the 6 H environment is in sample SiC163. It seems then that there is a relation between the annealing and the presence of the α-SiC polytypes.

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Figure 5.1.7 Solid state 29Si MAS–NMR spectra at 59.6 MHz of 1600°C (24 h) annealed silicon carbide (duty cycle, 500 s; 90 ° pulse, 5 μs; number of scans, 69 and 57). (Reprinted from Tougne et al. (1993) Diamond Relat. Mater., Evolution of the structure of ultrafine SiC-laser-formed powders with synthesis conditions, 2, 486–490, with permission from Elsevier Science.) Table 5.1.2 Surface of different silicon carbide batches as a function of temperature dwell (dwell duration: 1 h) Temperature (°C) SBET (m2/g) SiC171 SiC163 SiC177 1600 24.9 13.2 8.1 1800 3.7 7.3 3.3 2000 0.09 5.3 2.3 5.1.3 Si/C/N powders analysis 5.1.3.1 29Si and 13C solid-state MAS–NMR measurements Well-crystallised silicon nitrides have been analysed and 29Si chemical shifts of the two main α and β structures had been determined (Table 5.1.3). Average chemical shift identification of different Si environments in Si/C/N compounds has been done by comparison between pure silanes and pure silicon carbide or silicon nitride. Analysis of the results is given in Table 5.1.4. Besides, different authors have estimated the presence in less organised compounds of Si environments as SiCN3, SiC2N2 and SiC3N (Seitz et al. 1996; Gérardin et al. 1997; Verdecia

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Figure 5.1.8 Evolution of 29Si MAS–NMR spectra of sample 163 and 171 as a function of the temperature (dwell duration 1 h). Top of β-SiC peak had been cut to observe polytypes signal. Table 5.1.3 Chemical shift of 29Si in ppm/TMS for crystallised amorphous silicon nitrides. Values within parentheses are full-width at half-height of each peak Authors β-Si3N4 (singlet) α-Si3N4 (doublet) equal intensities Amorphous Carduner et al. 1987, 1990 −48.7 (2.5) −46.8 (1.8) −48.9 (1.8) −46.4 (27) Harris et al. 1990 −48.5 (1.3) −47.1 −49 Table 5.1.4 29Si average chemical shift of different silicon nitrides. Chemical shift measurements done on silane series (upper part) exhibit an average shift of −20 ppm comparatively to pure silicon carbide and silicon nitride. Consequently interpolation is done for intermediary environment of the Si (lower part) Silane Chemical shift (ppm/TMS) (CH3)4Si 0 (CH3)3SiN(CH3)2 6.9 to 5.0 (CH3)2Si(N(CH3)2)2 −1.3 to 1.85 (CH3)Si(N(CH3)2)3 −16.3to−17.5 Si(N(CH3)2)4 −28.1to−28.6 Silicon environment in solid state Average chemical shift (ppm/TMS) SiC4 −20 SiC3N −13 to −14 SiC2N2 −21 to −22 SiCN3 −36 to −37 SiN4 −46 to −49

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Figure 5.1.9 Chemical shift of 29Si in ppm/TMS for crystallised SiCxN4−x compounds as a function of the HTT of different ceramics from Seitz et al. (1996), Gérardin et al. (1997) and Verdecia et al. (1998). Symbols are ■ SiN4, ● SiCN3 and ◆ SiC2N2; highest temperature is for well-crystallised silicon nitrides (Carduner et al. 1987, 1990). et al. 1998). Chemical shifts of such Si environments are dependent on the first and second sphere of coordinence. This one is able to contain hydrogen atoms for as-formed or low heat-treated materials and consequently the chemical shift is susceptible to modifications. Figure 5.1.9 summarises different observations. Series HMDS41 to HMDS44, heat-treated from 1000°C to 1600°C in N atmosphere during 1 h, have been analysed using 29Si MAS–NMR. This method enables us to distinguish, through previous mentioned observations given in Figure 5.1.9, the evolution of the local structure around Si. Figures 5.1.10 and 5.1.11 give the general results of the survey. They are classified as a function of the initial C/Nini composition or of the heat-treatment temperature (HTT). Some remarks have to be made: • At HTT 1000°C, no important differences are observable. A broad massif centred around −50 ppm demonstrates an amorphous silicon nitride environment. Besides, a shoulder near −35 ppm can be attributed to SiCN3, the importance of which decreases with initial C/N ratio. A small peak at −110 ppm corresponds to silica.

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Figure 5.1.10 Solid state 29Si MAS–NMR spectra at 99.36 MHz, rotation frequency of 5 kHz, of HMDS file:///C|/Users/steppenwolf/Documents/educatio...ks/to%20burn/nanocomposites/files/page_122.html (1 of 2) [20/01/2010 10:04:22 AM]

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powders heat-treated during 1 h in N atmosphere, as a function of the initial C/N (atomic number) composition, at increasing temperature. (a) HMDS41 (duty cycle, 500 s; 90° pulse, 5 μs; number of scans, 160, 160, 160, 12, 84); (C/N)ini = 0.95; (b) HMDS42 (duty cycle, 500 s; 90° pulse, 5 μs; number of scans, 160, 103, 160, 75); (C/N)ini = 0.73; (c) HMDS43 (duty cycle, 500 s; 90° pulse, 5 μs; number of scans, 144, 125, 140, 120, 81); (C/N)ini = 0.67; (d) HMDS44 (duty cycle, 500 s; 90° pulse, 5 μs; number of scans, 150, 160, 102, 40); (C/N)ini = 0.57.

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Figure 5.1.11 Solid state 29Si MAS–NMR spectra of HMDS powders. Experimental conditions are given in file:///C|/Users/steppenwolf/Documents/educatio...ks/to%20burn/nanocomposites/files/page_123.html (1 of 2) [20/01/2010 10:04:23 AM]

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Figure 5.1.10 caption. (a) 1000°C/N2/1 h; (b) 1400°C/N2/1h; (c) 1500°C/N2/1 h; (d) 1600°C/N2/1h.

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Page 124 • At HTT 1400°C, a silicon nitride broad peak (SiN4 and SiCN3) emerges from the massif. It appears as a shoulder near −20 ppm (SiC2N2). X-ray diffraction patterns do not provide similar information, owing to the amorphous organisation of the material. • At HTT 1500°C, the structure of the massif is more evident. Three more or less well-resolved peaks are observed which are respectively attributed to SiN4, SiCN3 and SiC2N2. A shoulder near −18 ppm seems to appear which can be silicon carbide. • At HTT 1600°C, crystallised silicon nitride (α and β) are observed, the proportion of which increases as initial C/N decreases. A silicon carbide massif (SiC4, SiC3N) is clearly separated from this one. Two significant samples, HMDS42 and HMDS44, are analysed in more detail using least square decomposition method, assuming different occupations sites as given in Figure 5.1.9. Results are shown on Figures 5.1.12 and 5.1.13. Table 5.1.5 gives the relative proportions of the different sites. Neither SiC4 nor SiC2N2 are observed till HTT is 1000°C. A mixture of such sites is observed at a higher temperature. From HTT 1550°C, crystallised SiN4 (α and β) sites are clearly identified. In the silicon carbide broad line, SiC4 and SiC3N sites are present. The relative proportions of crystallised silicon nitride clearly depend on the initial C/Nini composition (C/Nini = 0.73 and 0.57, respectively, for HMDS42 and HMDS44). Such different evolutions are shown on the chemical composition (Figure 5.1.14). An important remark concerns the duty cycle, that is, the time between two accumulations, to obtain reliable spectra. This time has to be of the order of five times that of the spin-lattice

Figure 5.1.12 Solid state 29Si MAS–NMR spectra at 99.36 MHz, rotation frequency of 5 kHz, of (a) HMDS42 and (b) HMDS44 heat-treated at 1000°C, 1400°C and 1500°C during 1 h in N atmosphere and decomposition, using a least square method, into different assumed Si environments (El Kortobi 1998).

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Figure 5.1.13 Solid state 29Si MAS–NMR spectra at 99.36 MHz, rotation frequency of 5 kHz, of (a) HMDS42 and (b) HMDS44 heat-treated at 1550°C, 1600°C and 1700°C during 1 h in N atmosphere and decomposition, using a least square method, into different assumed Si environments (El Kortobi 1998). relaxation times of the species to observe. Figure 5.1.15 shows spectra obtained for different values of such time. One notices that the relative apparent proportion of silicon nitride is only significant for temperatures higher than 1550°C, which shows that the two massifs are associated with relatively short and long spinlattice relaxation times. Different authors, using MAS–NMR, had similarly used duty cycle as long as 30 mn to obtain reliable spectra (Dybowski et al. 1996). In the case considered here, it is necessary to reach 8000 s to estimate the relative proportion of silicon nitride in the different specimens. Amorphous silicon carbide massif relaxes more rapidly as shown in Figure 5.1.16. In Chapter 5.2, electron paramagnetic resonance (EPR) measurements show that the materials contain an important amount of paramagnetic centres. Besides, the study of the magnetisation as a function of the temperature as in the example on HMDS44 heat-treated at 1550°C, shows a Curie law dependence. This is in favour of localised paramagnetic centre organisation. As the above observation concerning the difference of spin-lattice relaxation times between silicon nitride and other species, it is obvious to consider that such centres are inside the amorphous silicon carbide structure. Same samples, HMDS42 and HMDS44, have been analysed using 13C MAS–NMR (Figure 5.1.17). The useful CP/MAS–NMR is useless in that temperature range, due to the lack of proton in the material. The difficulty in acquiring a good signal over noise spectra is related to the low-frequency measurement. This had been done to prevent side bands coming from a large CSA of the aromatic carbons. Different comments have to be made: • At HTT 1000°C, some differences are observable. A broad massif centred around −130 ppm is attributed to aromatic sp2 and oxygenated carbons. This is generally

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Page 126 Table 5.1.5 29Si chemical shift δ in ppm/TMS, obtained using a least square .tting method for HMDS42 and HMDS44 samples. Gaussian fit shape assumed for a theoretical decomposition of the experimental spectra. Δ is the Gaussian quadratic square deviation of each specie assumed HTT(° SiN4 SiCN3 SiC2N2 SiC3N SiC4 C) δ Δ %δ Δ %δ Δ % δ Δ % δ ö % (ppm) (ppm) (ppm) (ppm) (ppm) (ppm) (ppm) (ppm) (ppm) (ppm) HMDS42 1000 −53.6 18 67 −35 20 33 / / / / / / / / / 1400 −50 14 54 −34.6 13 30 −21 14 16 / / / / / / 1500 −48 14 34 −32 15 31 / / / / / / −14.5 15 35 1550 / / 9 / / / −21 8 6 −9 10 34.5 −17 8 31.5 1600 / / 13 / / / / / / −6 13 31.1 −18 10 56.5 1700 / / 10 / / / / / / −11 5 12.6 −17 6 74.4 HMDS44 1000 −53.6 17 80 −37 17 20 / / / / / / / / / 1400 −50 14 67 −36 14 16 −23 14 17 / / / / / / 1500 −53 12 40 −37 14 36 / / / / / / −20 13 24 1550 / / 48 / / / −22 7 5.5 −7 9 12.3 −17 10 34.2 1600 / / 40 / / / −23 5 6 −10 8.5 11.4 −17 8 42.6 1700 / / 42 / / / 21 6 3.2 −9 8 14.1 −17 6 40.8

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Figure 5.1.14 Atomic chemical composition SiyCxNz (x + y + z = 1) of HMDS42 and HMDS44 as a function of heat treatment temperature. Arrows indicate evolution from as-formed to highest treatment temperature. considered as ‘free carbon’. Besides, for HMDS42, a small peak near −20 ppm can be attributed to sp3 Si bonded C. • From HTT 1400°C, the amount of Si bonded C increases, which is in accordance with the 29Si MAS–NMR observation. Besides, the existence of other environments such as CSixNy are possible. • At HTT 1550°C, the line around −20 ppm, related to silicon carbide, is more important for both samples. Although the difference is noticeable concerning the ‘free carbon’ still present in the richer C HMDS42 sample than in HMDS44, were it is now absent. More generally, similar determinations are of interest to characterise silicon carbide fibres and consequently their mechanical properties (Hommel et al. 1990; Laffon et al. 1989). Some comparisons can be done between the atomic chemical compositions and 29Si NMR determinations. From Table 5.1.5, with the relative proportions of SiN4(p), SiCN3(q), SiC2N2(r), SiC3N(s) and SiC4(t), it is possible to calculate the average atomic compositions SiyCxNz (x + y + z = 1) as a function of HTT and to plot Figure 5.1.18, using the following formula:

(5.2) The above values are listed in Table 5.1.6. Such obtained plots show a convergence of the two types of analysis at high heat temperature treatment of both samples. The divergence observed at low HTT originates from

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Figure 5.1.15 Solid state 29Si MAS–NMR spectra obtained at different duty cycle (D1) values on HMDS44 and HMDS42, at different heat treatment temperatures. Spectra are adjusted so that the maximum of each spectrum is similar to the others. Signal over noise depends on the number of scans and consequently of the total time of accumulation, which is different for some spectra (El Kortobi 1998). an underestimation of C content. Those not being bonded to silicon C atoms are set in the amorphous ‘free carbon’ as it had been discussed above (Figure 5.1.17). This effect is not important for sample heat-treated at temperature higher than 1550°C, as it can be observed on the points corresponding to NMR measurements. They are in the vicinity of those in the chemical analysis. From the C content values obtained for chemical analysis and NMR, it is possible to estimate the proportion of ‘free carbon’ in total C. This calculation is only valid, within the experimental errors, in the temperature range 1000–1550°C (Figure 5.1.19). 5.1.3.2 Thermal evolution from NMR measurements This evolution can be roughly decomposed in three steps as follows: First step: HTT inferior to 1000°C This domain, from as-prepared powder till almost end of mineralisation, is mainly discussed in Chapter 2. The most efficient methods of

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Figure 5.1.16 Solid state 29Si MAS–NMR spectra obtained at different duty cycle (D1) values on HMDS44 heat-treated at 1600°C (LB = 30 Hz, NS = 16) (El Kortobi 1998). characterisation are in particular TGA and IR. It is characterised by emission of H2O, NH3 and CH4. It results in the formation of Si N with ammoniac (Gérardin et al. 1997; Bowen 1980) and Si—C bonds with methane (Gérardin et al. 1997, Hasegawa and Okumura 1983) emissions. Second step: 1000–1500°C As seen from 29Si MAS–NMR, it starts with a rich nitrogen (N) to Si environment (SiN4, SiCN3) depending on the initial C/Nini concentration. Besides, 13C MAS–NMR demonstrate the presence of C not bonded to Si, named ‘free carbon’. The proportion of such C decreases considerably till 1500°C. Parallelly, it is observed as an increase of C to Si environment. Third step: 1550–1700°C The crystallisation of the material is observed, with important loss in weight due to the departure of N, consecutive to a carbo-reduction. Demixtion occurs with the appearance of wellcrystallised α − β SiN4 and of amorphous silicon carbides (SiC4, SiC3N and SiC2N2). 5.1.4 Pre-alloyed Si/C/N/Al/Y nanopowders In Chapter 7 explanations are given on the interest to use additives for the sintering of silicon nitride based materials. The objective here is to characterise nanosized ready-to-sinter powders including additives. Two types of additives have been incorporated: alumina (Al2O3) and yttria (Y2O3).

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Figure 5.1.17 Solid state 13C MAS–NMR spectra at 25.15 MHz, rotation frequency of 5 kHz, of HMDS42 and HMDS44 at different heat treatment temperatures.


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Figure 5.1.18 Atomic NMR and chemical compositions SiyCxNz (x + y + z = 1) of HMDS42 and HMDS44 as a function of heat-treated temperature. Arrows indicate evolution from as-formed to highest treatment temperature.

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Page 131 Table 5.1.6 Atomic NMR composition CxSiyNz (x + y + z = 1) of HMDS42 and HMDS44 as a function of heat treatment temperature HTT(°C) C Si N HMDS42 1000 0.03896 0.4329 0.52814 1400 0.06987 0.43668 0.49345 1500 0.15315 0.45045 0.3964 1550 0.33175 0.47393 0.19431 1600 0.38462 0.48077 0.13462 1700 0.42157 0.4902 0.08824 HMDS44 1000 0.02155 0.43103 0.54741 1400 0.05652 0.43478 0.5087 1500 0.12796 0.47393 0.3981 1550 0.20354 0.44248 0.35398 1600 0.24658 0.45662 0.2968 1700 0.23767 0.44843 0.3139

Figure 5.1.19 ‘Free carbon’ evolution as a function of heat treatment temperature. Values are obtained owing to the formula: (xCA − xNMR)/xCA where xCA, xNMR are those of the chemical analysis and NMR, respectively. 5.1.4.1 Free carbon effect Samples HSAl03 and HSAl05, heat-treated from 1000°C to 1600°C in N atmosphere during 1 h, have been analysed using 29Si MAS–NMR (Figure 5.1.20). The following observations can be made: • From as-formed to HTT 1000°C, no important differences are observed. A broad peak around −20 ppm shows different environments of Si (SiC3N, SiC2N2 and/or SiC4). • At HTT 1400°C, a line narrowing is observed. Only SiC2N2 and/or SiC4 environments remain.

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Figure 5.1.20 Solid state 29Si MAS–NMR spectra at 99.36 MHz, MAS speed of 5 kHz, of (a) HSAl03 and (b) HSAl05 at different heat treatment temperatures. At 1600°C, no silicon nitride remains; compare to Table 3.1.3 (El Kortobi 1998). • At HTT 1600°C, the narrowing is more important. A structure is perceptible, which can be attributed to α and β silicon carbide. 13C MAS–NMR is able to improve such results. Only as-formed samples can be analysed using CP and MAS. The sites are CH3Si and Si—CH2—Si in the range 0–5 ppm. • At HTT 1000°C, as hydrogen is expulsed with heat treatment, 13C MAS–NMR sp3 and sp2 C are observed. • From HTT 1400°C to 1600°C, the ‘free carbon’ massif disappears progressively. The line around 24 ppm is proportionally improved. This confirms the presence of partially crystallised silicon carbide (Figure 5.1.21). 27Al MAS–NMR shows three aluminium environments for as-formed materials. Tetra, penta and octahedral environments are observed (Figure 5.1.22). • At HTT 1000°C, such three sites are still observable, with a significant increase of the octahedral environment. • From HTT 1400°C to 1600°C, besides tetra and octahedral sites, there is a broad massif at lower field, around 90 ppm. This can be attributed, owing to measurements done on defined samples (AlON and SiAlON) (Figure 5.1.23), to aluminium oxynitrides AlOxN4−x, with 0 < x < 4 (Smith 1992; Mackenzie et al. 1994; Fitzgerald et al. 1994). At HTT 1400°C, this massif contains AlO3N and AlO2N2 (96 ppm), AlON3 (106 ppm) and AlN4 (114 ppm). At HTT 1600°C, the peak corresponding to AlN4 is easier to observe (Figure 5.1.24).

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Figure 5.1.21 Solid state 13C MAS–NMR spectra at 25.15 MHz, rotation frequency of 5 kHz, of HSAl03 at different heat treatment temperatures (El Kortobi 1998).


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Figure 5.1.22 Solid state 27Al MAS–NMR spectrum at 130.32 MHz of (a) HSAl03 and (b) HSAl05. MAS speed 12 kHz, 1 μs pulse, duty cycle 1 s, number of scans 600 (El Kortobi 1998).

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Figure 5.1.23 Solid state 27Al MAS–NMR spectrum at 130.32 MHz of different reference samples. MAS speed 12 kHz, 1 μs pulse, duty cycle 1s, number of scans 600. At 1600°C, N is bonded to aluminium; compare to Table 3.1.3. So, the following remarks can be made: • Al—O and A—N bonds are observed. • N, initially bonded to Si as SiC3N and SiC2N2, is finally bonded to aluminium oxynitrides AlOxN4−x. In conclusion, the study of the local environment of 13C,27 Al and 29Si nuclei, enables us to understand why, within the proportions of reactants used for the synthesis, such powders did not produce a compound suitable for sintering (Chapter 7). The large amount of ‘free carbon’ favours the nitridation of aluminium owing to the reaction: This aluminium nitridation is responsible for the N loss in the Si environment. A new series of preparation is done with hydrogen and ammonia added on. This is to obtain the trapping of reactive free radical •CH3. The 27Al spectra show that the aluminium nitridation is noticeably decreased (Figure 5.1.24). It presents an intense line at 60 ppm, attributed to AlO4, sharper than that observed in a glass or conventional sintered sample. This line corresponds to a well-organised tetrahedral specie, only surrounded by oxygen atoms. This should be an aluminium solid solution in β-SiAlON. Similarly, the 29Si spectra demonstrate the presence of silicon nitride (Figure 5.1.25). The comparison between samples prepared without (HSAl07) and with (HSAl09) ammonia (Figure 5.1.26) shows clearly this

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Figure 5.1.24 Solid state 27Al MAS–NMR spectrum at 130.32 MHz of HSAl12 to HSAlY17. MAS speed 12 kHz, 1 μs pulse, duty cycle 1 s, number of scans 600. phenomenon not only for Si but also for aluminium. This support the analysis done above on the role of ‘free carbon’ effect which is clearly observable on 13C spectra (Figure 5.1.26). 5.1.4.2 Influence of hot pressing A comparison can be done between hot pressed and heat-treated nanopowders. As examples HSAl09 29Si and 27Al NMR measurements are shown in Figures 5.1.27 and 5.1.28. As the hot pressing had been done at a higher temperature than that of the nanopowder, silicon carbide has almost disappeared. The structured peak, corresponds to different silicon nitrides. As mentioned in Table 5.1.3, α- and β-Si3N4 respectively present a doublet and a singlet lines. Carduner et al. (1987) measured linewidth between 1.5 and 2.5 ppm (full-width at half-height) for well-crystallised material. Although the spectra are in accordance with these measurements, the decomposition, in order to have the relative proportion of β/ (α + β), is questionable. The powder heat treated at 1600°C is able to contain some amorphous part, which explains the higher intensity at −47 ppm. For the sintered powder, the β proportion estimation by X-ray diffraction is 30% (Table 7.6), which is qualitatively in accordance with the 29Si NMR observation. 27Al NMR spectra (Figure 5.1.27) show the AlxOyNz broad line with the typical peaks or shoulder characteristics of AlN4, AlO4 and AlO6 environments. A significant difference is observed with the relative proportions of the different species from one sample to another. This has to be compared with Figure 5.1.22 where alumina is bonded to variable environments containing N.

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Figure 5.1.25 Solid state 29Si and 13C MAS–NMR spectra at 99.36 MHz and 125.76 MHz of HSAl14, HSAl15 heat treated at 1600°C/N2/1h.29Si MAS speed 5 kHz, 5.2 μs pulse, duty cycle 500 s, number of scans 38, 4; 13C MAS speed 5 kHz, 5 μs pulse, duty cycle 500 s, number of scans 244.


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Figure 5.1.26 Solid state 29Si and 27Al MAS–NMR spectra at 99.36 MHz and 130.32 MHz of HSAl07, HSAl09 heat treated at 1600°C/N2/1 h. 29Si MAS speed 5 kHz, 5.2 μs pulse, duty cycle 500, 8000 s, number of scans 38, 121 and 4; 27Al MAS speed 14 kHz. 1.2 and 1 μs pulse, duty cycle 1 s, number of scans 3600. An underestimation of the amount silicon nitride is observed when a too short duty cycle is used (see Figure 5.15).

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Figure 5.1.27 Solid state 29Si MAS–NMR spectra at 99.36 MHz. MAS speed 5 kHz, 5.5 μs, duty cycle 8000 s, number of scans 10. Comparison between sintered and heat treated HSAl09 nanopowder.

Figure 5.1.28 Solid state 27Al MAS–NMR spectrum at 130.32 MHz. MAS speed 14 kHz, 1.2 μs pulse, duty cycle 1 s, number of scans 600. Comparison between sintered and heat treated HSAl09 nanopowder. 5.1.5 Conclusion The analyses of different nanopowders, as-prepared and heat treated at different temperature, show the efficiency of the solid-state NMR, particularly for not well-crystallised samples. The possibility of obtaining cross informations from different nuclei, as 13C,27Al and 29Si and eventually proton via CP will improve the structural analysis. This enables us to modify synthesis conditions as demonstrated on pre-alloyed Si/C/N/Al/ Y nanopowders.

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Page 138 References Apperley, D. C. Marshall, G. L. and Thompson, D. P. (1991) J. Am. Ceram. Soc., 74(4), 777. Bloch, F. (1958) Phys. Rev. 111, 841. Bowen, H. K. (1980) Mater. Sci., 44,1. Carduner, K. R., Blackwell, C. S., Hammond, W. B., Reidinger, F. and Hatfield, G. R. (1990) J. Am. Chem. Soc., 112, 4676. Carduner, K. R., Carter III, R. O., Milberg, M. E. and Crosbie, G. M. (1987) Anal. Chem., 59, 2794. Dando, N. R. and Tadayyoni, M. A. (1990) J. Am. Ceram. Soc., 73(8), 2242. Dybowsky, C., Gaffney, E. J., Sayir, A. and Rabinowitz, M. (1996) Colloids. Surf., A118, 171. El Kortobi, Youssef (1998) Thesis, Université P. M. Curie, Paris, France, http://corail.sudoc.abes.fr80 & SudocCatalogue no 053764552. Engelhardt, G. and Michel, D. (1987) High-resolution Solid-state NMR of Silicates and Zeolites, J. Wiley & Sons, New York. Finlay, G. R., Harman, J. S., Ricardson, M. F. and Williams, B. L. (1985) J. Chem. Soc., Chem. Commun., 159. Fitzgerald, J. J., Kohl, S. D. and Piedra, G. (1994) Chem. Mater., 6, 1915. Fyfe, C. A. (1983) Solid-state NMR for Chemists, C.F.C. Press Guelph. Gérardin, C., Taulelle, F. and Bahloul, D. (1997) J. Mater. Chem., 7(1), 117. Guth, J. R. and Petuskey, W. T. (1987) J. Phys. Chem., 91(20), 5361. Harris, R. K., Leach, M. J. and Thompson, D. P. (1990) Chem. Mater., 2, 320. Hartman, J. S., Richardson, M. F., Sherriff, B. and Winsborrow, B. G. (1987) J. Am. Soc., 109, 6059. Hasegawa, Y. and Okamura, K. (1983) J. Mater. Sci., 18, 3633. Hockey, J. A. (1965) Chem. Ind. (London) 57. Hommel, H., Miquel, J. L. and Legrand, A. P. (1990) L’Industrie Céramique, 849,5. Laffon, C., Flank, A. M., Lagarde, P., Laridjani, M., Hagege, R., Olry, P., Cotteret, J., Dixmier, J., Miquel, J. L., Hommel, H. and Legrand, A. P. (1989) J. Mater. Sci., 24, 1503. Lippmaa, E. and Samoson, A. (1982) Bruker Rep., 1,6. Lippmaa, E., Magi, M., Samoson, A., Engelhardt, G. and Grimmer, A. R. (1982) J. Am. Chem. Soc., 102, 4899. Mackenzie, K. J. D., Meinhold, R. H., White, G. V., Sheppard, C. M. and Sherriff, B. L. (1994) J. Mater. Sci., 29, 2611. Marsman H. (1981) NMR Basic Principles and Progress, 17 (eds P. Diehl, E. Fluck and R. Kosfeld), SpringerVerlag Berlin, p. 65. Müller, K. (1999) Precursor-derived Ceramics (eds J. Bill, F. Wakai and F. Aldinger), Wiley-VCH. Richardson, M. F., Hartman, J. S., Guo, D. and Winsborrow, B. G. (1992) Chem. Mater., 4, 318. Samoson, A. and Lippmaa, E. (1983) Phys. Rev., B28, 6567. Schmidt, W. R., Interrante, L. V., Doremus, R. H., Trout, T. K., Marchetti, P. S. and Maciel, G. E. (1991) Chem. Mater., 3, 257. Seitz, J., Bill, J., Egger, N. and Aldinger, F. (1996) J. Eur. Ceram. Soc., 16, 885. Smith, M. E. (1992) J. Phys. Chem., 96, 1444. Tateyama, H., Noma, H., Adashi, Y. and Komatsu, M. (1997) Chem. Mater., 9, 766. Templin, M., Friedrich, U., Wiesner, U. and Spiess, H. W. (1999) Precursor-derived Ceramics (eds J. Bill, F. Wakai and F. Aldinger), Wiley-VCH. Tougne, P., Hommel, H., Legrand, A. P., Herlin, N., Luce, M. and Cauchetier, M. (1993) Diamond Relat. Mater., 2, 486. Tuel, A., Hommel, H., Legrand, A. P., Chevallier, Y. and Morawski, J.C. (1990) Colloids. Surf., 45, 413. Vercedia, G., O’Brien, K. L., Schmidt, W. R. and Apple, T. M. (1998) Chem. Mater., 10, 1003. Wagner, G. W., Na, B. -K. and Vannice, M. A. (1989) J. Phys. Chem., 93, 5061. Wakai, F., Kodama, Y., Sakaguchi, S., Murayama, N., Izaki, K. and Niihara, K. (1990) Lett. Nature, 344, 421.

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Page 139 5.2 Paramagnetic defect states in silicon-based nanopowders Abdelhadi Kassiba and Stéphane Charpentier 5.2.1 Introduction Electron paramagnetic resonance (EPR) technique is a suitable tool in characterising the paramagnetic defects in the Si/C/N nanopowders. The EPR experiments are carried out both following the annealing temperature (realised separately up to 1800°C) and the specimen temperature (in situ from 4 to 1000 K). Different EPR frequency bands (4–95 GHz) and numerical EPR spectrum analyses are needed for an understanding of the defect nature and properties. However, when mixed compositions and structures coexist as in (Si/C/N) networks, the overlapping of the resonance lines leads to structureless EPR spectra. In such materials, it is possible to assess the concentration of paramagnetic defects and their properties such as being localised or delocalised. More quantitative analyses can be achieved on more homogeneous materials (Si/C) where structured EPR spectra are obtained. Precise insights on the paramagnetic defects are made and correlated with the material composition, structure and physical properties. The first set of investigations is devoted to SiC nanopowders synthesised by a laser pyrolysis of a gaseous mixture (SiH4, C2H2). The batches, referred as SiC242, SiC163 and SiC177, differ by the residence time of the reactants in the laser beam and by the gaseous precursor flux rate [SiH4]/[C2H2]. The defect concentration is relatively high in all the samples (NS ≈ 1020 spin g−1) and their nature seems correlated with the coexisting SiC polytypes. A consistent numerical analysis of the EPR spectra, over several EPR frequencies, was developed in order to identify the paramagnetic centres. Whatever the annealing stage and the EPR frequency band, three paramagnetic defects with well-defined ( , ) tensors coexist with contributions dependent on the annealing stages. Delocalised centres, thermally induced by varying the sample temperature, are also evidenced in these nanopowders. They would contribute to the significant electric conduction induced by annealing at Ta ≥ 1500°C. The second deals with Si/C/N materials (referred as HMDS40–45) where the coexisting structures (SiC, Si3N4) and the free carbon (C) content give rise to a large unpaired spin concentration brought by two types of paramagnetic centres. Their contributions depend on the synthesis conditions and on the annealing stages. We have shown that the rate of C on nitrogen (N) monitors the defect concentration and the annealing under N partial pressure (nitriding process) favours the creation of unpaired spins in the Si3N4-rich powders. The forthcoming EPR investigations focus on these two classes of silicon-based nanomaterials.

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Page 140 5.2.2 Theoretical background Paramagnetic defects in these Si/C/N covalent networks, consist mainly in unpaired electrons. Under an , each paramagnetic defect is described by an effective spin Hamiltonian which

external magnetic field reads:

where β represents the Bohr magneton, k refers to an interacting nucleus (nuclear spin Ik) with hyperfine diagonal tensor in a reference frame (xk, yk, zk) centred on the (k) nuclei. The tensor, has a diagonal form in its principal (magnetic) frame (X Y Z), with components in the order of ge = 2.0023 value; that is, is defined by angular variables the free electron Landé factor. In crystalline powders, the magnetic field (θ0, φ0) with respect to the magnetic axes (X, Y, Z). The local resonance fields Bres; that is, the EPR line position, is expressed as follows:

where h is the Planck’s constant and νEPR = 9.5 GHz in X-band EPR measurements. Be is the resonance position of a free electron. The EPR signal of paramagnetic defects would have the shape of a continuous absorption background with singularities at the positions

However, in current EPR experiments the derivative of the absorption line is recorded. The small departure Δgi = gi − ge, in a given (i = X, Y, Z) direction, results from a weak spin–orbit coupling (λ · the perturbation formula:

·

) through

and represent respectively the ground and the excited electronic states of energies E(0) and where E(n). Quantitative analysis of the tensor components cannot be undertaken safely in these heterogeneous Si/C/N media. However, the evolution of Δg shifts with particular thermal treatments that can be used to characterise the changes of the local environments around the involved paramagnetic centres. With regard to the hyperfine tensor components , they are determined experimentally from the hyperfine line positions. A brief theoretical justification of these parameters can be established in the following simple case. When the electronic spin density on the nuclei exhibits an axial symmetry around the zk-axis, the components of the

tensor are:

with

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Aiso represents the contact interaction parameter and Aaniso accounts for a dipolar interaction.

and

represent the orbital wave function of (s) and (p) types related to the (k) ion. (r, θ) are the radial and azimuthal coordinates of the unpaired electron in the reference frame centred on the nucleus. gn and βn represent respectively the nuclear g-factor and the nuclear Bohr magneton. The hyperfine lines are observed when the unpaired electron interacts with a nuclear spin in its near environment. For Si/C/N materials, the nuclear spin (I) is brought by 13C by 29Si or by 14N (I = 1, na = 99.6%). The intensity and the number of the hyperfine lines can be used to determine the nature of the interacting nuclei. Moreover, the integrated intensity of the EPR signal is proportional to the paramagnetic defect concentration (Ns) and to their magnetic susceptibility (χ). The temperature dependence of χ informs on the localised or delocalised nature of the paramagnetic species. Indeed, when χ varies with temperature following the Curie law (χ = C/T ), the unpaired electrons are localised in the host material sites. In the case of delocalised unpaired electrons, χ is temperature independent and ‘Pauli-like’ paramagnetism is expected. 5.2.3 Experimental details EPR experiments are carried out on Brucker spectrometer with different frequency bands [S(4 GHz), X (10 GHz), K (24 GHz), Q (34 GHz) and W (95 GHz)]. Investigations at different temperatures (4–300 K) are conducted on an Oxford cryostat. The high unpaired spin concentration in the samples limits the microwave power below 2 mW and the modulation field to 0.5 G. Saturation phenomena and EPR line shape distortions are thus avoided. The absolute concentration (Ns) of paramagnetic species is determined by a comparison with CuSO4 · 5H2O single crystal (≈3·1021 spin g−1). This absolute Ns estimate is within a range of four times more or less than the real value. 5.2.4 Si/C powders Three Si/C batches (SiC163, SiC177 and SiC212) synthesised from gaseous precursors (Table 2.2) are investigated. Whatever the nanopowder series are, the EPR spectra have common features and similar electrical properties. So, the forthcoming report is limited to the SiC212 batch referred below as SiC212 1200– 1800°C according to annealing temperature T = 1200–1800°C. The structural modifications, the powder composition and grain morphology are well known (Charpentier 1997, 1999 and 1999a). Figure 5.2.1(a) shows the room temperature EPR spectra at X-band in the SiC212 1200°C and 1650°C samples. High temperature measurements in the SiC212 1200°C sample up to 600°C under air did not show any change on the EPR spectrum features except for the obvious intensity decrease. The reversible behaviour of the EPR spectra, that is, before and after the oxidation rules out any influence of the oxidation process or thermal removal of the paramagnetic species. Furthermore, it seems that the defects are not connected with hydrogen which may be released by annealing at high temperatures. The defects are intrinsic to the material mainly caused by vacancies in the C or silicon (Si) sites of the coexisting C arrangements and SiC polytypes. Low temperature measurements are needed in order to assess the localised or delocalised nature of the defects. In the SiC212 1200°C sample, the X-band EPR spectrum is drastically

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Figure 5.2.1 (a) Typical X-band EPR spectra in SiC212 nanomaterials annealed at 1200°C and 1650°C. (b) Evolution of the SiC212 1200°C EPR spectrum with the temperature. (c) Experimental and calculated W-band EPR spectra in the SiC212 1200°C sample. The contributions from the different paramagnetic centres are also shown (relative contributions to the EPR spectrum DI: 21%; DII 49%; DIII 30%).

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Figure 5.2.2 Evolution with the temperature of the EPR signal integrated intensity weighted by T: (a) in the SiC212 1200°C sample and (b) in the SiC212 1600°C sample. broadened by lowering the temperature (Figure 5.2.1(b)). This is understood by the possible defect motion which is frozen at low temperature. A double integration of the EPR signal allows the determination of the overall intensity which is proportional to the spin susceptibility χ (Figure 5.2.2). A departure of χ(T) from the Curie law above 150 K, also observed in the SiC212 1650°C sample (Figure 5.2.2), indicates the presence of delocalised paramagnetic species. Furthermore, the peak-shaped behaviour of the product χ(T) × T around 30 K is consistent with a localisation of the defects and correlated unpaired spins. Consistent numerical analyses of the EPR spectra are hereafter conducted in order to determine the paramagnetic defect features in these SiC nanopowders. The paramagnetic centres are S = 1/2-defects file:///C|/Users/steppenwolf/Documents/educatio...ks/to%20burn/nanocomposites/files/page_143.html (1 of 2) [20/01/2010 10:04:35 AM]

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assumed to interact with p equivalent nuclei (nuclear spin

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Page 144 . Under an applied magnetic field

, the spin Hamiltonian is given by:

represents the hyperfine coupling tensor between the electronic and nuclear spins. The magnetic dipole transition intensities and their resonance positions are determined numerically from The basis is limited to

states with

due to the multiplicity of electronic and

nuclear spins in SiC materials. On the other hand, the na (=0.047) of the 29Si and (na = 0.01) for 13C, giving rise to a hyperfine coupling, impose the EPR spectrum of a given paramagnetic defect to be a superposition of the two contributions. One is related to a spectrum with a hyperfine interaction, weighted by the p × na, superimposed to a spectrum without the hyperfine interaction as for 28Si and 12C nuclei weighted by (100 − p × na). A homemade program for EPR spectrum adjustments in crystalline powders was developed. All the magnetic field orientations (θ0, φ0) are allowed (0 ≤ θ0 ≤ π/2, 0 ≤ φ0 ≤ π) due to the random orientations of the crystalline sites in the SiC powders. The calculated resonance positions are used with suitable line widths in order to adjust the experimental EPR spectra. Whatever the annealing temperature or the operating νEPR frequency bands are, the EPR spectra are accounted by using three paramagnetic defects with the spectroscopic data ( , ) summarised below at room temperature: • DI centres: • DII centres: • DIII centres: no hyperfine coupling. Furthermore, the EPR spectra in X-band are characterised by individual resonance line widths less than 2 G for DI and DII species and close to 8 G for DIII centres. The relevance of the EPR signal adjustments is shown for the SiC212/1200°C sample in the W-band, highly sensitive to the anisotropy of the defect tensors (Figure 5.2.1(c)). The hyperfine constants are determined from the S-band measurements where the hyperfine lines are well resolved. 5.2.5 Assignment of the paramagnetic species in Si/C nanopowders tensors, the behaviour with the sample temperature of the EPR spectrum line The components of and width and integrated intensity are the relevant information in identifying the paramagnetic species. The anisotropic DI centres are characterised by a tensor with an axial symmetry along X-direction (Δgx = gx − ge = 0.0021 and Δgy = gz = 0). Qualitatively, such shifts are not consistent with the assignment of the unpaired spins to dangling bonds (DB). Indeed, taking the DB along X-direction, a symmetry argument implies that the shifts would be Δgx = 0 and gy, gz ≠ 0. This conclusion is supported by the hyperfine coupling constant (A = 11·10−4 cm−1) which is proportional to the residence probability of the unpaired electron on the Si or C nucleus. The theoretical hyperfine constants for free ions are

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Page 145 (Luchsinger et al. 1997): and From the above EPR spectrum analysis, the experimental isotropic hyperfine tensor for DI and DII centres gives is:

Then, the probability of the unpaired electron being in the s orbital of Si or C

No more than 1% of the spin density is accounted for by the localisation on one Si or C nucleus. So, the paramagnetic defects DI and DII in these SiC nanopowders can be associated to unpaired electrons delocalised inside the cubic and hexagonal crystalline sites of the coexisting SiC polytypes (3C—SiC and αSiC). The possible origin of these paramagnetic centres consists in charged vacancies in the SiC sites favoured by the initial C in excess in these materials (Kityk 2000). The elimination of the DI centres by annealing above 1500°C (Figure 5.2.3) seems correlated with the structural modification α-SiC → 3C—SiC. The paramagnetic centre DIII is characterised by an isotropic tensor with components close to 2.0031(2), that is, close to the average value ((gx + gy + gz)/3) of DI or DII defects. The EPR spectrum contribution from DIII centres consists of a broad and unstructured resonance line (Figure 5.2.1(c)). Such features can result from defects involved in amorphous or disordered regions. However, by varying νEPR, we have not observed any line broadening

Figure 5.2.3 Variation with annealing of the absolute contributions from DI, DII and DIII paramagnetic centres. The vanishing of the DI contribution in the SiC212 1600°C sample coincides with the structural change α → β-SiC.

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Figure 5.2.4 Evolution with the EPR frequency band of the relative contributions of the different paramagnetic species. DI exhibits a stationary variation while that of DII and DIII indicates a conversion between these centres. (ΔHpp) connected with a spread of the

components which may be caused by a structural disorder

around DIII defects. Indeed, G in X-band changes only to of g-values works, the ratio between the line widths will be close to

G in W-band. If a spread

Such a criterion rules out the location of DIII defects in disordered regions. On the other hand, when νEPR increases, that is, for short EPR characteristic times, the relative proportion of DII and DIII exhibits an opposite evolution (Figure 5.2.4). Similar behaviours are observed by varying the SiC212 1200°C and 1650°C sample temperature. As shown in Figure 5.2.5, the relative contributions to the EPR spectrum from the paramagnetic centres exhibit an anti-coincidence behaviour between the anisotropic (DI, DII) and the isotropic DIII centres (Figure 5.2.5). These results are understood by a conversion between the localised paramagnetic DII species and the DIII ones. The departure from the Curie law of the overall spin susceptibility would result from the DIII centres being delocalised in the material. This can be partly responsible for the significant macroscopic electric conduction which appears in these nanopowders annealed above 1500°C (Kassiba 2000) as discussed in Chapter 6. To sum up, EPR gives relevant information on the paramagnetic defects in SiC materials. Consistent analyses can be achieved when the EPR signal exhibits resolved details depending on the operating conditions (νEPR, T, etc.). A correlation can be made between the delocalised paramagnetic centres and the electrical properties in SiC nanopowders as reported in Chapter 6. 5.2.6 Silicon carbonitride (Si/C/N) powders X-band EPR spectra in HMDS series (Table 2.5) exhibit similar features in all the as-formed powders. The spectra are composed of a broad Gaussian line (peak to peak line width

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Figure 5.2.5 Variation with the sample temperature of the relative contributions of the anisotropic (DI, DII) and the isotropic (DIII) centres. A conversion between defects is evidenced in the (a) SiC212 1200°C sample and in the (b) SiC212 1600°C sample. ΔHpp = 9 G) superimposed on a narrow Lorentzian line (ΔHpp = 4 G). The resonance position at g = 2.0030 (2), is reported in a wide class of (Si, C, N) based materials where dangling bonds (DBs) are attached to C or N atoms (Warren 1990; Demichelis 1991; Suzuki 1995 and Xingguo Li 1996). Measurements down to 50 K in the as-formed HMDS40 powder did not show any resolved hyperfine lines and any effect due to a slowing down of the paramagnetic centre motions. The integrated EPR spectrum intensity varies with the temperature according to I (T)= C/T +Ip (Figure 5.2.6). This is supported by a coexistence of localised defects giving rise to the (C/T) contribution and delocalised ones with a temperature-independent contribution (Ip). The latter is connected with Pauli-like paramagnetism which can originate from delocalised electrons on the free C arrangement present in a high concentration in the as-formed powders.

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Figure 5.2.6 Evolution v. the temperature of the product I * T, I represents the EPR signal integrated intensity. (●) HMDS45 annealed at 1500°C in 50% Ar and 50% N partial pressure. (▲) HMDS40 as-formed powder. The slope of the curves is connected with the importance of delocalised paramagnetic species.

Figure 5.2.7 Typical X-band EPR spectra in HMDS40–45 annealed at 1500°C in 50% Ar and 50% N partial pressure. In HMDS45, Si3N4-rich material, a creation of a large defect concentration is induced by the annealing and nitriding process. With regard to the behaviour with annealing of the paramagnetic centres in the HMDS series, 1500°C is a critical annealing temperature. Indeed, while the EPR broad Gaussian component exhibits an enhanced intensity only in the HMDS43–45 batches (Figure 5.2.7), the sharp Lorentzian line is dramatically reduced in all the batches after annealing at 1500°C

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Figure 5.2.8 Paramagnetic centre concentrations: (a) in the as-formed HMDS40–45 powders and (b) in the


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annealed powders at 1500°C. Black and white bars refer respectively to the contribution from two types of defects involved in these materials.

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Figure 5.2.9 EPR spectra in HMDS42–44 v. the annealing temperature. (Figure 5.2.8(b)). A careful analysis of the EPR spectra in HMDS44–45, reveals that the spectra consist of a sharp line and shoulders which undergo a better resolution after annealing above 1500°C (Figure 5.2.9, El Kortobi 1998). Since these materials are N rich (a low C/N atomic ratio), the observed EPR spectrum shoulders would originate from hyperfine interactions between unpaired spins and 14N nuclei. A crude EPR signal deconvolution in HMDS44, annealed at 1600°C, is made in terms of a sharp central line and three equivalent and equidistant lines with a separation of about 5 G (Figure 5.2.10). According to EPR investigations of DB centres in amorphous Si nitride (Warren et al. 1991), Si DBs interacting with the nearest N neighbours are characterised by tensor (g = 2.0028, g⊥ = 2.0032) and by 4.6 G as a hyperfine coupling constant. Furthermore, (SiN3)− are the most thermally stable DBs in SiC—Si3N4 networks heat treated above 1673 K under partial N pressure (Li et al. 1996). So, in HMDS44–45 annealed above 1500°C, DBs such as (SiN3)− would probably bring the large concentration of the paramagnetic centres. Their magnetic susceptibility follows the Curie law in the temperature range (4–100 K) (Figure 5.2.9, El Kortobi 1998) and indicates that the involved DBs are localised. Structural reorganisations and composition changes of the HMDS networks after annealing at 1500°C were noticed by NMR investigations (Chapter 5.1). The behaviour of the paramagnetic

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Figure 5.2.10 Comparison between the experimental EPR spectrum in HMDS44 annealed at 1600°C and the calculated ones by a superposition of a central line (C1) and triplet lines (N1, N2, N3). (C1) is associated to paramagnetic centres implying C atoms. (N1, N2, N3) lines are induced by hyperfine coupling of unpaired electrons on Si atoms with equivalent N nuclei in centres such as (SiN3)−. centres seems intimately connected with these modifications. What can be addressed below is the correlation between the characteristic EPR lines, that is, the paramagnetic species and the sample composition. On one hand, the narrow Lorentzian EPR line can be partially caused by delocalised paramagnetic defects as suggested by the temperature variation of the EPR integrated intensity (Figure 5.2.11). Motional narrowing and exchange between localised and delocalised spins by a thermal activation were reported in C black particles embedded in a polymer matrix (Adriaanse et al. 1997). Similarly, in C rich material (HMDS40–41– 42), the importance of the narrow Lorentzian lines would result from DBs and delocalised unpaired electrons on C atoms. On the other hand, the broad EPR line is consistent with paramagnetic species implying N atoms. Indeed, when the material composition is 90% of Si3N4 type (HMDS44 annealed at 1600°C), the experimental spectrum is characterised by an isotropic g tensor (2.0030(2)) and hyperfine triplet lines due to the coupling of defects with I (14N) nuclear spins. The large spin concentration created by annealing under N partial pressure can be explained by the large rate of structural defects induced by the nitriding process and structured reorganisation. These features are similar to those reported in ultrafine SiC—Si3N4 composite powders (Li et al. 1996). According to these authors, annealing above 1500°C gives rise to different behaviours of the EPR signal depending on the atmosphere used. Under argon, the defect concentration decreases with annealing while the opposite behaviour is obtained for annealing under Ar and 50% N2. We have shown that more homogeneous materials are obtained by annealing at 1500°C and above, HMDS40–41–42 (SiC-rich materials) and HMDS43–44–45 (N-rich materials).

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Figure 5.2.11 Integrated EPR line intensity in HMDS44 annealed at 1600°C v. the inverse sample temperature (1/T ). The Curie law is revealed and indicates the localised nature of the coexisting paramagnetic centres. A large concentration of paramagnetic defects is created in HMDS45–44–43 consecutive to the annealing and the nitriding process. The opposite behaviour is observed in HMDS40–41–42 where a dominant SiC structure is realised as well as the removal of the free C content. References Adriaanse, L. J., Brom, H. B., Michels, M. A. J. and Brokken-Zijp, J. C. M. (1997) Phys. Rev. B55, 15, 9383. Charpentier, S., Kassiba, A., Emery, J. and Cauchetier, M. (1999) J. Phys.: Condens. Matter, 11, 4887. Charpentier, S., Kassiba, A., Bulou, A., Monthioux, M. and Cauchetier, M. (1999a) EPJ: AP, 8, 111. Charpentier, S., Kassiba, A., Fusil, S., Armand, X., Cauchetier, M., Fayet, J. C. and Emery, J. (1997) Appl. Magn. Reson., 12, 255. Demichelis, F., Pirri, C. F., Tresso, E., Rigato, V. and DellaMea, G. (1991) J. Non-Cryst. Solids, 128, 133. El Kortobi, Youssef (1998) Thesis, Université P. M. Curie, Paris, France. http://corail.sudoc.abes.fr80 & SudocCatalogue no 053764552. Kassiba, A., Tabellout, M., Charpentier, S., Herlin, N. and Emery, J. R. (2000) Solid State Commun., 115, 389. Kityk, I., Kassiba, A., Tuesu, K., Charpentier, S., Ling, Y. and Makowska-Janusik, M. (2000) Mater. Sci. Eng. B, 77(2), 147. Li, X., Chiba, A., Nakata, Y., Nagai, H. and Suzuki, M. (1996) Mater. Sci. Eng., A219 95. Luchsinger, R. H., Yu Zhou and Meier, P. F. (1997) Phys. Rev. B55, 1, 6927. Suzuki, M., Hasegawa, Y., Aizawa, M., Nakat, Y. and Okutani, T. (1995) J. Am. Ceram. Soc. 8,1. Warren, W. L., Lenahan, P. M. and Curry, S. E. (1990) Phys. Rev. Lett. 65, 2, 207.

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Page 153 5.3 Electronic structure studied by X-ray photoemission (XPS) and soft-X-ray (SXS) spectroscopies Adriana Gheorghiu de La Rocque, Georges Dufour, François Ténégal and Christiane Sénémaud 5.3.1 Principles X-ray Photoemission Spectroscopy (XPS) provides the electronic distributions of the core levels and valence band (VB) of solids. As concerned the VB states, energy distribution curves (EDCs) correspond to the total density of states (DOS) modulated by appropriate photoionisation cross-sections:

σx being the cross-section and Nx(E) the partial DOS of x state. Thus, in some cases, information concerning particular symmetry states can be arbitrarily enhanced. Due to the small escape depth of emitted photoelectrons (0.5–10 nm), the experiments are sensitive to the superficial layer of the samples. X-ray spectroscopy probes separately the valence and the conduction DOS. In X-ray Emission (XES), radiative transitions between the VB and a localised inner vacancy are investigated. Initial core holes are created by means of either electron impact or X-ray photons, possibly synchrotron beam. The radiative de-excitations of the inner vacancy are governed by dipole selection rules which imply that only valence states with particular angular momentum are involved. Moreover, as the core level concerns a specific atomic site, the DOS around each atomic species of a compound are investigated separately. The information obtained by XES concerns the volume. The spectral intensity of an X-ray emission is, in first approximation:

where Mif(E) is the matrix element which does not affect the main features of an XES spectrum over a is the DOS with (1±1) symmetry; L(hν–E) is the Lorentzian distribution of limited spectral range; the core hole. By X-ray absorption spectroscopy, transitions between a core level and the conduction band states are analysed. In the spectral region close to an absorption edge: X-ray Absorption Near Edge Structure (XANES), the photoabsorption coefficient μ is in first approximation proportional to the density of the empty states Nunocc above the Fermi level (EF), with particular angular momentum around a specific site:

By XANES either the volume or only the surface of the sample is concerned according to the technique chosen (Total Electron Yield (TEY), transmission or Fluorescence Yield (FY)).

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Page 154 For both XES and XANES spectra, the contribution of the Lorentzian profile of the core hole induces a broadening of the spectra, which limits the resolution of the method and implies to study only transitions involving a sufficiently sharp core level. For a valuable comparison of XPS and SXS data, it is particularly important to set both XPS–XES (or XPS– XANES) experimental curves in a common energy scale referring to the Fermi level. This is directly obtained in XPS as EDCs are referred to EF. In the case of XES and XANES spectra, the spectral curves are obtained in a transition energy scale (EVB–EX) or (ECB–EX) and the energy position of the Fermi level can be obtained by determining the core level binding energy (BE) EX, provided that relaxation effects are neglected. 5.3.2 Experiments 5.3.2.1 X-ray Photoemission Spectroscopy (XPS) XPS spectra were obtained by using Mg Kα incident radiation (hν = 1253.6 eV). An electrostatic hemispherical analyser was used in the fixed analyser transmission mode (FAT) with a pass energy equal to 40 eV. Under these conditions, the full-width at half-maximum (FWHM) of the Ag 3d5/2 line is 1 eV. The core level spectra were scanned with a 0.1 eV step and the XPSVB with a 0.2 eV step. In some cases, an energy shift of the photoelectron peak due to a charging effect was observed (Shirley 1972; Jentoft et al. 1999). In order to take this effect into account, we calibrate the BE scale with reference to the C 1s line observed from hydrocarbonated species, which are always present at the surface of the sample. The C 1s line is set at 285.0 eV. In our case, the corrections are more complicated because carbon (C) is also present in the studied materials. The energy scale was then adjusted so that O 1s peak, which corresponds essentially to silicon dioxide, coincides approximately for all spectra. In this procedure, it is assumed that oxygen atoms are essentially present in silicon dioxide bonding. Let us note that in our experimental work the oxygen proportion measured by XPS is only a few percent in as-formed (AF) samples and slightly higher, 10–15%, in annealed samples, due to surface contamination (Gheorghiu et al. 1997). The interpretation of the core level lines is based on a decomposition of each spectral distribution into Voigt functions; the spectra are superposed to a background due to the inelastic electron scattering. For a mathematical treatment of core level spectra, a subtraction of this background is performed by assuming that its variation is linear or polynomial. No correction of background was made in the case of the VB spectra. 5.3.2.2 Soft X-ray Spectroscopies (XES–XANES) In the case of Si/C/N powders, we measured the Si Kβ (3p → 1s) emission line which probes the Si 3p VB states and Si K absorption line (1s → np) which provides the Si p conduction band (CB) states. The X-ray emission spectra Si Kβ were analysed with a bent-crystal vacuum spectro-meter (Sénémaud et al. crystal (2d = 850.766 pm) in the first order of 1989) using as a monochromator either a quartz reflection, or a gypsum crystal (020) (2d = 1519.84 pm) in the second order of reflection bent to a radius of 250 mm. The detector was a proportional counter (500 mb Ar/CH4 90/10) with an entrance slit fixed to 50 μm. The spectra were scanned with a 0.2 eV wide step. The total instrumental broadening, mainly due to the monochromator itself, is estimated to be about 0.3 eV (Sénémaud and Ardelean 1990). The emissive target was prepared by depositing a thin layer of laser synthetised powder onto

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Page 155 a copper anode. The spectra were excited by electron bombardment with exciting conditions 5 kV, 5 mA. The thickness sampled is determined by the penetration of the incident electron beam. In our experimental conditions, we estimate the thickness of the emitting layer to about 0.5 μm. Thus in this case the whole volume of the powder particles is concerned. The near edge photoabsorption (XANES) measurements were carried out at Laboratoire pour l’Utilisation du Rayonnement Electromagnétique – Orsay (LURE). The Super-Aco storage ring was operating at 800 MeV, and the average current was approximately 200 mA. A two-crystal monochromator equipped with InSb (111) crystals (2d = 748.06 pm) was used. The photoabsorption spectra corresponding to the Si K-edge were obtained in the transmission geometry and scanned with a 0.2 eV-wide step. The detector was a proportional counter filled with a 300 mb Ar/CH4 (90/10) mixture. The absorbing screens were obtained by depositing powders onto 5-μm-thick millipore polycarbonate membranes. The spectra of the incident beam were obtained through a membrane without powder. The instrumental resolution is evaluated to be approximately 0.4 eV (Bouisset et al. 1991). 5.3.3 Study by XPS–SXS of Si, Si/C, Si/N and Si/C/N laser synthetised nanopowders produced from gaseous precursors The samples were prepared by laser pyrolysis of appropriate precursors in the same conditions as reported previously (Gheorghiu et al. 1992), summarised in Table 2.2. 5.3.3.1 Core levels The Si 2p spectrum of Si50 powder (Figure 5.3.1) exhibits a strong asymmetric line accompanied on the highBE side by a low intensity component due to SiO2. This oxide contribution has approximately the same relative intensity as that of a native oxide layer at the surface of a monocrystalline Si wafer (Himpsel et al. 1988). The main peak, located at 99.3 eV,

Figure 5.3.1 Si 2p core levels for Si50 sample.

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Figure 5.3.2 Si 2p core levels for (a) SiC151 and (b) SiC152a samples. corresponds to the 2p1/2,3/2 doublet from Si—Si bonds and is well fitted by a sum of two Voigt lines separated by 0.6 eV, with an intensity ratio 1/2 which accounts for the spin–orbit splitting. The total width at half-maximum of the line is 1.2 eV, close to the value obtained with a single-crystal wafer and in the same experimental conditions: 1.1 eV (Rochet et al. 1986). The remaining intensity in between Si and SiO2 lines reveals probably the existence of an intermediate oxide SiOX (Hollinger et al. 1988). A tailing could result from inhomogeneities in the sample. For pure Si powders, the XPS results show that the sample is a wellorganised silicon (Si), which is in agreement with X-ray diffraction (XRD) results (Table 2.2). In Si/C samples, SiC151 and SiC152a (Figures 5.3.2(a) and (b), the Si 2p spectra have identical shapes: quasisymmetric with a high BE tail. The decomposition achieved on the basis of two components: one at the energy position of the maximum intensity, the second at higher BE, corresponding to SiO2 (102.9 ± 0.1eV), shows that the energy of the main line is located at 101.3 and 101.2 eV for SiC151 and SiC152a, respectively. The Si 2p line is slightly sharper in SiC151 than in SiC152a. Finally, a low proportion of SiO2 is present at the surface for both samples. It is noteworthy that the intensity is negligible at the BE corresponding to Si—Si bonds. Thus in Si/C, heteropolar Si—C bonds are present, without any segregation of Si atoms. In SiC151, the local order is slightly better than in SiC152a and commercial β-SiC (Driss-Khodja et al. 1992) (fewer bond angle and bond length fluctuations). Consequently, XRD measurements led to the same conclusions (Cauchetier et al. 1988). The C 1s BE from powders is close to the value that characterised the C—Si bonds in β-SiC (Driss-Khodja et al. 1992). The C 1s line is shifted towards high BE in powder samples as compared to β-SiC. As the energy distance E(Si2p) − E(C1s) varies, this effect cannot be interpreted in terms of a Fermi-level shift but rather by the presence of H atoms in the network, in agreement with previous measurements (Fang and Ley 1989). This is in agreement with the identification of Si Hn groups by infrared (IR) spectroscopy in samples prepared under the same conditions (Cauchetier et al. 1990). For Si/C powder materials, the XPS measurements reveal that only Si—C bonds are present in Si—C4 and C— Si4 groups and that there is no segregation of Si atoms in the network. They reveal, moreover, that the samples contain a very low proportion of oxide. The atomic

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Figure 5.3.3 Si 2p core levels for (a) SiN7 sample and (b) α-Si3N4 powder. structure and the long-range order of the material depend on the preparation conditions. Thus, an increase of the laser power used for the powder synthesis from 220 to 600 W (Table 2.2) induces an increase of the flame temperature and consequently of the local order; the Si 2p becomes slightly sharper. The Si 2p peaks observed from the SiN7 sample (Figure 5.3.3(a)) compared to commercial α-Si3N4 (Figure 5.3.3(b)) show that, in both cases, the spectrum consists of a wide, asymmetric line, the high-binding energy tail resulting from the presence of SiO2. The decomposition of the α-Si3N4 spectrum in two components presented in Figure 5.3.3(b) shows a main Si 2p component, 1.7 eV wide at 101.8 eV BE, accompanied by a contribution from SiO2 at 103.1 eV having an intensity of about 11%. In this case the spin–orbit splitting of Si 2p is not observed due to the increase of Si 2p FWHM as compared to pure Si. The Si 2p spectrum from SiN7 sample has a similar shape but is significantly broader than the commercial α-Si3N4 one. The peak decomposition gives a main peak, 2.0 eV wide, located at 101.7 eV, and a SiO2 contribution with a relative intensity of about 12% (Figure 5.3.3(a)). Moreover, a low BE feature is present at 99.0 eV which we interpret as due to the existence of a low proportion (about 2%) of Si—Si clusters. Thus, in the SiN7 powder, most Si atoms are bonded to N atoms in Si—N4 units which is characteristic of the silicon nitride network. The broadening of the corresponding Si 2p component as compared to the α-Si3N4 one (2.0 instead of 1.7 eV) can be attributed to the existence of structural disorder in the network, introduced by fluctuations in the bond lengths and/or bond angles as compared to their values in the crystalline phase. The presence of H atoms in the material is possible and could contribute to such a disorder. The N 1s line is observed at approximately the same BE value (397.7 eV) in both the powders and the stoichiometric α-Si3N4 samples indicating the N—Si3 configuration, but the width is significantly broader in the powder than in crystalline α-Si3N4 (1.7 eV). This result confirms the existence of a chemical order in the powder sample, the broadening of the line being attributed to the local topological disorder, in agreement with the above results concerning the Si 2p line. The existence of N—H bonds can contribute to a broadening of the N 1s line as previously observed in Si—NX: H samples prepared by photo-chemical-vapour deposition (PCVD) (Driss-Khodja et al. 1989).

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Figure 5.3.4 Si 2p core levels for (a) SiCN12 and (b) SiCN27 samples. We conclude that the SiN7 powder contains mainly Si—N4 groups, N atoms being surrounded by three Si atoms. The local bonding is similar to that of silicon nitride α-Si3N4. A few Si atoms are involved in Si—Si bonds; the proportion of oxide is very low (12% in the powder). The presence of Si—Si bonds in the network, as well as the possible existence of hydrogen in the sample, can contribute to the introduction of short-range disorder around both Si and N atoms. The Si 2p spectra from the two Si/C/N samples, SiCN12 and SiCN27 (Figures 5.3.4(a) and (b)) consist of a rather broad asymmetric line with different FWHM: 2.2 and 1.7 eV, respectively. According to the conditions of decomposition, we infer that, besides the small SiO2 contribution around 103.3 eV, the SiCN12 spectrum results from the contribution of two main lines, while SiCN27 spectrum corresponds to a single Voigt line. In the case of SiCN12 spectrum, the two Voigt curves, 1.5 and 1.7 eV wide, are located at 101.1 and 101.9 eV, respectively; the intensity ratio of the components is about 0.7. These two components are close to Si 2p lines from β-SiC and α-Si3N4, respectively. An additional line, lower in intensity, at 99.6 eV can be attributed to Si—Si. In contrast, the main peak of SiCN27 is 1.7 eV wide and is located at 101.6 eV. It is accompanied by two smaller components at 99.4 and 103.3 eV that we attribute to Si—Si bonds and SiO2, respectively. This result clearly shows that the local bonding is quite different in the two Si/C/N samples under consideration. In one case, SiCN27, for which C/N = 0.22, most Si atoms have the same local structure SiCxN4−x, Si being mainly bonded to N atoms. The value of the FWHM is the same as in the crystalline αSi3N4 and reveals that the network is well ordered. It is noteworthy that the Si 2p line from this sample is sharper than that from the Si/N7 sample. Consequently, the modifications of the gas mixture (addition of CH3NH2), which leads to an increase of temperature, contributes to give a better ordered material. In the SiCN12 sample, with C/N = 0.93, our results show clearly that both Si—N4 and Si—C4 groups are present, with approximately the same proportion. The C 1s BE corresponding to C—Si bonds is practically the same in both samples. However, the FWHM increases significantly in mixed compounds as compared to Si/C, indicating a loss of local order around C atoms induced by the presence of N atoms.

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Page 159 The N 1s BE value remains the same for all samples containing N atoms, revealing a preferential N—Si3 arrangement in Si/C/N samples. A decrease of its FWHM is observed for mixed compounds as compared to Si/ N, which confirms the increase of local order induced by C atoms. In the case of composite Si/C/N, our results show that different materials can be prepared depending on the preparation conditions, especially the initial gases concentration. The powders are formed either by a mixture of Si—C4 and Si—N4 units or by a single type of local Si environment. For a C/N ratio of about 1 in the initial gas mixture, the existence of two types of local Si bonding, Si—C4 and Si—N4, with a proportion of about 0.7–1 is revealed by our XPS measurements; a few Si atoms are involved in Si Si bonds. For an initial C/N ratio in the gas mixture equal to 0.22, the XPS spectra reveal that most Si atoms are involved in a single type of chemical environment; the Si 2p core level has a BE close to the value corresponding to α-Si3N4; the Si—C bonds are also present, as revealed by the C 1s line. Consequently our results suggest the existence of local C—Si—N3 arrangement around Si atoms. 5.3.3.2 XPS valence band results The VB spectra of powders Si/C (SiC151), Si/N (SiN7) and Si/C/N (SiCN12, SiCN27) exhibit a broadband (Figure 5.3.5) noted A extending between 0 and about 13 eV, followed by peaks B and C, at about 15 and 19 eV, respectively, which have different relative intensities in the considered samples. The distances from these features to EF are given in Table 5.3.1. The SiC151 VB spectrum is almost identical to a β-SiC spectrum, as already reported (Driss-Khodja et al. 1992). Part A exhibits two peaks labelled A2 and A4 located at 5.7 and 9.5 eV, respectively. From a comparison with theoretical densities of states (Robertson 1992) these peaks are attributed to C 2p states mixed with Si 3p and Si 3s, respectively; a wide minimum separates A4 from peak B, located at 14.4 eV BE, which corresponds to C 2s states mixed to Si 3p and Si 3s. The SiN7 spectrum is very similar to the silicon nitride (Johnson Mattey, 99.99% powder) XPSVB spectrum previously reported (Sénémaud et al. 1993). The upper part A exhibits

Figure 5.3.5 XPS VB spectra of laser-synthesised SiC151, SiN7, SiCN12 and SiCN27 samples.

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Page 160 Table 5.3.1 Energy positions of the XPSVB features (eV) for laser synthesised samples A1 A2 A′ A3 A4 A″ A5 B C SiC151 5.7 9.5 14.4 SiN7 3.3 7.5 11.4 18.9 SiCN12 7.0 10.8 14.9 19.3 SiCN27 3.9 7.6 11.4 19.0 three structures labelled A1, A3 and A5 located at 3.3, 7.5 and 11.4 eV, respectively. From the previous study of Si3N4 XPS and SXS spectra (Sénémaud et al. 1993), in agreement with theoretical DOS(Robertson 1992), peak A1 is unambiguously attributed to N 2pπ lone pair, peaks A3, A5 are associated respectively to N 2p–Si 3p states and N 2p–Si 3s states. Peak C, well separated from the A band, is due to N 2s states mixed with Si 3s–Si 3p states. The VB spectra from SiCN12 and SiCN27 are quite different in shape (Driss-Khodja 1993). The spectrum from sample (SiCN12) with C/N = 0.93 exhibits only two structures A′ and A″ at 7.0 and 10.8 eV. The energy position of A′ is intermediate between that of A2 from SiC151 and A3 from SiN7. Similarly, A″ is located between A4 (SiC151) and A5 (SiN7) (Table 5.3.1). The low BE edge of the spectrum is more contrasted than in the SiN7 spectrum; however, it is less abrupt than in SiC151. It is noteworthy that both peaks B and C are observed in this spectrum at 14.9 and 19.3 eV. Their energy positions correspond to those of peaks B from SiC151 and C from SiN7 suggesting that both C 2s and N 2s mixed to Si 3s, 3p states are present. For SiN27, which corresponds to C/N = 0.22, the spectrum is rather similar to the SiN7 one. The upper band A exhibits three structures (see Table 5.3.1); it is followed by a sharp peak C located at 19.0 eV, well separated from band A; consequently, the VB states correspond essentially to those of Si—N bonds found in silicon nitride. Let us note that the minimum between A and C is not so marked as in SiN7 spectrum and that a shoulder is clearly observed around 15 eV. 5.3.3.3 SXS-XPS results In Figures 5.3.6–5.3.9, we have plotted, in the same BE scale with EF as origin, the XES Si Kβ, the XPS VB and the XANES spectra for powders SiC151, SiN7 and SiCN12 and SiCN27. Let us recall that the XES Si Kβ curves give the Si 3p VB states distribution and the XANES curves the unoccupied Si p states distribution. Si Kβ from SiC151 powder (Figure 5.3.6) is very similar to silicon carbide Si Kβ spectrum observed previously (Driss-Khodja et al. 1992; Zahorowski et al. 1989). The maximum intensity located at 5.3 eV from EF corresponds to peak A2 of XPSVB spectrum, the decrease of intensity towards EF is rather abrupt; its inflection point coincides with that of the low BE edge of XPSVB spectrum. The high BE edge is less steep and shows a shoulder at about half-maximum, that is at an energy corresponding approximately to the energy of peak A4. Finally a small intensity peak, well resolved from the main band, is observed at 15 eV from EF, showing that Si 3p states are present deeper in the VB. This small peak coincides with peak B of XPSVB spectrum. The XANES spectrum of SiC151 presents a two-step edge extending from about 0 to 5.3 eV, where a sharp maximum occurs. A secondary maximum occurs at 11.5–12 eV and a strong peak at 19 eV from EF.

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Figure 5.3.6 XPSVB, XES and XAS spectra from SiC151 sample adjusted in a common energy scale with EF as origin.

Figure 5.3.7 XPSVB, XES and XAS spectra from SiN7 sample adjusted in a common energy scale with EF as origin. Si Kβ from SiN7 (Figure 5.3.7) is a quasi-symetric band; its overall shape is close to that of silicon nitride Kβ emission band previously reported (Sénémaud et al. 1993); however, its FWHM is wider, 6.5 instead of 5.5 eV. The energy position of the maximum, 8.4 eV from EF gives the mean BE of the Si 3p states; it is the same as in the nitride. This maximum coincides approximately with peak A3 of XPSVB spectrum. Consequently, this result confirms that the peak A1 of XPSVB spectrum corresponds to pure N states. We note that the decrease of intensity on the low BE tail of Si—Kβ is not steep but exhibits some tailing; we attribute this effect to the presence of a few Si Si bonds in this sample. Let us recall that these like-atom bonds had already been revealed by the study of Si 2p core level from the same sample (Figure 5.3.3(a)). The low intensity peak located at 19.5 eV is in coincidence with peak C of XPSVB spectrum and reflects the presence of Si 3p states mixed to N 2s ones in this part

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Page 162 of the VB. The XANES spectrum of SiN7 consists of a steep edge whose inflection point is located 1 eV above EF, followed by a maximum at 3.8 eV from EF. A second maximum is observed around 22 eV from EF. In SiCN12 and SiCN27, the Si Kβ emission band is a quasi-symmetric band with FWHM 7.7 and 6.5 eV, respectively (Figures 5.3.8 and 5.3.9) (Driss-Khodja 1993). In SiCN12 sample, the main peak is accompanied by a shoulder at 8.5 eV, in coincidence with SiN7 Kβ maximum; a faint feature around 5 eV could be a signature of SiC151 Si Kβ maximum. The distance from EF of the low BE edge, measured at half amplitude, is about the same as in SiC151. The SiKβ spectrum is broader than in binary compounds and can be considered approximately as a sum of both the SiC151 and the SiN7 spectral distributions. It is noteworthy that, in the high BE range, a small feature is observed at the energy positions of B. The photoabsorption edge of SiCN12 has a continuous slope, slightly less abrupt than in SiN7 spectrum. Beyond the first maximum, two peaks are observed at about 3.8 and 5.3 eV from EF, that is, at approximately

Figure 5.3.8 XPS VB, XES and XAS spectra from SiCN12 sample adjusted in a common energy scale with EF as origin.

Figure 5.3.9 XPS VB, XES and XAS spectra from SiCN27 sample adjusted in a common energy scale with EF as origin.

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Page 163 the same position as in SiC151 and SiN7 photoabsorption curves, respectively. Beyond this energy range, a slight feature is noted at 11.5 eV from EF and a strong maximum at 19.6 eV from EF. In SiCN27, the Si Kβ spectrum is narrower than in SiCN12; its shape is rather similar to SiN7 Si Kβ. The maximum is approximately in coincidence with peak A3 of the XPS VB curve, while no feature appears at the energy position of A1. The photoabsorption curve exhibits a strong maximum at 3.6 eV from EF, that is, close to that observed in SiN7 photoabsorption, giving the appearance of a ‘white line’. A shoulder is observed at 7 eV and a wide maximum at 21 eV from EF. 5.3.3.4 Discussion Complementary information on the valence and conduction states of solids can be obtained by using both SXS and XPS spectroscopies. XES and XANES spectra correspond in first approximation to the convolution of the Si p occupied or unoccupied state distributions by the Lorentzian distribution of the Si 1s core-level distribution, which is about 0.5 eV wide (Sénémaud et al. 1993). The XPS spectra correspond to the total valence distribution modulated by the photoionisation cross-sections which, in our case, emphasises s states as compared to p states. For binary systems Si/C and Si/N, the experimental XPS and XES–XAS spectral curves are qualitatively well reproduced by the DOS calculations performed by J. Robertson in the tight-binding formalism for crystalline βSiC (Robertson 1992) and α-Si3N4 crystalline phases (Robertson 1991). Differences in relative intensities of the peaks are due to differences in photoionisation cross-sections (Driss-Khodja et al. 1996). These conclusions are in total agreement with those obtained for the same sample from the XPS core levels and EXAFS studies (Gheorghiu et al. 1992). In Si/C, our results confirm unambiguously the existence in the powder material of chemically ordered sp3 bonded network with Si—C heteropolar bonds. The Si 3p states whose position corresponds to Si Kβ emission are mainly located at the top of the VB, in coincidence with peak A2 of XPS. The sharp XPS peak A4 at 9.5 eV, at the same energy as a low feature of Si Kβ, can be associated to sp states. From Robertson’s calculations, we have deduced that this peak is due to mixed Si 3s and C 2p states; from our results Si 3p states are also involved in this energy range. A sharp minimum, corresponding to the ionicity gap, separates the maximum A4 from peak B; this energy gap is characteristic of Si—C bonds in a crystalline network. The minimum is expected to be partially filled in the case of amorphous SiC, as predicted from a theoretical study of α-SiC density of states performed by Finocchi et al. (1993). Peak B is largely C 2s in character, with a small, but non-negligible Si 3p contribution as clearly revealed by our XES data, in excellent agreement with theoretical predictions. As concerned the conduction band, an overall agreement is found between the photoabsorption curve and the theoretical DOS, for which the absorption edge is followed by two peaks attributed to antibonding Si 3p–C 2s and Si 3p–C 2p states distant by about 3 eV ; the non-resolution of these features in the experimental curve is probably due to the core-level broadening. Beyond the first peak, the experimental photoabsorption curve shows well-resolved peaks that reveal the existence of a wellorganised network. This is in agreement with the EXAFS results (Gheorghiu et al. 1992) we have obtained for this powder, which clearly revealed the existence of Si—C4 groups. The influence of hydrogen incorporation in SiC matrix has been recently studied in a theoretical simulation by Finocchi and Gattia (1994). For a-SiC : H alloys at stoichiometric

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Page 164 composition, containing 20% hydrogen, these authors predict that hydrogen atoms are bonded to both C and Si atoms and that the presence of H favours the formation of a tetrahedral network and thus enhances the chemical order in the material. This is well supported by our results. In Si/N powder, the VB distribution observed experimentally is very close to the Si3N4 one. Our results confirm that the top of the VB corresponds to an N 2pπ lone pair orbital, in agreement with theoretical results (Robertson 1991). The lone pair orbital is characteristic of the planar configuration of N atoms surrounded by three Si atoms. This feature is followed by mixed Si 3p–N 2p and Si 3s–N 2p states; N 2s states, mixed to Si 3s–3p states appear as a well-resolved peak. A broadening of the Si 3p distribution, when going from Si3N4 (Sénémaud et al. 1993) to Si/N powders is noted. Two effects can contribute to this broadening. The first effect is due to a loss of long-range order as compared to well-crystallised silicon nitride; a topological disorder attributed to fluctuation of bond lengths and bond angles has been evidenced in this material from a previous study (Gheorghiu et al. 1992). Let us note that this effect does not modify the overall shape of the VB distribution, which is essentially determined by the short-range order (Kärcher et al. 1984). The second effect is the presence of like-atoms bonds Si Si in a low, but not negligible proportion, as already revealed by the Si 2p core level study. Thus Si/N powders are characterised by both structural and chemical disorder. The presence of hydrogen in the network is possible, according to the preparation conditions. In a previous study of a-SiNX : H alloys, we had shown that hydrogen atoms are preferentially bonded to N atoms (Driss-Khodja et al. 1989). In this case, the Si states distribution should not be significantly modified. From the results obtained by Ley for hydrogenated a-Si—NX samples (Kärcher et al. 1984), hydrogen incorporation involves a recession of the VB maximum, which remains probably small for the low H concentration corresponding to our samples. Moreover, due to the low value of H 1s cross-section under our experimental conditions (hν = 1253 6 eV), the effect of hydrogen on XPSVB spectrum remains probably quite negligible. The first photoabsorption peak gives the distribution of the unoccupied Si p states; its shape corresponds to a disordered network and supports the previous EXAFS results on the same samples (Gheorghiu et al. 1992). In this energy range, the DOS calculations for silicon nitride (Robertson 1991), which covers an energy range of about 12 eV from EF, predicts the existence of a broad maximum between 5 and 10 eV from EF; the shift between the experimental and the calculated data that we have already mentioned (Gheorghiu et al. 1992) can be explained by the experimental gap value introduced in the calculation. It is interesting to note that the presence of hydrogen in the network favours a higher chemical and structural order in the case of Si/C systems; on the contrary, it involves a tendency for more disorder in the case of Si/N powder; in both cases the laser power used for the synthesis is 600 W (Table 2.2). In Si/C/N powders, the valence and conduction band distributions are quite different according to the composition. For C/N value equal to 0.93 (SiCN12), our results indicate clearly that Si atoms are involved in both carbide and nitride configurations. This is evidenced from XPSVB spectrum which reveals that both C 2s and N 2s states, mixed to Si sp states, are present; on the XESSi Kβ spectrum, we observe a small peak at the energy position of B (Si 3p–C 2s states). Moreover, these results are well supported by the photoabsorption spectrum of SiCN12, which shows two features close to the edge in coincidence with the first absorption maximum of SiN7 and SiC151, respectively, next features being close to those observed in SiC151 spectrum.

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Page 165 In SiCN27 with C/N = 0.22, the spectral distributions are rather close to SiN7 ones, but slight differences are noted which clearly indicate that the network is not identical to silicon nitride: (i) the XESSi Kβ spectrum (Figure 5.3.9) shows that the Si 3p maximum is slightly shifted towards low BE as compared to SiN7 Si 3p, without broadening; this result shows that additional Si—C bonds are present in the network, a few Si—N bonds being replaced by Si—C bonds; (ii) the minimum between A5 and B on the XPSVB spectrum is slightly enhanced as compared to SiN7 one. The energy position of this minimum corresponds to that of peak B of Si151C spectrum and consequently this additional shoulder reveals that C 2s states due to a few Si—C bonds are present in the network. On the photoabsorption curve, the sharp maximum following the photoabsorption edge corresponds to the SiN7 first maximum; it is characteristic of a crystalline Si3N4 network. A maximum is located at about 20 eV, close to both SiN7 and SiC151 maxima. The presence of a shoulder at about 7 eV from EF could result from Si —C bonds. Consequently, for sample SiCN27, our experimental results are consistent with the existence of a single type of chemical environment of Si atoms with both Si—N and a few Si—C bonds in the network. This is in quite good agreement with our previous results obtained from XPS core levels and EXAFS studies (Gheorghiu et al. 1992). Therefore, our results reveal that in ternary systems the local environment of Si atoms depends essentially on the preparation conditions of the powder. For C/N ratio close to 1, Si atoms are involved in both Si—C4 and Si —N4 configurations. EXAFS data for these samples (Gheorghiu et al. 1992) indicated that atomic distances Si —C and Si—N are close to the corresponding crystalline values. Moreover, they clearly revealed the presence of β-SiC microcrystallites in the material; we suggested that the microcrystals could be contained in an amorphous phase of SiN4 units, H atoms being present in NHX, CHX or SiHX groups. For an initial ratio C/N equal to 0.22, we confirm a local atomic arrangement with both C and N atoms surrounding the same Si atom, in agreement with XPS core levels and EXAFS results (Gheorghiu et al. 1992). 5.3.4 Effect of heat treatment on the electronic structure of nanometric Si/C/N ex-HMDS and pre-alloyed Si/C/N/Al/Y powders studied by XPS spectroscopy 5.3.4.1 HMDS nanopowders The samples were prepared by laser pyrolysis of hexamethyldisilazane. The experimental method, already described (Cauchetier et al. 1994; Martinengo et al. 1996), is based on the resonance between the emission of a continuous-wave CO2 laser, at 10.6 μm, and the IR absorption of the liquid precursor, in the form of an aerosol produced by an ultrasonic generator. The Si/C/N ex-HMDS powders have an average grain size in the nanometric range (20–40 nm). Their composition, determined by chemical analysis, is adjusted by varying the partial pressure of the carrier gas (NH3). Two series of powders in the intermediate range were studied, namely HMDS43 and HMDS44, which correspond to chemical composition C/N equal to 0.67 and 0.58, respectively, for AF samples. The XPS measurements were performed on AF powders and after heat treatment at Ta = 1500°C and 1600°C during 4 h under N2. The AF powders were stored under inert atmosphere. The analysed samples were pressed pellets prepared from pure laser synthetised powders.

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Page 166 Core levels The Si 2p, C 1s and N 1s core levels spectra from HMDS43 and HMDS44 powders are reported in Figures 5.3.10–5.3.12 for AF samples and for powders after heat treatment at 1500°C and 1600°C for 4 h. The Si 2p spectra (Figures 5.3.10(a) and (b)) are broad and asymmetric peaks which reveal the presence of several contributions in each spectrum. A decomposition of the spectra into Voigt functions is achieved, by assuming that each contribution corresponds to stoichiometric species: pure Si (99.5 eV), SiC (101.1 eV), silicon nitride Si3N4 (101.8 eV), or silicon dioxide SiO2 (103.2 eV) and no intermediate bonds such as Si CXN4 −X are considered (Sénémaud et al. 1998). For AF samples (Figures 5.3.10(a) and (b)), the maximum of the peak is close to the BE of Si3N4. The overall shape of the spectrum is significantly wider in HMDS43 than in HMDS44, for which a marked asymmetry towards low BE is observed. The Si 2p peaks can be fitted with a sum of two Voigt functions, the energy separation of which is close to the energy distance SiC—Si3N4 and, consequently, are attributed to a sum of silicon carbide and silicon nitride contributions, respectively. In agreement with the results of chemical analysis of the samples, the relative proportion Si3N4/SiC (noted G) is much higher in HMDS44 than in HMDS43; G is respectively 2.2 and 4.2 for HMDS43 and HMDS44. The two components are broad (about 2 eV) and their energy position is not well defined (within ± 0.5 eV). This result reveals that the AF powders exhibit a high degree of chemical disorder. The existence of mixed tetrahedra with C and N atoms around a same Si atom, which would correspond to peaks located at intermediate energy positions, is likely (Gheorghiu et al. 1997, 1998b). Hydrogen is present in AF powders (Musset et al. 1994) and the presence of H bonds in the network can contribute to increase the chemical disorder and consequently the width of the lines. After a heat treatment of powders at 1500°C, the Si 2p line from both samples exhibits a high BE tail which can be attributed to the presence of silicon dioxide. Besides this SiO2

Figure 5.3.10 Si 2p core levels for AF samples and for 1500 °C and 1600°C annealed samples (a) HMDS43 and (b) HMDS44.

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Page 167 contribution, a shift of the maximum of the line towards higher BE is observed in HMDS44 as compared to HMDS43. The decomposition of the lines into peaks from silicon carbide and silicon nitride shows that the apparent shift is due to the variation of the relative proportion of the Si3N4 to SiC contributions (G ratio) which is respectively 1.1 and 2.1. In HMDS43, the relative proportion of nitride and carbide is approximately the same whereas in HMDS44, Si—N bonds are dominant. Moreover, the widths of the components are significantly lower in HMDS43 than in HMDS44 (about 1.6–1.7 instead of 1.9 eV). In 1600°C annealed samples, the presence of SiO2 is still noted with about the same proportion as in 1500°C annealed sample. It is noteworthy that the factor G decreases considerably from AF samples to 1500°C and 1600°C samples: G is respectively 0.4 and 0.8 for 1600°C annealed HMDS43 and HMDS44 samples. The contribution from SiC is dominant for both samples; it is higher for HMDS43 than for HMDS44; once again, this result is in good agreement with the chemical analysis results. The energy positions of SiC and Si3N4 contributions are obtained, in this case, close to the reference binding energy values; their energy widths are significantly narrower than in AF samples. It is noteworthy that Si 2p lines from SiC are only 1.6 and 1.5 eV wide respectively in 1600°C annealed HMDS43 and HMDS44 samples; these values correspond to a perfectly well-ordered SiC network. The nitride contribution is narrower in HMDS43 (1.6 eV) than in HMDS44 (1.8 eV); in both cases the result reveals a well organised network. C 1s spectra observed from HMDS43 and HMDS44 samples (Figures 5.3.11(a) and (b)) exhibit quite different shapes according to the composition and to the heat treatment. In HMDS43 AF sample, C 1s consists of a quasi-symmetric broad line; in HMDS44, the peak is split into two components. The decomposition into Voigt functions reveals in both cases the existence of two peaks. In HMDS43 and HMDS44 the high BE component corresponds to C—N (285.9 eV) (Dufour et al. 1994), possibly mixed to C—H (285.0 eV) bonds; the low BE component is clearly resolved from the high BE one in HMDS44 and thus it can be associated

Figure 5.3.11 C 1s core levels for AF samples and for 1500°C and 1600°C annealed samples (a) HMDS43 and (b) HMDS44.

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Page 168 to C—Si (283.2 eV) (Driss-Khodja et al. 1992). In HMDS43, the low BE component is closer to the high BE component and reveals more probably the presence of C—C bonds (sp2). After a 1500°C annealing of sample HMDS43, the C 1s line is strikingly sharp and the energy maximum of the narrow peak corresponds to C—C bonds (284.3 eV) (Diaz et al. 1996). The small width of this peak (only 1.2 eV) reveals unambiguously the presence of small crystallites. This sharp peak is accompanied by a low BE shoulder corresponding to C—Si bonds in small proportion and by a tailing towards high BE which reveals the presence of C—N bonds (or C—H). A 1500°C annealing of HMDS44 sample gives quite a different shape for C 1s spectra. The line is well fitted by both C—C and C—N contributions; a low proportion of C—Si bonds is present as revealed by the low BE component. The high BE tailing of the spectrum can be interpreted by the existence of a few C—O species at the surface of the sample. After a 1600°C annealing, the most intense line is attributed to C—Si (carbide) for both HMDS43 and HMDS44 samples. Towards higher BE, the tailing is interpreted by the presence of C—N bonds. The presence of C—C bonds, the energy position of which is intermediate between those of C—N and C—Si contributions, is not excluded. N 1s spectra from HMDS43 and HMDS44 samples are reported in Figures 5.3.12(a) and (b). In AF HMDS43 powder, a broad line is observed which can be fitted with a sum of two Voigt functions having approximately the same intensity. The low BE peak can be associated to N—Si3 group (398.0 eV) (Gheorghiu et al. 1992); the higher energy component can be associated to N—H and N—C bonds; these two contributions are only separated by 0.5 eV and located at respectively 400.1 and 399.5 eV (Bischoff et al. 1987; Rossi et al. 1994). In fact, AF powders contain a non-negligible proportion of H atoms (Musset et al. 1994) which are essentially present in N—Hn groups. For HMDS44, the width of the observed line is noticeably lower than in HMDS43. A tailing of the line on the low BE side is noted. The decomposition reveals the existence of a strong peak attributed to N—H and N—C bonds and a lower intensity peak associated to N—Si bonds. Consequently, the significant decrease of the total width of the line is due to an appreciable decrease of the low BE contribution.

Figure 5.3.12 N 1s core levels for AF samples and for 1500 °C and 1600°C annealed samples (a) HMDS43 and (b) HMDS44.

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Page 169 After a 1500°C annealing of the sample, most H atoms are eliminated and, consequently, the proportion of N —H bonds is much lower. In this case, the N 1s peak is asymmetric and reveals the existence, in both samples, of two components that we attribute to N—Si and N—C bonds, respectively. In HMDS43, the low BE (N—Si) peak is more intense than the high BE (N—C) one. On the contrary, in HMDS44, the high BE component dominates. After a 1600°C annealing, both HMDS43 and HMDS44 spectra exhibit contributions at approximately the same energies. In HMDS43, the relative intensities are about the same as at 1500°C. In HMDS44, the relative intensities are completely modified, as compared to 1500°C, and the N—Si contribution becomes the most important. The splitting of the peak is well observed in 1600°C annealed samples. This effect is principally due to a narrowing of N—Si component in both samples. XPS valence band spectra The VB distributions from HMDS43 and HMDS44 powders in AF samples and in samples annealed at 1500°C and 1600°C are reported in Figures 5.3.13(a) and (b). The VB of the AF samples consists, for both spectra, of a wide band A in the range 0–14 eV; the maximum of this band is located at about 12 eV BE. This first band is followed by a deep minimum around 15 eV; an intense peak D is observed at 20 eV, and a lower structure E at about 26 eV. For 1500°C annealed samples, the overall shape of the spectral curves remains approximately the same and only small modifications are observed as compared to AF powders. The maximum A is not significantly shifted but it becomes slightly broader, especially in HMDS43. The minimum between A and D is broader and slightly filled up; in HMDS43, the sharp minimum is replaced by a 1 eV wide step. Peak D is noticeably narrower; the energy position of features D and E are not modified. The minimum between D and E is more pronounced. These modifications are approximately the same in both samples. Striking changes of the VB distributions are observed after the annealing at 1600°C, in both HMDS43 and HMDS44 samples. The maximum of peak A is shifted by 2 eV towards

Figure 5.3.13 XPSVB spectra for AF, 1500°C and 1600°C annealed samples (a) HMDS43 and (b) HMDS44.

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Page 170 lower BE for both samples; it becomes sharper and a shoulder at about 7.5 eV BE is clearly observed. The most striking modification, as going from Ta = 1500–1600°C, consists of a dramatic decrease of peak D intensity and the appearance of an additional peak labelled B at about 15 eV. It is noteworthy that peak B is well resolved from A and from D, and that wide minima occur between A, B and D. The intensity of B is slightly higher than that of D for both HMDS43 and HMDS44 samples. As far as peak E is concerned, its intensity is strongly enhanced for 1600°C annealed samples as compared to 1500°C annealed ones. The modifications of the spectral distributions observed as Ta increases from 1500°Cto 1600°C can be discussed by reference to VB distributions of the stoichiometric compounds SiC (Driss-Khodja et al. 1992) and Si3N4 (Sénémaud et al. 1993), which are reported in Figures 5.3.14 and 5.3.15. The energy positions of features A, B and D are given in Table 5.3.2 for HMDS43 and HMDS44 powders (AF, Ta = 1500°C and 1600° C), comparatively to stoichiometric compounds SiC and Si3N4. Let us recall that in SiC, peak A corresponds to C 2p states mixed to Si 3s and Si 3p states and peak B to C 2s states mixed to Si 3s and Si 3p states (DrissKhodja et al. 1992) In Si3N4, peak A is attributed to N 2p states mixed to Si 3s and Si 3p states, the lower BE feature being attributed to the N 2pπ lone pair; peak D is associated to N 2s states mixed to Si 3s and Si 3p states (Sénémaud et al. 1993). For HMDS43 and HMDS44 nanopowders annealed at 1600°C, the shape of band A is intermediate between those of SiC and Si3N4: the maximum of the peak is located between the maximum of A in SiC and in Si3N4. Both peaks B and C are observed for the samples annealed at 1600°C; they correspond to Si 3s, 3p states mixed to C 2s states (B) and respectively N 2s states (D); they reveal unambiguously the existence of both Si—C4 and Si—N4 groups in these systems. The results obtained from the analysis of the N 1s core level have shown (see Section 4.1.1) that in 1500°C annealed samples, N atoms are bonded to both Si (N—Si bonds) and to C atoms (N—C bonds). The contributions of these bonds in the VB are located in the energy range corresponding to peak D (about 20 eV BE). For these samples annealed at 1500°C, the study of C 1s core level has shown that C atoms are essentially bonded to C atoms (C—C bonds) or to N atoms (C—N bonds); the contribution of C—Si bonds is very low (even negligible), as revealed by the very low intensity in the energy range concerned. The presence of C—C bonds (either sp2 or sp3) induces VB states around 17 eV BE, accompanied by a sharp peak at 11.5 eV in diamond (Mac Feely et al. 1974). These states are located well apart from

Figure 5.3.14 XPSVB spectrum for stoichiometric SiC.

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Figure 5.3.15 XPSVB spectrum for stoichiometric Si3N4. Table 5.3.2 BE referred to the Fermi level EF for HMDS43 and HMDS44 HMDS43 and HMDS44 A(eV) B(eV) D(eV) AF 12.0 — 20.0 Ta = 1500°C 12.0 — 20.0 Ta = 1600°C 10.0 14.5 20.0 SiC (5.7) 9.5 14.4 — Si3N4 (3.9 7.6) 11.4 — 18.9 C—Si(Si—C4) contribution which occurs at 14.4 eV. As a consequence, the contribution of C—C bonds induces in the VB a broadening of peak D and a filling up of the minimum between peak D and peak A; no separate peak is expected. Consequently, the great intensity of peak D, in AF and 1500°C annealed samples, is consistent with an associate contribution of N—C and C—C bonds in the VB. The slight filling up of the VB distribution in the energy range of peak B reveals the low contribution of C—Si bonds to the VB. After annealing at 1600°C, there is a complete reorganisation of bonds around C and N atoms which are involved mainly in Si—C4 and Si—N4 groups. It is to be noted that some effects may be due to rehybridisation of orbitals (e.g. C sp2–C sp3). The data obtained for the VB distribution, as a function of heat treatment, are in good agreement with the results observed for the Si 2p, N 1s and C 1s core levels. In particular, after a 1500°C annealing, the presence of the strong peak at about 20 eV BE observed in the VB spectrum can be interpreted by the fact that N atoms are involved in both N—Si and N—C bonds; the existence of both types of bonds involving a broadening of the corresponding N 2s peak of the VB around BE = 20 eV. In these samples, C atoms are essentially involved in C—N and C—C bonds, as revealed by the analysis of core levels spectra, and this can explain why the VB distribution exhibits no contribution in the BE range corresponding to C 2s states from C— Si bonds.

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Page 172 Our results are in quite good agreement with a structural model composed of SiN3C tetrahedra linked through C—N bonds obtained from EXAFS measurements on the AF SiCN-HMDS samples (Ténégal et al. 1996). 5.3.4.2 Si/C/N/Al/Y nanopowders The samples were prepared by laser pyrolysis of a liquid mixture of hexamethyldisilazane (HMDS) and dissolved solid Al,Y-isopropoxides. This method is already described in Section 5.3.4.1 (HMDS powders). Amorphous ultrafine (20–40 nm) nanopowders are produced by laser pyrolysis of a liquid precursor. As for the SiCN-HMDS powders, the Si/C/N/Al/Y nanopowders may have variable N contents depending on the proportion of NH3 in the carrier gas (which is a mixture of argon and ammoniac). The powders were annealed at 1600°C for 1 h under N2 atmosphere. The atomic compositions of the studied powders (amorphous and annealed) as well as their designation are reported Table 5.3.3. The Al + Y content does not exceed 1 at% in the AF powders. Evolution of the Al and Y content in the superficial layers with annealing treatment CORE LEVELS In order to get information about atomic concentrations in the superficial layers of the powder grains and about their evolutions with annealing treatment, the intensities of the photoemission peaks (core levels) were compared. The core levels were normalised with respect to the number of measured scans. Figure 5.3.16 gives the Al 2p and Si 2p core levels for the intermediate composition (C/N = 0.78) for the AF and annealed sample and one can observe a strong increase of the Al 2p/Si 2p intensities ratio with annealing. From the measurements of the peak intensities and from the values given by Scofield (Scofield 1976) for the photoionisation cross-sections, we deduced the Al/Si atomic ratio in the superficial layers:

(5.3.1) where and are the intensities of the Al 2p and Si 2p photoemission peaks, and the photoionisation cross-sections given by Scofield (Scofield 1976). Table 5.3.4 reports the results of the intermediate composition. Al/Si atomic bulk (global) ratios are also given. The comparison between bulk and superficial ratio (from superficial layers) underline a strong increase of the Al/Si superficial ratio with annealing, whereas the bulk ratio tends Table 5.3.3 Atomic composition of the studied Si/C/N/Al/Y/O nanopowders Powders Si N C O C/N Al Y (at%) (at%) (at%) (at%) (at%) (at%) HSAlY17, AF 33.5 33.5 26 6.2 0.78 0.57 0.14 HSAlY16, 1600°C/1 h/N2 41.3 52.1 4.9 1.5 0.09 0.14 0.02 HSAlY17, 1600°C/1 h/N2 44.4 29.5 24.1 1.6 0.82 0.34 0.11 HSAlY18, 1600°C/1 h/N2 47.3 8.3 41.9 1.9 5.05 0.41 0.18

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Figure 5.3.16 Al 2p and Si 2p core levels for a powder with an intermediate C/N ratio (HSAlY17). (a) For the AF sample (C/N = 0.78) and (b) for the annealed one (1600°C, C/N = 0.82). Table 5.3.4 Surface ratios deduced from XPS analysis compared to bulk ratios from atomic global compositions for HSAlY17 IAI2p/ISi2p AP/Si surface ratio (XPS) AP/Si bulk ratio AF 0.004 0.01 0.02 1600°C/1 h/N2 0.022 0.03 0.01 to decrease. This result shows that Al atoms move towards the surface of the grains, when annealing is performed, and chemical shifts of the Al 2p and Si 2p peak by 1 and 0.5 eV (Figure 5.3.16) reveal that the chemical environment of Si and Al atoms changes as they reach the surface.

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Page 174 As we want to evaluate in which proportions aluminium content increases in the superficial area of the grains, the atomic ratio of Al (1600°C/AF) is deduced from the measurements of the Al 2p peak intensities for the AF sample and for the annealed one (1600°C). As we compare ratio of the same atomic species (Al or Si) in two different samples, then equation (5.3.1) is no longer valid and differences of densities, textures (amorphous or crystalline powders) which modify λ (the elastic mean free path for the measured electron kinetic energies), and of the size of the grains, must be taken into account in the evaluation of the atomic surface ratio. For the crystalline sample (1600°C), the density is about 1.5 times more than the one found in the AF sample (amorphous powder) and then the elastic mean free path is reduced in the crystalline sample. In a first approximation, the number of atoms probed by XPS for the two samples is almost the same: in the annealed sample, the distances covered by the electrons are shorter than in the AF sample but the number of atoms per unit of volume is increased. In the crystalline sample (1600°C), the increase of the average size of the grains (by 30–40%) induces a reduction of the atomic fraction seen by XPS: sensitivity to superficial layers is increased for the annealed sample. Table 5.3.5 reports the intensities ratio for Al 2p and Si 2p core levels as well as the atomic ratios given for the bulk. Collectively, Si content increases slightly (ratio of 0.75) when annealing is performed whereas the Al content (ratio of 1.68) decreases. The analysis of superficial ratios (relative to superficial layers) peak intensities confirms a strong increase of the Al content in the superficial layers after annealing: for the annealed powder (Figure 5.3.16(a), Table 5.3.6), the intensity for the Al 2p peak is five times greater than that found in the AF sample. The strong difference between atomic compositions and measured Table 5.3.5 Ratios of the atomic concentration between the AF sample and the annealed one (1600°C) for Al and Si. Measurements were performed in the same experimental conditions for both samples (same preparation, same detection angle). Samples HSAl16-17-18 IAF/I1600°C AF/1600°C surface (XPS) AF/1600°C bulk A1 2p 0.17 0.17 1.68 Si 2p 1 1 0.75 Table 5.3.6 Results of the simulations with Voigt functions of the Al 2p and Y 3d core levels Powders Al 2p Y 3d AlO4 + AlO6 AlN4 3/2 5/2 E(eV) Γ(eV) E(eV) Γ(eV) E(eV) Γ(eV) E(eV) Γ(eV) HSAlY17, AF 75.7 2.6 — — — — — — HSALY16, 1600°C/1 h/N2 75 1.8 73.8 1.8 160.9 1.7 159.0 1.8 HSALY17, 1600°C/1 h/N2 74.9 1.9 73.8 1.6 160.9 2.0 159.0 1.8 HSALY18, 1600°C/1 h/N2 74.8 1.5 73.8 1.5 160.7 1.9 158.6 1.8 References: YSiAlON 74.9 1.8 73.8 1.7 160.7 2.1 158.6 2.2 Sintered sample 74.9 2.2 73.8 2.1 160.0 1.8 158.2 1.7

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Page 175

Figure 5.3.17 Y 3d core levels (3/2 and 5/2) for the intermediate powder (HSAlY16, HSAlY17, HSAlY18) (AF and the annealed sample).

Figure 5.3.18 Al 2p core levels for the annealed samples (1600°C (HSAlY16, HSAlY17, HSAlY18) and results of the simulations with Voigt functions. values for Al 2p confirms the increase of Al concentration at the surface of the grain. The same result is obtained for yttrium. Measurements performed on the as-prepared sample show that Y 3d core levels are not detected (Figure 5.3.17). On the other hand, XPS spectra display a clear Y 3d signal for the annealed sample (Figure 5.3.17), whereas yttrium global content does not change. As the temperature increases until 1600°C, SiC and Si3N4 crystallites appear and Al and Y atoms are totally or partially rejected at the surface of the crystallites.

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Page 176

Figure 5.3.19 Al 2p core levels for the references compounds (sintered sample and YsiAlON) as well as the reconstructions. Figure 5.3.16 shows that the Si 2p and Al 2p XPS peaks move towards lower BE when annealing is performed, indicating that Al and Si atoms are bonded to less electronegative atoms at 1600°C than in the AF powder. The Al 2p peak is located at 75.8 eV in the AF sample and its large width at half maximum (2.6 eV) indicates that Al atoms are embedded in a disordered network. This energetic position indicates the presence of voids in the network. After annealing, the peak position is between 75 and 74 eV depending on the C/N atomic ratio of the sample (Figure 5.3.18). Figure 5.3.18 presents the results of the simulations performed on Al 2p core levels for the annealed samples. Three compositions were analysed: a rich C powder (C/N = 5.05), an intermediate one (C/N = 0.82) and a rich N powder (C/N = 0.09). The results obtained for the reference compounds (YSiAlON vitreous phase and a sintered sample) are also given (Figure 5.3.19). The solid sintered sample is a mixture of silicon nitride (Si3N4), SiC, β-SiAlON (crystalline phase) and YSiAlON. First of all, the width of the Al 2p peaks is quite similar in the nanopowders (1600°C) and in the reference compounds. The results of the simulations show that Al 2p peaks can be simulated using two Voigt functions, the energetic position of which is that found for aluminium in AlN and Al2O3 (Muilenberg 1979; Briggs and Seah 1996): AlN4 and AlO4 + AlO6 sites are detected in the annealed powders. The proportion of Al—N bonds increases with the C/N atomic ratio. Concerning reference compounds, one finds the same contributions to Al 2p core levels but the proportion of AlN4 sites is lower than that found in the powders. These results are in excellent agreement with those found by XAS analysis by Flank et al. (1999) on the same powders, (see part XAS): EXAFS results show that crystalline AlN is present in the powders at 1600°C. The XANES study revealed also that Al atoms are embedded in an amorphous phase in the as-prepared powders. The results of our simulations are given in Table 5.3.6. The FWHM of the contributions is quite similar for all the compounds. However, we observed the lowest width for the rich C powder, which is in agreement with a preferential formation of crystalline AlN in that powder.

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Page 177 Y 3d core levels (3/2 and 5/2) for the rich C powder are presented in Figure 5.3.20 (annealed powder) as well as those from the YSiAlON phase. The position and width of the peaks are similar for both compounds. The results of the simulations, given in Table 5.3.4 for the nanopowders and for the references, show that the chemical environment for yttrium atoms in the nanopowders is close to that found in the YSiAlON vitreous phase. The comparison of the simulations for the YSiAlON reference phase and for the sintered compound reveals a slight difference of the final parameters (E and Γ), due to differences in composition between the YSiAlON reference and the YSiAlON phase inside the sintered composite. The results of the decompositions (annealed samples) are given in Figure 5.3.21 for Si 2p core levels, in Figure 5.3.22 for N 1s core levels and in Figure 5.3.23 for C 1s ones. Si 2p core levels are modelised with four components relative to Si—Si4, Si—C4, Si—N4 and Si—O4 sites in crystalline compounds (Si, SiC, Si3N4 and SiO2). As the content of C increases, the proportion of crystalline SiC increases and the proportion of Si— C4 sites increases (Figure 5.3.21). A weak contribution from Si—Si bonds is observed for the rich N powder. The study of Si 2p core levels confirms the results of XRD analysis, which revealed the existence of Si3N4 and SiC crystalline phases with an increasing proportion of SiC when the C content of the powder increases. These results are in agreement with those of the previous structural studies performed on SiCN nanopowders (Sénémaud et al. 1998; Musset et al. 1994; El Kortobi et al. 1996; Ténégal et al. 1996, 1998) with different C/N ratios. The study of the N1s core levels (Figure 5.3.22) confirms the existence of N—Si bonds (near 398 eV), N—Al bonds (396.8 eV) and N—C bonds (399–399.5 eV) in the annealed powders. The main contribution to the N 1s core level comes from N—Si bonds (even for the rich C powder). The proportion of N—C bonds is more important in the powder with intermediate C/N ratio. Contribution at higher BE (near 401 eV) is attributed to N—O bonds. The existence of crystalline AlN at the surface of the grains is confirmed by analysis of the N 1s core levels.

Figure 5.3.20 Y 3d core levels for annealed HSAlY16, HSAlY17, HSAlY18 as well as the results of the simulations with Voigt functions.

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Page 178

Figure 5.3.21 Si 2p core levels for the annealed samples HSAlY16, HSAlY17, HSAlY18. One can observe the increase of the SiC content with the increase of the C/N ratio.

Figure 5.3.22 N 1s core levels for annealed HSAlY16, HSAlY17, HSAlY18. The simulations confirm the existence of an AlN crystalline phase at the surface of the grains. N—C bonds are also detected.

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Page 179 The study of C 1s core levels confirms the existence of C—N bonds (286 eV), while C—H bonds from surface contamination are also detected as well as contributions from C—Si bonds in silicon carbide. Contrary to the case of N 1s core levels, the N and C rich powders display very different shapes for C 1s core levels. Indeed, for the C rich powder, the main contribution to the core level comes from C—Si bonds in silicon carbide, the shoulder at 285–286 eV is expected to come from C—H and C—N bonds, but for the rich N powder, the intensity of the C—Si contribution strongly decreases and the main contribution to the C 1s core level is that from C—N bonds. These results show that C—N bonds are stable through annealing treatments. These results are in good agreement with the previously detailed structural study of Si/C/N nanopowders produced by laser pyrolysis of hexamethyldisilazane (Ténégal et al. 1996, 1998, see the XAS part) which underline the existence of homogeneous amorphous C substituted Si3N4-type structures (1–2 nm) for annealing temperatures below 1500°C. Mixed tetrahedra SiCXN4−X and C—N bonds were detected by XPS, RMN and XAS (Sénémaud et al. 1998; El Kortobi et al. 1996; Ténégal et al. 1996). When the annealing temperature increases beyond 1500°C, the N departure allows the local formation of SiC germs and powders crystallise in a mixture of Si3N4 and SiC, the proportions of which are dependent on the initial C/N ratio (in the AF powder). Our results show that when the C content decreases, the stable phase for C atoms is no longer a silicon carbide phase but a C substituted silicon nitride one. XPSVALENCE BAND SPECTRA VB spectra are given in Figure 5.3.24 for the three powders annealed at 1600°C. For the intermediate composition (HSAlY17), the corresponding spectra for powders without Al and Y is that of HMDS43 or HMDS44 given in Section 4.1 (Figures 5.3.13 and 5.3.14). The overall shape of the VB spectra for HSAlY17 (Figure 5.3.24) is similar to that obtained for HMDS43 and HMDS44 (Figures 5.3.13, 5.3.14 in Section 4.1.2). In SiC, peaks A and B are

Figure 5.3.23 C 1s core levels for annealed HSAlY16, HSAlY17, HSAlY18. C—N bonds are dominant in the rich N sample at 1600°C.

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Figure 5.3.24 VB distributions in the pre-alloyed HSAlY16, HSAlY17, HSAlY18 powders annealed at 1600°C. respectively attributed to mixtures of C 2p states with Si 3s–Si 3p states and C 2s states with Si 3s–Si 3p states. In silicon nitride, A and D are respectively attributed to mixtures of N 2p states with Si 3s–Si 3p states and N 2s with Si 3s–Si 3p states (see Section 4.1.2). The incorporation during the laser synthesis of Al and Y atoms does not change the shape of the VB. However, a slight decrease of the D peak (N 2s states mixed with Si 3s–Si 3p states) in the pre-alloyed powder (HSAlY17) is observed. The VB spectra for the rich N powder (HSAlY16) is quite different from that found for the silicon nitride reference (Figure 5.3.15 in Section 4.1.2). Modulations from 12 eV to the Fermi level observed on the silicon nitride VB spectra are not detected on the VB of HSAlY16. We expect to explain this difference by the presence of C and Al atoms in the network of silicon nitride which leads to a smoothing of the structures near the top of the VB. The presence of Al atoms in superficial layers either in AlN or SiAlON(C) crystalline phases at 1600°C can explain the VB smoothing from 12 eV to the Fermi edge. The VB distribution for HSAlY18 is close to that found in crystalline SiC (Figure 5.3.14, Section 4.1.2). References Bischoff, J. L., Kubler, L. and Bolmont, D. (1987) J. Non-Cryst. Solids, 97–98, 1407. Bouisset, E., Esteva, J. M., Karnatak, R. C., Connerade, J. P., Flank, A. M. and Lagarde, P. (1991) Phys. Rev., B24, 1609. Briggs, D. and Seah, M. P. (1996) in Practical Surface Analysis, 2nd edn, vol. 1; Auger and X-ray Photoelectron Spectroscopy, John Wiley & Sons. Cauchetier, M., Croix, O. and Luce, M. (1988) Adv. Ceram. Mater., 3, 548. Cauchetier, M., Croix, O., Herlin, N. and Luce, M. (1994) J. Am. Ceram. Soc., 77, 993. Cauchetier, M., Croix, O., Luce, M., Baraton, M. I., Merle, T. and Quintard, P. (1990) J. Eur. Ceram. Soc., 8m, 215. Diaz, J., Paolicelli, G., Ferrer, S., and Comin, F. (1996) Phys. Rev., B54, 8064. Driss-Khodja, M., Gheorghiu, A., Dufour, G., Roulet, H., Sénémaud, C. and Cauchetier, M. (1996) Phys. Rev., B53, 4287. Driss-Khodja, M., Dufour, G., Gheorghiu, A., Roulet, H., Sénémaud, C., Cauchetier, M., Croix, O. and Luce, M. (1992) Mater. Sci. Eng., B11, 97.

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Page 181 Driss-Khodja, M., Le Corre, A., Sénémaud, C., Gheorghiu, A., Thèye, M. L., Allain, B. and Perrin, J. (1989) J. Non-Cryst. Solids, 114, 489. Driss-Khodja, M. (1993) Thèse de l’Université Pierre et Marie Curie, Paris, 93PA066357; http://corail.sudoc. abes.fr80. Dufour, G., Poncey, C., Rochet, F., Roulet, H., Sacchi, M., De Santis, M. and De Crescenzi, M. (1994) Surf. Sci., 319, 251. El Kortobi, Y., Sfihi, H., Legrand, A. P., Musset, E., Herlin, N. and Cauchetier, M. (1996) Coll. Surf. A: Phys. Eng. Asp., 115, 319–327. Fang, R. C. and Ley, L. (1989) Phys. Rev., B40, 3818. Finocchi, F. and Gallia, G. (1994) Phys. Rev., B50, 7393. Finocchi, F., Gallia, G., Parrinello, M. and Bertoni, C. M. (1993) Physica, B85, 379. Flank, A. M., Armand, X., Cauchetier, M. and Mayne, M. (1999) J. Synchr. Rad. 6, 512. Gheorghiu, A., Dufour, G., Sénémaud, C., Herlin, N., Musset, E., Cauchetier, M. and Armand, X. (1997) J. Phys. III, 7, 529. Gheorghiu, A., Sénémaud, C., Roulet, H., Dufour, G., Moréno, T., Bodeur, S., Reynaud, C., Cauchetier, M. and Luce, M. (1992) J. Appl. Phys., 71, 4118. Gheorghiu-de La Rocque, A., Dufour, G., Bonnefont, P. A., Sénémaud, C. and Cauchetier, M. (1998a) MRS Symp. Proc., 501, 115. Gheorghiu-de La Rocque, A., Dufour, G., Sénémaud, C., Cauchetier, M. and Legrand, A. P. (1998b) Proceedings of International Symposium of Carbon Science and Technology for New Carbons, Tokyo, p. 424. Himpsel, F. J., Mac Feely, F. R., Taleb-Ibrahimi, A., Yarmoff, J. A. and Hollinger, G. (1988) Physics and Chemistry of SiO2and Si/SiO2Interface (eds C. R. Helms and B. E. Deal), Plenum Press, New York. Hollinger, G., Sferco, S. J. and Lannoo, M. (1988) Phys. Rev., B37, 7149. Jentoft, C., Weinberg, G., Wild, U. and Schölgl, R. (1999) Precursor-derived Ceramics (eds J. Bill, F. Wakai and F. Aldinger), Wiley-VCH. p. 175. Kärcher, R., Ley, L. and Johnson, R. L. (1984) Phys. Rev., B30, 1896. Mac Feely, R. F., Kowalczyk, S. P., Ley, L., Cavell, R. G., Pollak, R. A. and Shirley, D. A. (1974) Phys. Rev., B9, 5268. Martinengo, H., Musset, E., Herlin, N., Armand, X., Luce, M., Cauchetier, M., Roulet, H., Gheorghiu, A., Dufour, G. and Sénémaud, C. (1996) Silicates Industriel, 1–2,9. Muilenberg, G. E. (1979) Handbook of X-ray Photoelectron Spectroscopy, Perkin Elmer. Musset, E., Cauchetier, M., Herlin, N., Luce, M., Gheorghiu, A., Dufour, G., Roulet, H. and Sénémaud, C. (1994) High Temp. Chem. Proc., 3, 535. Robertson, J. (1991) Philos., Mag., B63, 47. Robertson, J. (1992) Philos. Mag., B66, 615. Rochet, F., Rigo, S., Froment, M., D’Anterroches, C., Maillot, C., Roulet, H. and Dufour, G. (1986) Adv. Phys., 35, 237. Rossi, F., Andre, B., van Veen, A., Mijnarends, P. E., Schut, H., Labohm, F., Dunlop, H., Delplancke, M. P. and Hubbard, K. (1994) J. Mater. Res., 9, 2440. Scofield, J. H. (1976) J. Elect. Spectro. Rel. Phen., 8, 129–137. Sénémaud, C. and Ardelean, I. (1990) J. Phys. Condens. Matter., 2, 8741. Sénémaud, C., Driss-Khodja, M., Gheorghiu, A., Harel, S., Dufour, G. and Roulet, H. (1993) J. Appl. Phys., 74, 5042. Sénémaud, C., Gheorghiu de La Rocque, A., Dufour, G. and Herlin, N. (1998) J. Appl. Phys., 84, 4945. Sénémaud, C., Laporte, D., André, J. M., Khérouf, R., Paquier, P. and Ringuenet, M. (1989) Proc. ECO 2 Xray Instrumentation, (ed. R. Benattar), 1140, 297. Shirley, D. A. (1972) Phys. Rev., B5, 4709. Ténégal, François, Bouchet, B., Bellissent, R., Herlin, N., Cauchetier, M. and Dixmier, J. (1998) Philos. Mag., A78, 4, 803. Ténégal, F., Flank, A. M. and Herlin, N. (1996) Phys. Rev., B54, 12029. Zahorowski, W., Wiech, G., Mell, H. and Weiser, G. (1989) J. Phys. Condens. Matter, 1, 9571.

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Page 182 5.4 Local order in Si/C/N nanopowders studied by X-ray absorption spectroscopy Study at the Si—K edge François Ténégal, Anne-Marie Flank, Adriana Gheorghiu-de La Rocque and Christiane Sénémaud 5.4.1 X-ray absorption principles and data analysis 5.4.1.1 Principles X-ray absorption spectroscopy (XAS) refers to the measurement of the X-ray absorption coefficient μ as a function of photon energy E. μ is given by the Beer–Lambert law: where I0 and I are the intensities of the incident and transmitted beams, respectively and x is the sample thickness. μ is a measurement of the absorption cross-section σ which is given by the Fermi golden rule formula. As the energy of the incident X-ray beam reaches the binding energy of a core level, one can observe a strong increase of the absorption coefficient, corresponding to the excitation of an electron from its core level to bounded vacant states or to the conduction band. The edge position depends on the chemical environment of the absorbing atom. Transitions occur within the dipolar approximation for the absorbing atom. For a K-shell absorption (1 s electron), due to the dipole selection rules, only transitions to p-states are allowed (Δl = ± 1) and the X-ray absorption spectrum reflects the local partial unoccupied p-density of state (DOS) around absorbing atoms. As the energy gradually increases, if the absorbing atoms are isolated, the absorption coefficient decreases monotonically and its shape can be evaluated by a Victoreen function. If the atoms are embedded in a condensed phase (liquid or solid), then one observes modulations which have been shown to depend on the local environment around the absorbing atom (Figure 5.4.1). The photoelectron which has a variable kinetic energy (depending on the photon energy) is scattered by the neighbours of the excited atom, and builds up interference patterns that contain structural information. In the early 1970s, Lytle et al. (1975) showed that a subsequent Fourier transform (FT) of the extended X-ray absorption fine structure (EXAFS) oscillations yields a pseudo-radial distribution function. As the mean free path of the photoelectron depends on its kinetic energy (Figure 5.4.2), one will distinguish two regimes. In the EXAFS regime (Ec> 50 eV) (Figure 5.4.1), the interaction is in a first approximation calculated using the single scattering approach. The initial state for the electron is the localised core level corresponding to the absorption edge (K, LI, LII, LIII, …). The final state is that of the ejected electron, which can be represented as an outgoing spherical wave originating from the X-ray absorbing atom which is backscattered by the neighbouring atoms. The final state is the sum of the outgoing and incoming waves that gives

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Page 183

Figure 5.4.1 Absorption coefficient (background was subtracted) for β-SiC (cubic). When the energy of the Xray beam reaches the binding energy of a core level (E0 = 1841 eV), then an abrupt increase of μ is observed. As the energy increases (E > 1841 eV), modulations of the coefficient are clearly evidenced.

Figure 5.4.2 Mean free path for the electron as a function of its kinetic energy. rise to the sinusoidal variation of μ known as EXAFS. For an unoriented sample with small or approximately Gaussian disorder, the EXAFS oscillations are given by (Teo 1986)

(5.4.1)


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Page 184 k is given by

where E0 is the threshold energy, μ0(k) is the atomic absorption (for isolated atoms) and μ(k) the absorption coefficient. The summation is made over the concentric coordination shell j around the absorbing atom at a distance Rj. Nj is the number of atoms of the same kind inside the coordination shell j. Nj and Rj are structural parameters. The electronic parameters Fj(π, k, Rj) (the backscattering amplitude for the shell j) and j(k, Rj) + 2δc (the phase function) were calculated for each kind of atomic pairs and tabulated by Teo and Lee (1979), McKale et al. (1986) and Rehr and Albers (1990). They can be extracted from model compounds. is an amplitude reduction factor which accounts for the intrinsic inelastic loss (concerning the absorbing atom), exp(−2Rj/λj(k)) is a reduction factor accounting for extrinsic inelastic process is a reduction factor accounting for thermal and (concerning neighbouring atoms) and structural disorder, where σj is the Debye–Waller factor. The EXAFS formula (5.4.1) neglects multiple scattering events which were introduced in the EXAFS theory by Lee and Pendry (1975). Later on, Rehr and Albers (1990) gave a formulation of the EXAFS signal as a sum of the contributions from each possible scattering path Γ (single or multiple scattering path) for the photoelectron:

(5.4.2)

R is the effective path radius (the total length of the path Γ is 2R). σΓ is the DW factor for the Γ path, is its effective scattering amplitude and Γ (k) + 2δc(k) its effective phase shift. The second formulation of the EXAFS signal accounts for multiple scattering events which must sometimes be included in the EXAFS analysis (Figure 5.4.3). The FEFF code developed by Rehr and Albers calculates the EXAFS signal using the second expression (5.4.2) (Rehr and Albers 1990). The calculation requires an atomic coordinate file from which a muffin-tin potential is calculated. The paths are numbered and selected using amplitude

Figure 5.4.3 Representation of a single scattering process (1) for the photoelectron and a multiple scattering process (2) for the photoelectron.

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Page 185 filter paths importance criteria (Zabinsky et al. 1995). All the contributions Γ are individually calculated and gathered to give the EXAFS signal. Multiple scattering events become fundamental in the ‘low kinetic energy’ regime (Ec< 50 eV) where the mean free path of the photoelectron can be large. In that region, the photoabsorption spectrum is very sensitive to the medium-range order, as it is illustrated in Figure 5.4.4, where the XANES part of the XAS spectra for SiC in the crystalline beta form and in the amorphous phase are compared. The loss of mediumrange order in the amorphous material washes out the structures in the near edge area. As well, some features of the near-edge structures may appear as fingerprints of the crystalline phases present in the sample (Figure 5.4.5).

Figure 5.4.4 XANES part for β-SiC and amorphous silicon carbide (a-SiC) prepared by RF sputtering. XANES features are strongly smoothed in the case of a-SiC.

Figure 5.4.5 Comparison of Si—K edge XANES spectra for β-SiC, α-Si3N4, β-Si3N4.

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Page 186 5.4.1.2 Experimental An X-ray absorption experiment requires a white X-ray source of reasonable intensity to collect data. One fundamental reason for the development of the technique, in the 1970s, was the advent of the synchrotron radiation X-ray sources of high brilliance and intensity, which were improved since then. The photoabsorption measurements (HMDS powders) were carried out on the SA32 beamline (Super-ACO, LURE-Orsay) equipped with a double-crystal (InSb (111)) monochromator which allows an energy resolution of 0.7 eV. The incident beam was monitored by measuring the total electron drain current of an aluminium foil located downstream the monochromator. The photoabsorption EXAFS spectra at the Si—K edge were recorded at 298 K, were collected in the transmission mode as a function of the incident photon energy, in the 1800–2500 eV energy range and with a 1 eV step. The transmitted beam was measured with an ionisation chamber filled by a low pressure of air (200 Torr). Measurements for the powders produced from gaseous precursor were carried out on the SB3 beamline which was equipped with a two crystal InSb (111) monochromator. The photoabsorption spectra were recorded at room temperature, in the transmission mode as a function of the incident photon energy, in the 1750–2550 eV range with a 2 eV step. The transmitted beam was measured with a proportional counter filled with a 220 Torr Ar/CH4 (90/10) mixture. More accurate measurements were done in the near edge region (1830–1910 eV, 0.2 eV step), as the XANES part of the absorption spectra has been shown to be a good fingerprint of the degree of crystallinity and of the nature of the silicon phases present in the compound (Sainctavit et al. 1991). As a good approximation very near the edge, which is sufficient to draw qualitative inferences, a constant background was subtracted in the pre-edge region from the as-recorded spectra. They were then normalised far from the edge (at 1910 eV). As an example, Figure 5.4.5 presents the XANES spectra of β-SiC,α-Si3N4, and β-Si3N4. All the XANES spectra were recorded during the same shift, and references were checked before and after the measurements, as well as crystalline Si and crystalline and amorphous SiO2 already reported in the literature (Lagarde and Flank 1986; Lagarde et al. 1992, 1993). 5.4.1.3 EXAFS data analysis The first step is the extraction of the oscillatory part χ(k) from the absorption coefficient μ. The atomic absorption μ0 is modelised by a polynomial (4th or 5th order) fit of the modulations. The spectra are normalised using the Lengeler–Eisenberger (1980) approximation. The normalisation coefficient is given by the Heitler formula:

(5.4.3) where μ1(E0) − μ0(E0) is the experimental value of the jump at the threshold (E0). μ0(E) is the atomic absorption given by the polynomial fit. Once the oscillatory part is extracted, FT in R space of the k2 or k3 weighted oscillations χ(k) is performed. A Kaiser (τ = 2.5) or Hanning window is applied to the knχ(k) oscillation. FTs are calculated over the same energy range (2–10 Å−1) for all the samples and references compounds. Extraction of the oscillatory part and FT calculations must be executed using the same parameters for all the compounds that will be compared. FT calculations give according to the expression (I) a pseudo-radial distribution function (uncorrected from the phase shift) of the coordination shells surrounding the absorbing Si atoms (Figure 5.4.6). The sorted and back transformed

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Figure 5.4.6 Modulus of the FT of k2χ(k) for β-SiC(2–10 Å−1). The first peak corresponds to the first coordination shell around Si atoms (4 C atoms at 1.89 Å), the second to the second coordination shell (12 Si atoms at 3.07 Å). Peak A is a multiple scattering structure known as a fingerprint of the cfc structure. It is a focusing effect due to the alignment of 3 Si atoms. Phases shift are not corrected and the peaks do not appear at the real distances. in the k space peak, corresponding to a given coordination shell, is at first fitted using the electronic parameters (amplitude and phase functions) extracted from model compounds in which bonding configurations similar to those being investigated are present: β-SiC for Si—C pairs (4 at 1.89 Å) and Si—Si second neighbour pairs (12 at 3.08 Å), α-Si3N4 for Si—N pairs (4 atoms at an average distance of 1.74 Å), and α-SiO2 for Si—O pairs (4 at 1.61 Å) (Wyckoff 1963). The experimental signal extracted for the first shell is fitted using the single scattering formula (5.4.1) for EXAFS. For the first shell, given the available energy range of the EXAFS spectra Δk equal to about 8 Å−1, and the ΔR equal to 1.4 Å used in the back-FT, the number of free parameters (2 · ΔR · Δk/π) turns out to be seven. The free parameters during fitting process were the coordination number and the interatomic bond length. A three-shell fit deals with six free parameters. The description of the other contributions (beyond the closest neighbours) must include multiple scattering events and formula (5.4.2) must be used. 5.4.2 Study by EXAFS at the Si—K edge of Si, SiN, SiC and Si/C/N laser-synthetised nanopowders produced from gaseous precursors Six as-formed nanopowders (Si, SiC151, SiC152a, SiN7, SiCN12 and SiCN27) were measured. The Si nanopowder was formed from a SiH4—Ar gas mixture. XRD measurements revealed the existence of c-Si in the as-formed nanopowder. Two SiC (SiC151 and SiC152a) powders were synthetised from CH3SiH3 gaseous precursor. They were differentiated by the temperature reaction which induces changes in X-ray diffraction (XRD) patterns, as

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Page 188 observed previously in the case of the SiH4—C2H2+n mixtures (Luce 1990). The XRD measurements performed on SiC151 and SiC152a revealed the existence of β-SiC crystallites in the nanopowders. The SiN7 powder was formed from a SiH4—NH3 mixture and the two SiCN powders (SiCN12 and SiCN27) from SiH4— CH3NH2—NH3 mixture in variable proportions. Thus, NH3 was only used for the synthesis of SiCN27 and allowed the increase of the N content in SiCN27 (C/N at. = 0.22) compared to SiCN12 (C/N at. = 0.93). XRD measurements revealed the existence of α and β-Si3N4 crystallites in SiN7, SiCN12 and SiCN27 powders and of β-SiC crystallites in SiCN12. Note that a small proportion of c-Si was also detected in the SiN7 powder. 5.4.2.1 Si, SiN7, SiC151 and SiC152a nanopowders Figure 5.4.7 presents the FT of the Si nanopowder. The first peak (A) located near 2 Å corresponds to the first coordination shell surrounding the Si atoms. It is related to the four Si atoms which build the tetrahedral environment of the absorbing Si atom (RSi—Si = 2.35 Å). Beyond this first intense peak (R > 3 Å), one finds a set of small peaks, related to the intermediate-range organisation in the material, which are characteristic of that of crystalline Si (Lagarde and Flank 1986). These results are in good agreement with those found by XRD which revealed the existence of c-Si. The FT curves obtained for SiN7 and α-Si3N4 are displayed in Figure 5.4.8 and one can take qualitative information of their comparisons. First, one can observe a strong decrease in the amplitude of the peaks located at R > 2 Å for SiN7 compared to α-Si3N4. This reduction is an effect of the static disorder. A reduction in the amplitude of the first peak (A) is also observed in SiN7. An attempt was made to simulate the contribution of this peak on the basis of a one-shell model. The only parameters which were varied are the interatomic bond lengths and the coordination number. The DW factor was not allowed to vary during the fitting procedure, which means that we considered this factor equal to that found in the crystalline compound α-Si3N4 for the Si—N bond. A slight variation (


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Figure 5.4.8 FTs of k3χ(k)for α-Si3N4 and the SiN7 nanopowder. Table 5.4.1 First-shell structural parameters obtained for the different samples. ΔN/N = 10%, ΔR= 0.02 Å Sample Si—C Si—N NSi—C+Si—N NSi—C RSi—C(Å) NSi—N RSi—N(Å) SiN7 — — 3.1 1.73 3.1 SiC152a 3.4 1.88 — — 3.4 SiCN12 2 1.89 1.7 1.74 3.7 SiCN27 1 1.87 2.9 1.73 3.9 of the energy edge shift was allowed. The result of the simulation, given in Table 5.4.1, indicates that the coordination number is close to three which is lower than in α-Si3N4. The interatomic distance does not change compared to α-Si3N4. The reduction of the coordination number can be explained by Si—H bonds. Indeed, H with its single electron shows no measurable scattering and the main effect of its presence in the neighbourhood of Si atoms is a reduction in the effective measured coordination number. We expect that an effect of static disorder on the Si—N bond length is excluded. Indeed, the DW factor of the Si—N bonds in αSi3N4 is strong (the bond length distribution spreads from 1.70 to 1.79 Å) and no change of its value was observed in amorphous SiNx compounds in previous studies. The FT curves of SiC151 and SiC152a are displayed in Figure 5.4.9. SiC151 exhibits the well-known structures of the crystalline β-SiC phase (Laffon 1986). Particularly, the peak labelled D is characteristic of the fcc structure of the Si subnetwork in β-SiC. This is the ‘focusing effect’ peak which is related to the alignment of three Si atoms distant to each other of 3.08 Å in the β-SiC structure. SiC151 is fully crystallised in the β-SiC phase. The spectrum of SiC152a presents the structures of the β-SiC phase too, but a strong reduction of the amplitude of the peaks located at R > 2 Å is observed. The first shell contribution to the

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Figure 5.4.9 FTs of k3χ (k) for SiC151 and SiC152a. FT of SiC152a was simulated using the electronic parameters extracted from SiC151 (4 C at 1.89 Å). The result given in Table 5.4.1 is the decrease of the coordination number. One finds 3.4 C atoms around one Si atom. This result gives evidence of the presence of crystallite of β-SiC in SiC152a. The reduction of the coordination number may have several explanations. First, the H content is higher in the SiC152a sample, than in the SiC151 sample, and induces a decrease of the coordination numbers and so, the H effusion during the pyrolysis process is less important in SiC152a. We expect that, the existence of Si—H bonds in the as-formed powder is responsible for this decrease, since the gaseous precursor contains SiH3 groups. Second, the structural disorder is more important in SiC152a than in SiC151 due to bond length and/ or angular fluctuations. Smaller crystallites (1–2 nm) of silicon carbide in SiC152a would also lead to a decrease of the coordination number around Si atoms. For SiC151, the proportion and/or the size of the βSiC clusters are more important and the features of the EXAFS spectra are those of the crystalline β-SiC phase. 5.4.2.2 SiCN12 and SiCN27 FTs of SiCN12 and SiCN27 are given in Figure 5.4.10. The two compounds exhibit the features of ordered structures. Beyond the first peaks of the FT, one finds a set of peaks related to the intermediate-mediumrange organisation in the material. The presence of the focusing effect peak in SiCN12 indicates that, contrary to SiCN27 for which the peak is absent, this powder is crystallised partially in a β-SiC phase. This result is in agreement with the XRD measurements which detect α, β-Si3N4 phases for 1 and SiCN27 and the β-SiC phase only for SiCN12. The analysis of the other peaks is difficult because of the multiplicity of the possible configurations for Si atoms, in such SiCN compounds, and because a fine analysis of this part of the pseudo-radial distribution function requires the introduction of multiple scattering contributions. Concerning the first peak related to the first coordination shell, a two-shell fit (Si—N + Si—C) was realised using electronic parameters extracted from model compounds (SiC151 and α-Si3N4). The results are given in Table 5.4.1. One

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Figure 5.4.10 FTs of k3χ (k) for SiCN12 and SiCN27. Table 5.4.2 Influence of the flame temperature T during the pyrolysis process on the proportion of Si—H remaining bonds in the as-formed nanopowders assuming that the reduction of the coordination number is completely due to Si—H bonds (ΔN/N = ±10%) SiN7 SiC152a SiC151 SiCN12 SiCN27 T (°C) 1250 1150 1530 1550 1615 NSi−H 0.9 0.6 0 0.3 0.1 finds that the NSi—C/NSi—N ratio follows the C/N ratio. For SiCN12 (C/N = 0.93), the first coordination shell of one Si atom is composed of 2 C atoms and 1.7 N atoms, whereas for SiCN27 (C/N = 0.22), it is composed of 2.9 N atoms and 1 C atom. The coordination numbers are close to 4 for the two compounds which indicate that, as for SiC151, there is not much H bonded to Si. We expect that the proportion of Si—H bonds in the material is correlated to a pyrolysis reaction parameter which is the flame temperature. Indeed, for SiC151, SiCN12 and SiCN27, this temperature is about 1500–1600°C (Table 5.4.2) whereas for SiC152a and SiN7, in which a higher reduction of the coordination number (20–30%) in the first coordination shell is found, it is about 1100–1200°C. The results of the fits show that both C and N atoms are present in the first coordination shell of Si atoms in SiCN12 and SiCN27. For SiCN12, both SiC4 and SiN4 tetrahedra are present in the compounds, since XRD measurements show the existence of β-SiC and α, β-Si3N4 crystallites in the powder. For SiCN27, SiN4 tetrahedra are present in α, β-Si3N4 crystallites, but we cannot conclude about the existence of SiC4 or SiCxN4−x tetrahedra, since a two-shell fit of the EXAFS data provides an average picture of the local arrangement of C and N atoms in the first coordination shell of the absorbing Si atoms.

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Page 192 5.4.3 Short-range order determination in amorphous SiCN-HMDS nanopowders: evolution as a function of the C/N ratio and annealing treatment This part reports a detailed structural analysis of nanosized SiCN powders produced from the liquid precursor hexamethyldisilazane (HMDS). We focused our study on the intermediate compositions, for which a delay in the crystallisation is observed. XAS was used to describe the short- and medium-range order in the amorphous powders. The amorphous → crystal transition which occurs beyond 1500°C was also studied by Xray and neutron diffraction (see Section 5.4.4 ND and XRD). The goal of this study is to give a model for the structural evolution of these powders when the annealing temperature increases. The informations obtained from absorption and diffraction measurements are related to the structure of the powders at the atomic scale and are then particularly adapted to the study of amorphous compounds. The XAS experiments at the Si—K edge give a radial distribution function (RDF) around the Si atoms in the material, and diffraction measurements give a total radial distribution function. The analysis of these RDF allows the determination of structural parameters, such as bond lengths or coordination numbers, and of information related to the intermediate-medium-range order, which can be modelised and tested through XAS and diffraction spectra calculations. This detailed study of the amorphous phases was performed to obtain structural information, which could explain the delay in the crystallisation and give information about the mechanism that leads to the crystallisation of the nanopowders. 5.4.3.1 Mixed tetrahedra and C—N bonds Short-range order deduced from the analysis of the first coordination shell of Si atoms Figure 5.4.11 shows the extracted EXAFS signal as a function of the powder composition for the as-formed samples. The FT curves obtained for the various C/N ratio samples are

Figure 5.4.11 Si K EXAFS raw spectra for β-SiC, α-Si3N4 and Si/C/N with variable C/N. The spectra have been vertically displayed for convenience. The corresponding concentration of the samples are displayed (Table 2.5).

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Figure 5.4.12 (a) FT modulus of the k2 weighted EXAFS oscillations for the two references β-SiC, α-Si3N4, and for the Si/C/N samples; (b) imaginary part of the FT for the Si/C/N samples as a function of C/N. reported in Figures 5.4.12(a) and (b), and one can take out some qualitative information of their observation: absence of long-range order (for R values higher than 4 Å), short and intermediate orders very slightly dependent on the C/N ratio, as evidenced clearly by both the magnitude and the imaginary part of the FT, which are very similar up to 4 Å. In order to determine the local order around Si atoms, a two-shell fit has been calculated on the basis of Si—C and Si—N pairs. The structural parameters obtained as a function of the C/N ratio are summarised in Table 5.4.3.


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Page 194 Table 5.4.3 First-shell structural parameters obtained with a two-shell fit (Si—C + Si—N) for the different samples. ΔN/N = 10%, ΔR= 0.02 Å Sample Si—C Si—N NSi—C+Si—N NSi—C RSi—C(Å) NSi—N RSi—N(Å) HMDS40 1.4 1.88 2.1 1.73 3.5 HMDS43 0.9 1.9 3.2 1.73 4.1 HMDS44 0.9 1.87 3.2 1.72 4.1 HMDS45 0.7 1.88 3.3 1.73 4.0 We first used a fitting procedure where the two interatomic distances have been kept fixed to their crystalline values, within 0.02 Å. As we used experimental electronic parameters, the values of the adjustable disorder factors are just the difference between the unknown and the model compound where is included in the experimental amplitude. This difference, in all cases, is found to be negligible by the fitting procedure. One has to evaluate the DW factor for Si—C and Si—N bonds by fitting the first coordination shell in model compounds with the calculated backscattering amplitudes and the phase shifts taken from McKale et al. (1988). We found a DW factor value less than 5·10−2 Å, consistent with the value found in ZnSe (Diop and Grigenti 1995), and in c-Ge (Dalba et al. 1993). Then, during the fitting procedure, the damping factors Δσ2 were kept equal to 0, which means that we considered a DW factor for the unknown structure equal to those of the reference phases. As for the values of the energy edge shift ΔE0 = E0 alloy − E0 model, we checked that small variations of this parameter (±2 eV) leave the fit unaltered, therefore, this parameter was not varied during the fits. Then in a second step, the only parameters which were varied during the fit procedure were the interatomic distances and the partial coordination numbers. The different interatomic distances appear to be constant over the C/N range. The values obtained for the different pairs may be well compared with the model values. Such constancy has already been observed in hydrogenated amorphous silicon carbide alloys (Pascarelli et al. 1992). The total coordination number of Si is expected to be close to 4, as there is limited bonding of Si to H. Actually, hydrogen with its single electron shows no measurable scattering, and its structural effect can only be evidenced when it is located between the absorber and the backscatterer in well-defined sites. In disordered structure as the one described here, the main effect of the presence of H in the samples would result in a reduction of the effective coordination. The results concerning the total coordination number are 4 ± 10%, an acceptable uncertainty for EXAFS measurements. The proportion of the different pairs (partial coordination numbers) follows approximately the chemical composition, as it can be seen in Figure 5.4.13, which shows the evolution of the ratio of the partial coordination numbers NC/NN as a function of C/N. One can observe that this evolution is linear with a slope equal to 0.5. As the chemical composition of the powders relies on the presence of oxygen in the powders, fits including Si —O pairs were also performed. Table 5.4.4 reports the results obtained with a three-shell fit and the result of the adjustment for HMDS43 is given in Figure 5.4.14. The introduction of Si—O bonds improves definitely the quality factor of the fit (from 1.3 × 10−2 to 5.5 × 10−3). The proportion of Si—O pairs, that never exceeds 10%, increases with the O atomic content of the alloy.

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Figure 5.4.13 Variation of the ratio of the partial coordination numbers NC/NN with C/N: I is for parameters given in Table 5.4.3 and II for those in Table 5.4.4. Table 5.4.4 First-shell structural parameters obtained with a three-shell fit (Si—C + Si—N) for the different samples. ΔN/N = 10%, ΔR = 0.02 Å Sample Si—C Si—N Si—O NSi—C+Si—N+Si—O NSi—C RSi—C(Å) NSi—N RSi—N(Å) NSi—O RSi—O(Å) HMDS40 1.6 1.89 1.9 1.73 0.2 1.61 3.7 HMDS43 0.9 1.88 2.6 1.75 0.7 1.62 4.2 HMDS44 0.9 1.87 2.7 1.73 0.8 1.58 4.4 HMDS45 0.5 1.88 3.0 1.75 0.8 1.63 4.3

Figure 5.4.14 Best fit of the filtered first peak (0.6–2 Å) of the FT for HMDS43 (as-formed).

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Page 196 The estimation of the uncertainties in the fitted parameters is a non-trivial problem, because of the correlation between the different parameters, Rij and ΔE0 on one hand, and Nij and σij on the other. One should also consider the role played by the anharmonicity of pair potentials, which means the deviation from a gaussian behaviour. In another strongly covalent system, amorphous Ge, the study of the local disorder using the cumulant expansion leads to corrections, to interatomic distances and to coordination numbers (Dalba et al. 1993). These corrections are fully within the error bars we have quoted here. So we did not consider any anharmonic behaviour, which should be in any case very difficult to evidence in a two-shell disordered model. Furthermore, a non-negligible uncertainty of the first-shell fit procedure comes from the fact that C and N, acting as backscatterers, are neighbours in the periodic classification table, which means that they cannot easily be distinguished. Last, a two-shell model, with partial coordination numbers NC and NN, provides an average picture of the local arrangement of the C and N atoms in the first coordination shell of the absorbing Si atoms. It does not allow the discrimination of a locally demixed phase with SiC4 and SiN4 tetrahedra arrangement (complete chemical ordering), or a random stacking of mixed SiCxN4−x tetrahedra. This point will be clarified by looking at the near edge structure in the next section. The Si—K XANES studies In X-ray absorption, an electron absorbs an incident photon and is promoted to an empty state in the conduction band. Within the dipolar approximation, only transitions corresponding to Δl = ±1 between the orbital quantum numbers of the initial and final states are allowed. The experimental Si—K edge spectrum therefore reflects the p local DOS on a Si atom. For instance, a sharp and intense resonance is observed at the threshold of spectra of molecular species. This resonance has been associated for molecular spectra to a Si 1s electron promoted to σ* orbitals with appreciable Si 3p character (Hitchcock et al. 1993) and its energy position and intensity vary with the nature of the first Si neighbours. An alternative way to interpret the experimental XANES spectra is to use the full multiple scattering formalism, where the effect of the core hole is also taken into account. This has been applied with success to reproduce the polarisation dependence of XANES of α-quartz, for which very large clusters are required (Briois et al. 1993), and of β-SiC (Sainctavit 1993) using the CONTINUUM code developed by Natoli et al. (1980). It has been shown that the structures B and C (Figure 5.4.5) in the SiC XANES spectrum can be related to the size of the microcrystals through EXAFS information and full multiple scattering calculations. As a passing note one can observe the very peculiar shape of the β-SiC XANES. It has to be compared with the one of Si—(CH3)4 (Bodeur et al. 1990), where Si is also surrounded by 4 C, the XANES of which exhibit a prominent edge resonance. The effect of the longrange order in β-SiC appears as a corner cut-off of the edge. In all Si reference compounds, the structures, present between the resonance at the edge and the broad resonance located about 20 eV above the threshold, are strongly dependent on the structural medium-range order. The different calculations described above show definitely this point. One can notice as an illustration that α- and β-Si3N4 can be differentiated by taking a glance at the XANES: the β form, where the unit cell is half as large as that of the α phase, presents the most structured XANES spectrum. The second point to notice is the chemical shift existing when going from pure Si to silicon oxide. The threshold position is strongly dependent upon the nature of the atoms which built the tetrahedral environment of Si, and is directly related to the electronegativity of these atoms, that is, to the more or less large density of p conduction states on Si sites.

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Page 197 This allows us to give qualitative conclusions concerning the XANES spectra of the Si/C/N compounds (Figure 5.4.15(a)). First of all the shape (which remains similar whatever the C/N ratio) of the near edge domain is constituted of a well-defined threshold, with a broad resonance (A) followed only by a broad structure (C) at about 15 eV above the edge. The fact that all the XANES structure between A and C is washed out indicates a medium-range disorder between the tetrahedra, and clearly confirms the amorphous state of the as-formed sample. Second, looking at the energy position of the absorption threshold, we see that it varies continuously but stays at an intermediate position between that of SiC and Si3N4 v. the C/N ratio. We think that this is a direct evidence for a mixed environment of a SiCxN4−x type. Actually, a demixed environment would lead to a two-steps edge, as it can be observed for samples heated above the crystallisation temperature. Figure 5.4.15(b) is a clear illustration of the demixing process accompanying the beginning of the crystallisation: the XANES

Figure 5.4.15 (a) Evolution of the Si—K XANES spectra of Si/C/N powders for different C/N ratio (b) Si—K XANES spectrum for the two references β-SiC and α-Si3N4 and for the HMDS44 sample after heating above the crystallisation temperature. The last one appears clearly as a combination of the two references ones.

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Figure 5.4.16 Variation of the edge position with the C/N ratio. becomes very structured near the threshold, and looks more like a linear combination of both SiC and Si3N4 spectra: the two kinds of tetrahedral environment are now present in different proportions, depending on the annealing conditions. Therefore, concerning the as-formed samples, a local chemical ordering can be excluded by just looking at the XANES spectra. Figure 5.4.16 shows the evolution of the edge position when the C/N ratio increases, and C atoms are replaced by N atoms in the Si environment. Such a mixed configuration for Si atoms was also evidenced previously for laser synthesised Si/C/N nanoparticles, using gaseous Si precursor (Gheorghiu et al. 1992) or an aerosol precursor with a low C/N value (Driss-Khodja et al. 1996). Such a mixed environment for Si atoms has already been suggested from NMR studies by Suzuki et al. (1995), who concluded to the presence of SiCN3 tetrahedra in Si/C/N systems at high N content. EXAFS – Analysis of the intermediate order: modelisation of the experimental data Beyond the first peak of the pseudo-radial distribution function (FT), which corresponds to the first coordination shell of the Si atoms, one can observe (see Figure 5.4.12) two small peaks in the range 2–3.5 Å (uncorrected from phase-shift). These peaks are very little dependent on the C/N ratio, and this means that the intermediate order is in average independent of the chemical composition. Focusing on one sample (HMDS43 with a C/N equal to 0.7), an attempt has first been made to fit these peaks with Si—Si contributions, with a wide distance distribution, as one can expect in an amorphous covalent network where the close environment is well defined, but where there can be an enhancement of the bond angle distortions. Figure 5.4.17(a) shows the very bad quality obtained for the fit with only Si—Si second environment. Fitting the experimental data by Si—N pairs at a distance of 2.78 Å only, or by Si—C pairs at 2.84 Å, strongly increases the fit quality (Figure 5.4.17(b)). The improvement obtained by this solution evidences the fact that most of the silicon tetrahedra are not in majority linked by their corners. Such a two-shell simulation assumes also the presence of C—N bonds in the amorphous network. From these preliminary results, a model is proposed (Figure 5.4.18), where mixed tetrahedra are randomly linked through C—N bonds.

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Figure 5.4.17 (a) Best fits of the filtered second peak (2–2.8 Å) of the FT (Figure 5.4.13) for HMDS43 sample with one Si—Si shell; (b) One-shell fit of the filtered second peak corresponding to Si—C at 2.84 Å as second neighbours for Si atoms. A fit corresponding to Si—N at 2.78 Å gives also a good agreement with the experiment.

Figure 5.4.18 Proposed model for the amorphous atomic structure of Si/C/N with C/N = 0.7.


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Page 200 This model was built with a molecular mechanic software (CERIUS). An analysis of various two, three and four bodies interactions allows an evaluation of the forces on each particle. The energy expression deduced for N-body systems, such as our model, is calculated by superposition of single terms (bond, angle, torsion, inversion). The exact nature of each term depends on the particular loaded force field. Here, we used the DreidingII force field (Mayo et al. 1990), which is recommended for structures with novel combination of elements and those for which there is little or no experimental data. A computational module allows then the minimisation of the energy expression of the model and leads to an optimum geometry. This average model is composed of SiN3C tetrahedra linked through C—N bonds, in order to form six-edges polyhedra. The only constraints imposed before the relaxation process are the first neighbours distances for Si—C and Si—N bonds. They were fixed to those found by the experimental analysis (1.89 Å for Si—C and 1.74 Å for Si—N). The coordinates of the atoms in the model are then extracted and used as input to the FEFF code (Rehr et al. 1992), in order to compare the corresponding calculated EXAFS signal, which includes multiple scattering paths, to the experimental one. In that calculation, the default values were used for the input parameters. The calculation is performed with Si1, Si2 and Si3 playing the role of the central atom (see Figure 5.4.18). The number of paths is limited to four, and the maximum distance (Rmax) to 6 Å. The first shell gives the main contribution. Beyond it, one finds a large distribution of Si N and Si C two-legs paths between 2.6 and 2.85 Å, followed by three-legs paths corresponding to Si/C/N and Si/N/C, the distances of which are found between 2.9 and 3.15 Å. The first Si—Si distances are found around 3.3—3.4 Å. It happens that the best fit is obtained when averaging the results on Si2 and Si3 as central atom. Figure 5.4.19 shows the good agreement obtained by this way. The calculated result obtained with Si1 as central atom leads to an enhancement of the second neighbours peaks. This amplitude difference, due to the fact that Si2 and Si3 do not have all their second neighbours, may have two explanations. First, this could evidence a static and/or dynamic disorder on the outer shells distance, due to distortions of the valence bond angles of C and N first neighbours. As it is well known in EXAFS and has been discussed previously,

Figure 5.4.19 Comparison of the FT (modulus and imaginary part) for the experimental spectrum of the HMDS43 Si/C/N sample, with the FEFF calculation using the atomic coordinates of the model proposed in Figure 5.4.18.

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Page 201 such a disorder appears as a reduction of the coordination number, then the environment of Si2 and Si3 could represent a more realistic situation than Si1. This amplitude reduction can also be interpreted as an anharmonic effect originating in the size of the nanometric particles. But nanometric Si/C/N powders concern particles of an average diameter of about 30 nm, with only about 10% of the atoms in surface. The only disorder taken into account intrinsically in our simulation is that coming from the calculated positions of the atoms, with no further parameter: this accounts for the structural disorder and its deviation to a gaussian distance distribution. Second, this lower coordination number might be real and could come from dangling bonds of N and C, or to a bonding to H atoms. In terms of an EXAFS analysis, these two structural situations are indistinguishable, and can be both present in the material. Finally, such a limited model is not able to reproduce the XANES part of the spectrum, probably due to the fact that the muffin-tin approximation in the calculation of the potential is not appropriate. 5.4.3.2 Evolution of the local order during the pyrolysis process 1400°C Our purpose is then to understand the structural evolution before the crystallisation, when the sample is treated under N2 at 1400°C for 24 h, 1500°C for 4 h and up to 1600°C. The observed changes are rendered on the FT as modifications of the second and third neighbours in the intermediate-range order (Figure 5.4.20). The comparison of the FT of the as-formed powder and the annealed ones underlines the slight evolution of the short-range order after annealing at 1400°C (Figure 5.4.20), while more important change occurs in the intermediate-range order. In order to determine the local order around Si, a two-shell fit of the first peak has been performed on the basis of Si—C and Si—N pairs. We obtained the structural parameters presented in Table 5.4.5. As for the as-formed samples, a two-shell model with partial NSi—C

Figure 5.4.20 FT (modulus and imaginary part) of the k2 weighted EXAFS oscillations for HMDS44, asformed, and annealed 24 h at 1400°C.

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Page 202 Table 5.4.5 First-shell structural parameters obtained with a two-shell fit (Si—C + Si—N) for two compositions and for two annealing temperatures (1400°C for 24 h and 1500°C for 4 h). ΔN/N = 10%, ΔR= 0.02 Å Sample Si—C Si—N NSi—C+Si—N NSi—C RSi—C(Å) NSi—N RSi—N(Å) HMDS42-1400 1.4 1.87 2.8 1.74 4.2 HMDS42-1500 2.3 1.88 1.7 1.75 4 HMDS44-1400 1.2 1.87 2.8 1.73 4 HMDS44-1500 1.7 1.87 2.3 1.73 4

Figure 5.4.21 Proposed model for the amorphous atomic structure of the sample annealed at 1400°C (24 h). and NSi—N provides an average picture of the local arrangement of the C and N atoms in the first coordination shell of the absorbing Si atoms. Beyond the first peak of FT (Figure 5.4.20), the structures relative to intermediate-range order evolve with the annealing temperature. Unlike the results obtained for the as-formed powders, and as early as 1400°C, one has to introduce Si—Si bonds as second neighbours. A new 3D structure based on EXAFS results (short range and intermediate order) was built and then tested through multiple scattering calculations at the Si—K edge of central Si atoms of the structure. The model presented in Figure 5.4.21 is related to the 1400°C annealed temperature. It is composed of rigid SiN3C tetrahedra (Si—N = 1.74±0.02 Å and Si—C = 1.89 ±0.02 Å linked by C—N bonds with large distance distributions (1.46–1.55 Å) or partially linked by their corner. The evolution compared to the as-formed model is the possibility to link the SiN3C tetrahedra by their corner. This model was in a second step tested on 3 Si central atoms through multiple scattering calculation using the FEFF code. The average calculated spectrum and its FT presented in Figures 5.4.22(a) and (b) are compared to the corresponding experimental parts. The agreement between experiment and calculation is good. Particulary, the evolution beyond the first peak is well reproduced. The evolution in the way in which tetrahedra are linked is confirmed by the calculations: Si atoms as second neighbour of Si

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Figure 5.4.22 Comparison of (a) k × χ (k) and (b) FT (modulus and imaginary part) for the experimental spectrum of the sample (HMDS44), annealed at 1400°C for 24 h with the FEFF calculation, using the atomic coordinates of the model proposed in Figure 5.4.21. atoms appear at 1400°C. This result is consistent with the increase of the density which is also observed in the powders. Concerning the static disorder, the Si atoms taken as absorbing atoms in the FEFF calculations still have an incomplete second coordination shell, which simulates the static disorder. The increase of the coordination number compared to the as-formed model simulates the decrease of the static disorder when the annealing temperature increases. 1500°C The same approach on the XAS spectra, corresponding to the 1500°C annealed sample, shows this time an evolution of the NC/NN ratio. Figure 5.4.23 presents the FT of HMDS44 annealed at 1400°C (24 h) and 1500° C (4 h). Strong changes occur in the first coordination

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Figure 5.4.23 FT of the k2 weighted EXAFS for HMDS44 annealed at 1400°C (24 h), 1500°C

Figure 5.4.24 Evolution of the Si K XANES spectra at 1500°C with different C/N ratio. shell around the Si atoms. The amplitude of the first peak in the FT decreases at 1500°C (4 h). Table 5.4.5 gives the results obtained for HMDS42 and HMDS44 with a two-shell fit (Si—C and Si—N) of the first peak. The evolution of the ratio NC/NN as a function of the C/N ratio is still linear. The first coordination shell of the Si atoms is enriched in C. This result is confirmed by the XANES study of the samples annealed at 1500°C (4 h). For HMDS42, HMDS43 and HMDS44, the spectra still exhibit a unique threshold, the energetic position of which is intermediate between those of SiC and Si3N4 (Figure 5.4.24). A small shift towards lower energies of the threshold position is observed (


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Page 205 a first approximation by fitting the second peak (1.9–3.2 Å) of the FT, one can conclude to an increase of the number of Si—Si second neighbours corresponding to a structural reorganisation, which may appear as an early beginning of crystallisation. This increase of the number of Si—Si second neighbour pairs is confirmed by FEFF calculations performed on a cluster, where SiC2N2 tetrahedra are in a majority linked by their corners. The Si atoms chosen as central atoms in the calculations have now a nearly complete second coordination shell and the decrease of the structural disorder is going on. These results assume that strong modifications of the short- and medium-range order in the powders occur at 1500°C in the amorphous state. Beyond 1500°C The FT corresponding to the samples annealed at 1500°C (Figure 5.4.23) exhibits the well-known focusing peak at about 5.8 Å, which corresponds to the crystallised SiC phase (ZnS blende structure) of the 8th shell. The compound is crystallised in a mixture of β-SiC and Si3N4 the proportion of which can be estimated by fitting the first shell. The observation of the XANES spectra reinforces the conclusions drawn from the fitting procedure. Until 1500°C (for a 4 h annealing time), the shape of the near edge domain is constituted of a well-defined threshold (Figure 5.4.24), with a broad resonance (A) followed by a broad structure (C). The lack of structures between A and C is the sign of a disordered structure. The unique threshold inserted between those of β-SiC and α-Si3N4 is a proof of the presence of mixed tetrahedra. Above this temperature, or with longer annealing time, the XANES becomes more structured showing already the characteristic structures of β-SiC, and a two steps threshold (Figure 5.4.15(b)). We can conclude to the presence of a mixture of β-SiC and α or β-Si3N4. From the chemical composition and assuming a complete crystallisation, one can estimate the proportions of the two different phases, about 35% of β-SiC, and 65% of α-orβ-Si3N4 for HMDS44. In order to check this assumption, mixtures of two crystallised powders (β-SiC and β-Si3N4) for different proportions were tested. Figure 5.4.25 compares the XANES of a mixture constituted of 70% of Si3N4 and 30% of SiC and the one of the HMDS44 powder treated at 1600°C for 4 h. The agreement between the two spectra confirms that the crystallisation is achieved. At this point one has to underline that the slight differences observed on the two XANES of α- and β-Si3N4 does not allow one to determine definitely which phase is present in the final powder. 5.4.4 Local structure of pre-alloyed Si/C/N/Al/Y nanopowders studied by XAS at the Al—K edge using fluorescence yield detection 5.4.4.1 Introduction The aim of the present part is to determine the atomic environment of Al both in as-formed SiCNAlY and annealed powders, in order to get information on a possible formation of the (Y)SiAlON liquid phase during sintering, necessary for producing dense ceramics and leading to an intergranular phase after cooling. We compare it with the structure found in crystalline Al(O)N, with YSiAlO(N) amorphous phases, and with a sintered sample where the intergranular phase is present.

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Figure 5.4.25 Comparison of the XANES spectrum of the sample annealed at 1600°C with one of the mixture of the two crystallised phases (30% of β-SiC and 70% of β-Si3N4). 5.4.4.2 Experiment Al—K edge (1559 eV) XANES and EXAFS spectra were recorded at room temperature at beamline SA32 of the 800 MeV positron storage ring SuperACO at LURE, by collecting the total electron yield for the references compounds and the fluorescence yield for the dilute samples. We used a mono-element Ge detector (EurisysMesures), with an energy resolution of 150 eV. The monochromator was a scanning two-crystals set-up with quartz (1010) crystals. As a consequence, the EXAFS data were recorded only up to 1830 eV. For the XANE Spart of the absorption spectra, a constant background has been subtracted in the pre-edge region, and except for α-Al2O3, the spectra were normalised to the atomic absorption above the threshold. All Al XANES spectra were calibrated with an Al metallic foil at the inflection point (1559 eV). The extraction of the EXAFS oscillations was performed in a classical way: modelisation of the atomic background by a fifth-order polynomial, and normalisation using the Heitler approximation. FTs of the k2 weighted oscillations were calculated over the same energy range for all compounds and references. The chemical compositions of the studied samples as well as those of the amorphous references are presented in Table 5.4.6. 5.4.4.3 Results As we are dealing with amorphous compounds, where XANES multiple scattering calculations in the ‘muffintin approximation’ are difficult to perform, our approach was to measure aluminium oxides and nitrides, as well as amorphous YSiAlO(N), pure or embedded in the sintered material. From direct comparison with that in model compounds, the Al environment of unknown compounds can be ‘fingerprinted’ in terms of differences in XANES (Ildefonse et al. 1994, 1998 and references therein). Experimental Al spectra of references compounds are plotted in Figure 5.4.26(a). The differences observed between berlinite (AlPO4, with AlIV—O) and corundum (α-Al2O3, with AlVI—O) have been fully investigated and simulated previously (Cabaret et al. 1996). In AlN, Al is coordinated to four N atoms. Its XANES spectrum presents a large resonance located 3 eV below the characteristic white line

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Page 207 Table 5.4.6 Elemental analysis (at%) of the amorphous reference phases, and of the pre-alloyed samples = as-formed, AN = annealed at 1600°C) Si C N O Al Y HSAl17-AF 33.5 26 33.5 6.2 0.6 HSAl17-AN 44 24 29.9 1.6 0.4 YSiAlON 12.3 — 7.5 54.7 18.5 YSiAlO 6.8 — — 63.6 17.7

(AF 0.2 0.1 12.3 11.9

Figure 5.4.26 (a) Al—K edge XANES spectra for references where Al is fourfold coordinated (AlPO4 and AlN), sixfold coordinated (Al2O3), and in a mixture of sites (AlON), (b) Al—K edge XANES spectra for amorphous references. (Reprinted from Flank et al. (1999) with kind permission of International Union of Crystallography.)

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Page 208 of AlIV—O in AlPO4. The oxynitride phase AlON is usually described as a solid solution in the Al2O3—AlN structure, isostructural to γ -Al2O3, where it is expected that Al has a range of local tetrahedral and octahedral environments, including mixed anion (i.e. N and O) coordinations (Dupree et al. 1988). The XANES spectrum of this last system is characteristic of a mixture of sites, but we have to point out that the threshold position is more than 3 eV below the one found in corundum, in line with the presence of Al fourfold coordinated, and that the spectrum presents some structures characteristic of an ordered structure also evidenced by XRD. In Figure 5.4.26(b), XANES spectra of the amorphous phases YSiAlO(N) are compared with the Al spectrum of the sintered compound, where Al is supposed to be both in the intergranular phase YSiAlON and in the βSiAlON obtained after annealing treatment. The sintered compound (SC2X) is composed of 20 wt% of a SiC laser synthetised nanopowder and 62 wt% of a Si3N4 conventional powder (UBE-SNE10, α phase) and 12 wt % of Y2O3 and 6 wt% Al2O3. Except for a very small shift towards high energy values of the threshold of the sample without N (YSiAlO), all the spectra are identical, reflecting a structural similarity between these different phases. From previous observations (Ildefonse et al. 1998), the edges shapes: a small white line located at 1566.6 eV followed by a wide resonance, lead us to the assumption that the dominant site for Al atoms in these phases is fourfold with some amount of penta and hexa coordinated aluminium. But it is difficult to conclude about the existence of mixed environment. Actually the edge position in fourfold coordinated Al compounds depends on the nature of the second neighbour of Al atoms (Ildefonse et al. 1998). As well, qualitative conclusions can be drawn about the coordination number and the nature of the neighbours of aluminium in pre-alloyed Al/Y/SiCN powders. The XANES of the as-formed powder presents a double structure (Figure 5.4.27). The threshold position is located at the same position than for the amorphous references phases, that is, between AlPO4 and AlN. We cannot then conclude to a single Al environment. After annealing at 1600°C, the Al—K edge is much more structured (Figure 5.4.27): a resonance A appears at the same position than that of Al fourfold coordinated to O atoms indicating the presence of AlO4

Figure 5.4.27 Al—K edge XANES spectra for pre-alloyed Al/Y/SiCN nanopowders: evolution with annealing and comparison with references.

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Figure 5.4.28 Comparison of the FT (k2 weighted oscillations over 2–7 Å−1 EXAFS k range) for AlN and HSAl17 annealed at 1600°C. (Reprinted from Flank et al. (1999) with kind permission of International Union of Crystallography.) sites, the edge position is shifted to lower energy, when compared to the as-formed spectrum indicating the presence of AlN4 sites, and finally the resonances B and C also observed for the AlN crystalline structure prove the formation of AlN crystallites. An EXAFS analysis of these compounds has been undertaken at the Al—K edge. The pseudo-radial distribution function of the AlN phase presents four peaks. The first one comes from the first shell of 4 N atoms at R = 1.88 Å. The three following peaks are a mixture of the different contributions of the next shells. AlON’s FT exhibits three main peaks. The amorphous phases give both significant peak at a smaller distance than the first peak of AlN, but a similar amplitude. The FT of the sample HSAl-17 annealed at 1600°C (Figure 5.4.28) confirms the hypothesis drawn from the observation of the XANES: Al atoms are mainly present in a crystalline AlN type structure. 5.4.4.4 Discussion The unknown compounds have also been investigated by NMR, which reveals that Al can be found in three different sites in the as-formed compound: tetra, penta and octa coordinated to O atoms, while after annealing, four sites are still present: Al octa, Al tetra with AlO4, AlN3O, and AlN4. Because of the limited k range which does not allow a sufficient distinction between the AlN and the other phases, the EXAFS spectrum seems dominated by the contribution of Al atoms in the ordered phase AlN. But the shape of the XANES induces to claim the presence of other different disordered sites, consistent with NMR results. References Bodeur, S., Nenner, I. and Millie, P. (1990) Phys. Rev. A, 41, 252. Briois, V., Sainctavit, P. and Flank, A.-M. (1993) Jpn. J. Appl. Phys., 32, (Suppl. 32-2) 52–54. Cabaret, D., Sainctavit, Ph., Ildefonse, Ph. and Flank, A.-M. (1996) J. Phys. Condens. Matter, 8, 3691.

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Page 210 Cauchetier, M., Croix, O., Herlin, N. and Luce, M. (1994) J. Am. Ceram. Soc., 77, 993. Cauchetier, M., Croix, O., Luce, M., Baraton, M. -I., Merle, T. and Quintard, P. (1991) J. Eur. Ceram. Soc., 8, 215. Dalba, G., Fornasini, P., Diop, D., Grazioli, M. and Rocca, F. (1993) J. Non-Cryst. Solids, 164–166, 159. Diop, D. and Grisenti, R. (1995) Physica B, 208 and 209, 164. Driss-Khodja, M., Gheorghiu, A., Dufour, G., Roulet, H., Sénémaud, C. and Cauchetier, M. (1996) Phys. Rev. B, 53,8. Dupree, R., Lewis, M. H. and Smith, M. E. (1988) J. Appl. Cryst., 21, 109. Flank, A.-M., Armand, X., Cauchetier, M. and Mayne, M. (1999) J. Synchr. Rad., 6, 512. Gheorghiu, A., Sénémaud, C., Roulet, H., Dufour, G., Moreno, T., Bodeur, S., Reynaud, C., Cauchetier, M. and Luce, M. (1992) J. Appl. Phys., 71, 4118. Hitchcock, A. P., Tyliszczak, T., Aebi, P., Xiong, J. Z., Sham, T. K., Baines, K. M., Mueller, K. A., Lu, X. H., Baribeau, J. -M. and Jackman, T. E. (1993) Surf. Sci., 291, 349. Iawamoto, Y., Kikuta, K. and Hirano, S. (1998) J. Mater. Res., 13 (2), 353. Ildefonse, Ph., Cabaret, D., Sainctavit, Ph., Calas, G., Flank, A.-M. and Lagarde, P. (1998) Phys. Chem. Miner., 25, 112. Ildefonse, Ph., Kirkpatrick, R. J., Montez, B., Calas, G., Flank, A.-M. and Lagarde, P. (1994) Clays Clay Miner., 42 (3), 276. Lagarde, P. and Flank, A.-M. (1986) J. Physique, 47, 1389. Lagarde, P., Flank, A.-M. and Itie, J.-P. (1993) Jpn. J. Appl. Phys., 32 (Suppl. 32-2), 613. Lagarde, P., Flank, A.-M., Tourillon, G., Liebermann, R. C. and Itie, J.-P. (1992) J. Phys. I France, 2, 1043. Lee, P. A. and Pendry, J. B. (1975) Phys. Rev. B, 28, 95. Lengeler, B. and Eisenberger, P. (1980) Phys. Rev. B, 21, 4507. Lytle, F. W., Sayers, D. E. and Stern, E. A. (1975) Phys. Rev. B, 11, 4825. Mayne, M., Armand, X., Cauchetier, M., Doucey, B., Balhoul-Hourlier, D. and Goursat, P. (1998) Adv. Sci. Tech. Proceedings of the 9th CIMTEC, World Ceramics Congress Ceramics: Gelting into the 2000’s, Part B. (ed. P. Vincenzini), p. 211. Mayo, S. L., Olafson, B. D. and Goddard III, W. A. (1990) J. Phys. Chem., 94, 8897. McKale, A. G., Knapp, G. S. and Chan, S. K. (1986) Phys. Rev. B, 33, 841. McKale, A. G., Veal, B. W., Paulikas, A. P., Chau, S. K. and Knapp, G. S. (1988) J. Am. Chem. Soc., 110, 3763. Natoli, C. R., Miserner, D. K., Doniach, S. and Kutzler, F. W. (1980) Phys. Rev. A, 22, 1104S.1. Pascarelli, S., Boscherini, F., Mobilio, S. and Evangelisti, F. (1992) Phys. Rev. B, 45 (4), 1650. Rehr, J. J. and Albers, R. C. (1990) Phys. Rev. B, 41 (12), 8139. Rehr, J. J., Albers, R. C. and Zabinsky, S. I. (1992) Phys. Rev. Lett., 69 (4), 3397. Sainctacvit, P., Petiau, J., Laffon, C., Flank, A.-M. and Lagarde, P. (1991) in X-ray Absorption Fine Structure (ed. S. S. Hasnain), Ellis Horwood, New York, pp. 38–40. Suzuki, M., Hasegawa, Y., Aizawa, M., Nakata, Y., Okutani, T. and Uosaki, K. (1995) J. Am. Ceram. Soc., 78, 83. Teo, B. and Lee, P. A. (1979) J. Am. Chem. Soc., 101, 2815. Teo, B. K. (1986) EXAFS: Basic Principles and Data Analysis, Springer Verlag, Berlin Heidelburg, New York, Tokyo. Wyckoff, R. W. G. (1963) Crystal Structures, J.Wiley & Sons ed., Interscience Publishers. Zabinsky, S. I., Rehr, J. J., Ankudinov, A., Albers, R. C. and Eller, M. J. (1995) Phys. Rev. B, 52 (4), 2995.

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Page 211 5.5 Neutron and X-ray diffraction applied to the study of the medium-range order in SiCN nanopowders François Ténégal, Brigitte Bouchet-Fabre, Jean Dixmier and Robert Bellissent 5.5.1 Diffraction applied to the study of disordered compounds The early stages of the crystallisation of SiCN powders involve precipitation phenomena through atomic diffusion processes within the disordered network. The goal is to investigate the induced changes of the local order leading to nuclei of the equilibrium phases. The Fourier transform (FT) of the interference function S(q) yields the radial distribution function (RDF) G(r) that gives the values and number of interatomic distances between pairs of atoms. But since we measure the beam intensities and not the amplitudes we lose the phases of the waves and no information can be derived from G(r) on the statistics of higher order functions of 3, 4, etc. atoms. This knowledge is needed to build a three-dimensional image of the local order. For example, the G(r) results cannot tell if Si—C and Si—N pairs belong to mixed SiCxN4−x tetrahedra or to demixed SiC4 and SiN4 ones. It is the reason why the inverse calculation of S(q) has been performed for three-dimensional atomic models as suggested by the G(r) results and have been compared with the experimental ones. 5.5.2 Early stages of the crystallisation of Si/C/N-HMDS nanopowders studied by X-ray and neutron diffraction: Nano-ordered Si/C/N structures 5.5.2.1 Contrast effect between neutron and X-ray diffraction measurements Neutron and X-ray diffraction measurements were performed on SiCN nanopowders to take advantage of the drastic differences of the weighting factors on silicon (Si), nitrogen (N) and carbon (C) affecting diffraction spectra to selectively observe Si and N subnetwork. Such contrast effect was also used for the structural study of SiCN ceramics obtained by pyrolysis of polysilycarbodiimide polymer (Dürr et al. 1999). The total structure factor obtained from diffraction measurements can be written as (Warren 1969):

(5.5.1) where Sij (q) are the partial structure factors for the atomic pair i–j and Wij (q) are the weighting factor which depend on the composition of the studied samples and on the

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Figure 5.5.1 Weighting factors Wij(q) for the six pairs of atoms in the HMDS44 (1400°C for 24 h) samples, defining the profile of the X-ray and neutron high angle spectra. The values are also given for the Si—O pair. Note that Wij(q) varies with q in X-ray experiments but remains constant for neutron experiments. The important fact is that the subnetwork of Si is reinforced by X-ray while the N one is preferentially seen by neutron. energy of the incident radiation. A complete determination of the six partial structure factors (Si—Si, Si—N, Si —C, N—N, N—C, C—C) requires six experiments. Unfortunately, anomalous wide angle X-ray scattering measurements (AWAXS) are experimentally unreachable at Si, N and C photoabsorption edges and then the complete determination of partial structure factors is impossible. The difference in the nature of the radiation used in X-ray and neutron experiments is responsible for a strong contrast effect on N and Si amorphous subnetwork. Figure 5.5.1 compares the Faber–Ziman weighting factors (Waseda 1984) (Wij(q)) calculated for X-ray and neutron radiations for HMDS44 (annealed at 1400°C for 24 h). In the range of composition of the HMDS powders, the contrast is very strong between the X-ray spectra which reflects for 80% Si—Si and Si— N correlations, whereas the neutron spectra reflects for 80% N—N, N—Si and N—C correlations. 5.5.2.2 Experimental and data analysis Neutron The neutron experiments have been performed on the 7C2 spectrometer at Laboratoire Léon Brillouin (LLB) at Saclay (CEA-France). The neutrons wavelength (λ = 0.7 Å) was selected with a copper monochromator (Cu (111)). The diffracted intensity was measured with a 640 BF3-cells detector which allows to perform measurements with a constant Δθ path. The larger measurable scattering angle 2θ is 128°. The corresponding scattering vector q is equal to 16 Å−1(q = 4π sin θ/λ). The experimental data are corrected for absorption effect. The contribution of the container Ic(q) was subtracted from the measured diffracted intensity I(q). The Placzek correction was applied (Placzek 1952) to correct I(q) from inelastic scattering process. The spectra were then normalised by the coherent scattering cross-section

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Page 213 of the sample to obtain the structural factor S(q) (small contribution of the incoherent cross-section):

(5.5.2) where I(q) is the total elastic scattered intensity, S(q) is the total structure factor, N is the number of atoms, σ the coherent scattering cross-section (Neutron 1992) and α a constant. X-ray The wide angle X-ray diffraction experiments were performed at LURE-DCI (Orsay-France) on the wiggler DW31B beamline equipped with a 2-circles goniometer using a scintillator. A 2-crystal (Si (220)) monochromator with sagitally focalisation and a double setup of distant slits define the energy resolution. A 20.2 keV incident energy has been chosen in order to use a q window going from 1 to 16 Å−1 (the same as for neutron experiments). Due to low absorption coefficient of the studied materials, the spectra have been recorded in the reflection mode with a nearly grazing incidence (θ0 = 2.5°). The Compton contribution to the total signal has been removed using the analytic Hajdu’s formula (Hajdu 1971) adapted for light elements. The reduction of the recorded data after correction for absorption and asymmetric geometry has been made by integration assuming:

(5.5.3) where Icomp(q) is the modelised Compton contribution and N the number of diffracting atoms, I(q, E) is the total experimental scattered intensity. The elastic scattered intensity for one atom i(q, E) can be written as: (5.5.4) the total structure factor and fi are the

and where complex atomic scattering factors given by:

(5.5.5) are calculated for each element from tabulated values (Waasmaier and Kirfel 1995; Sazaki 1984). ci are the atomic fractions of each element in the powders. 5.5.2.3 Structural transition in the amorphous state Evaluation of the hydrogen content The strong incoherent (q dependent) and inelastic scattering of neutrons by hydrogen gives a form factor to the H-rich powders neutron spectra. Figure 5.5.2 presents crude spectra of the composition HMDS44 for the as-prepared sample and for annealed samples at 1000°C (4 h), 1500°C (4 h) under N2 atmosphere. The reduction of the form factor (decrease with q) with temperature is the consequence of the hydrogen effusion during annealing treatment. At 1500°C, H has totally effused. The residual decrease of the intensity at high q (1500°C) comes from the minor contribution

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Figure 5.5.2 Evolution of the HMDS44 original spectra after the usual corrections with the annealing temperature. Note the incoherent scattering form factor of H which allows to measure its content in the asformed powder (21.5 at%) and in the 1000°C sample (6 at%). At 1500°C, H has totally effused. of inelastic scattering process concerning the other atoms (Si, N, C). For the sample which present an important form factor, a quantitative evaluation of the atomic proportion of hydrogen in the powder can be performed (Dixmier et al. 1994). This can be achieved by making the assumption that at high q values, the measured intensity corresponds to the coherent contribution from Si, C, N atoms without the H contribution which has drastically decreased. The theoretical asymptotic scattering cross-section σ is calculated for one average atom using the tabulated scattering cross-section of the elements and the composition. The I(q = ∞) theoretical value is equal to α N σ, where N is the number of atoms (hydrogen excluded) and α is a constant. The atomic form factor for one H atom is evaluated using the inelastic scattering theory and the H content is then estimated by a trial and error method which consists in subtracting from the spectrum a certain proportion of the background, in order to force the curve to oscillate around the constant asymptotic value I(q = ∞). The subtracted intensity corresponds to the contribution from H atoms: α NHσH, where NH is the number of H atoms and σH the incoherent scattering cross-section for one H atom (σH = 80 barn for q = 0).

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Page 215 From this value and from the asymptotic value I(q = ∞), we calculated the H content. One will find 21.5 at% for the as-prepared powder and 6 at% for the 1000°C annealed powder. Structural evolution of Si/C/N nanopowders during the pyrolysis process: 1400°C → 1500°C SHORT-RANGE ORDER: EXISTENCE OF C—N BONDS The coupled X-ray and neutron diffraction experiments have been performed on the composition HMDS44 for two annealing temperatures (1400°C for 24 h and 1500°C for 4 h). X-ray absorption spectroscopy (XAS) measurements underlined strong modification in the short-range atomic structure around Si atoms between 1400°C and 1500°C. An increase of the NSi—C/NSi—N ratio is observed. The neutron (Sn(q)) and X-ray (SX (q)) structure factors for HMDS44 (1400°C for 24 h and 1500°C for 4 h) are reported in Figure 5.5.3. As expected, X-ray and neutron spectra, which respectively reflect silicon and nitrogen subnetwork, are very different. The structure factors exhibit a set of broad diffuse halo characteristics of covalent amorphous compounds. At low q values, on the neutron spectra, one can find a first sharp diffraction peak (FSDP) related to the nitrogen medium-range atomic arrangement in the powders. Dixmier (1992) showed that in a disordered network in which four coordinated atoms (Si or C) are connected to three coordinated atoms (N), the structure factor S(q) exhibits a main peak near 3 Å−1 generated by the first neighbour and two prepeaks at 2 and 1.5 Å−1 generated respectively by the mixture of four and three valencies. At 1400°C, we expect that the extreme sharpness of the peak located near 3 Å−1 is a signature of the stacking of disordered layers on coherent length of 2–3 nm. In contrast, the Si dominated X-ray spectrum exhibits a diffuse first peak, but small sharp diffraction structures for q > 6 Å−1. These small peaks indicate the existence in the amorphous state of well-defined long-range correlations in the Si amorphous subnetwork. At 1500°C, the shape of the structure factors change. The

Figure 5.5.3 Neutron (Sn(q)) and X-ray (SX(q)) interference functions for HMDS44 annealed at (a) 1500°C for 4 h and (b) 1400°C for 24 h . The two spectra (Sn(q) and SX(q)) reveal the different evolutions of the Si and N subnetwork with the annealing temperature.

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Page 216 prepeak of neutron spectrum is split into two components and the morphology of the successive haloes changes. For X-ray spectrum, the first halo becomes sharper. This evolution is the signature of the decrease of the chemical and topological disorder of the Si amorphous subnetwork when the annealed temperature goes from 1400°C to 1500°C. Strong structural modifications occur in the short-range and medium-range order (MRO) of the amorphous compound at 1500°C. These evolutions are consistent with those underlined by extended X-ray absorption fine structure (EXAFS) measurements concerning the evolution of the local order around Si atoms with annealing temperature. The EXAFS results show that the Si subnetwork is very disordered at 1400°C due to the mixture of the configurations for Si Si pairs (Si—N(C)—Si and Si—N(C)—C(N) —Si). The reduction of the number of Si—N(C) C(N)—Si bonds at 1500°C is responsible for the decrease of the topological and chemical disorder. The contrast effect is clearly visible on the corresponding pair of correlation functions G(r) presented in Figure 5.5.4. The first peak is not much affected (hetero-atomic bonds) but the second one is definitively shifted (homo-atomic pairs such as N—N or Si—Si). An attempt to simulate the first peak of the RDF (= 4πρ0r2G(r)) was done. The procedure starts with the modelisation of the distribution of j type atoms around an i type atom by a Gaussian distribution fij(r). This distribution is related to the partial i–j pair correlation function Gij(r) by: where ρ0 is the atomic density (at/Å3), cj the atomic fraction of the element j and Gij (r) the partial pair correlation function associated to the partial structure factor Sij(q). An FTof the r-space contribution is then performed to obtain the q-space contribution Sij(q) and the

Figure 5.5.4 Neutron (Gn(r)) and X-ray (GX(r)) pair correlations functions for HMDS44 annealed at (a) 1500° C for 4 h and (b) 1400°C for 24 h. The common integration limit S(qmax) = 15 Å of the FT generate nonphysical peaks before 1 Å at the same value of r in the both experiments but not with the same amplitude. We expect that due to the strong decrease of the X-ray scattering factors at high q, the neutron description of the short-range atomic structure is more accurate than the X-ray description (particularly for the shortest correlations). The contrast effect is well marked for the homo-atomic pairs N—N and Si—Si regrouped in the second peak near 3 Å.

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Figure 5.5.5 Result of the simulation (neutron) of the first peak of the RDF with only two contributions (Si—C and Si—N) for HMDS44 annealed at 1400°C (24 h). The corresponding structural parameters are given in Table 5.5.1. Faber–Ziman weighting factor Wij(q) of the i–j pair is applied. Then, one goes back to the real space by a new FT of the weighted contribution of the i–j pair (Wij(q)Sij(q)). If several contributions have to be taken into account to simulate the first RDF’s peak, the components are added. The obtained contribution is directly comparable to the experimental RDF since the limits of integration used to calculate the FT are in both cases the same (1.5–15 Å−1). Figure 5.5.5 presents the result of the simulation (neutron) using only two components (Si—C and Si—N) for the 1400°C annealed sample. Note that the values found for the centroid of Gaussian distributions are those found in the crystalline compounds SiC and Si3N4 for the Si—C bond length (1.89 Å) and the Si—N bond length (1.74 Å). The results of the EXAFS analysis revealed that the Si—N and Si—C bond lengths remain similar in the amorphous state compared to the crystalline state. These two contributions (Si—C + Si—N) cannot fit the experimental shape of the first RDF’s peak (Figure 5.5.5). A third shorter contribution must be added. Figure 5.5.6 shows the very good fit (neutron) obtained when a C—N (1.5Å) contribution is added. Following this procedure, two fits were realised for each spectrum taking as origin either Si or C and N atoms in the modelisation of fij(r). Table 5.5.1 gives the structural parameters obtained for the samples annealed at 1400°C (for 24 h) and 1500°C (for 4 h). One can observe that the proportion of C—N bonds decrease with annealing treatment. For the sample annealed at 1400°C, one finds 3N + 1C around one Si atom, and 2.3Si + 0.7C around one N atom. At 1500°C, the proportion of C atoms around one Si atom increases (2.2N + 1.8C) while the proportion of C—N bonds decreases. Let us recall that EXAFS results give 3N + 1C at 1400°C and 2N + 2C at 1500°C. The results obtained by coupling X-rays and neutron diffraction are in very good agreement with those obtained by EXAFS. C—N bonds are taken into account in EXAF Slocal order models and are detected by X-ray photoelectron spectroscopy (XPS) measurements. C—C bonds detected by XPS are not introduced in the simulation since the corresponding weighting factor WCC is lower than 5% (HMDS44). Their contribution to the experimental spectra cannot be unambiguously proved. C—C bonds from free carbon phase were reported in the study by Dürr et al. (1999) for Si/C/N ceramics obtained by pyrolysis

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Figure 5.5.6 Result of the simulation (neutron) of the first peak of the RDF with three contributions (Si—C, Si —N and C—N) for HMDS44 annealed at 1400°C (24 h). The corresponding structural parameters are given in Table 5.5.1. Table 5.5.1 First shells structural parameters obtained with a three-shell fit of the first peak of the radial distributions functions (ΔR = ±0.02 Å, ΔN/N = ±10%, Δσ = ± 0.01 Å). For each compound, two set of simulations (Si—N + Si+C + N—C or N—Si + C—Si + N—C) were realised corresponding to two different choices for the central atom taken for the modelisation of the distribution functions R(Å) N σ(Å) X-ray Neutron X-ray Neutron X-ray Neutron 1400°C (24 h) Si—N 1.76 1.74 2.8 2.8 0.07 0.07 (N—Si) 2.4 2.3 Si—C 1.89 1.89 1 0.9 0.07 0.05 (C—Si) 3.2 3 N—C 1.50 0.7 0.2 1500°C (4 h) Si—N 1.76 1.73 2.1 2.2 0.08 0.06 (N—Si) 2.2 2.4 Si—C 1.91 1.88 1.7 1.8 0.07 0.04 (C—Si) 3.6 3.4 N—C 1.51 0.4 0.18 of a polymer species (C/N = 1.3). In their study, the weighting factor ωcc was of the order of 20% and the peak related to C—C bonds was clearly observed in the neutron correlation pair functions. ANALYSIS OF THE MEDIUM-RANGE ORDER: THE TRANSITION 1400°C(α) → 1500°C(β) The existence of a layered structure evidenced at 1400°C for N atoms must already occur in the initial mineral residue at 1000°C. This should be induced by the possible sp2 bonding of

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Figure 5.5.7 Reduced neutron pair correlation function (4πρ0r(Gn(r) − 1)) in the MRO domain for HMDS44. The peak at 3.9 Å (A) is specific of the a stacking mode of α-Si3N4. This correlation disappears at 1500°C and is replaced by a new one at 3.6 Å. The stacking mode for the N atoms changes at 1500°C. C atoms but also by the flat three coordinated N environment. Figure 5.5.7 presents the MRO of Gn(r) (neutron) for HMDS44 for annealing treatments ranging from 1000°C to 1550°C. At 1000°C, one can observe two small bumps at 3.4 and 3.9 Å which become sharper at 1400°C. The existence in the pair correlation function of a correlation at 3.9 Å is attributed to the alpha stacking mode for the atoms in silicon nitride (α-Si3N4). Calculations using the Debye formula were performed on clusters of silicon nitride (Dixmier et al. 1994) which evidenced a correlation between the peak at 3.9 Å and the α stacking mode in silicon nitride. We performed a calculation of the atomic pair distribution in the two known crystalline forms of silicon nitride (α and β). The result is that the correlation at 3.9 Å in α-Si3N4 can be associated with N—N pairs at 3.82 Å between alternated layers. In the α form, the stacking mode (along the crystallographic c axis) is ABCDABCD…and N—N pairs at 3.82 Å exist between N atoms of the layer A and N atoms of layer C, and N atoms of the layer B and N atoms of the layer D. In the β form, the mode is ABABAB… This modification in the stacking mode leads to the disappearance of the N—N pairs at 3.82 Å and so to the annihilation of the correlation at 3.9 Å in Gn(r). Then, the correlation at 3.9 Å in Gn(r) is the signature of the α stacking mode for nitrogen atoms. At 1500°C, the correlation at 3.9 Å disappears while a new one appears at 3.6 Å on the experimental Gn(r). So, locally the increase of the NSi—C/NSi—N ratio at 1500°C is accompanied by a change in the sequential stacking mode of the layers. We expect that the α-mode is replaced by the β-mode at 1500°C. In the amorphous state, nanometric (2–3 nm) SiCN structures which are reminiscent of a crystalline topological order (α- and β-Si3N4) exist. The amorphous structure cannot definitely be considered as a mixture of an amorphous silicon nitride phase and free C phase since mixed SiCxN4−x tetrahedra are detected by a X-ray near edge structure (XANES) study at the SiK edge and by NMR measurements. Carbon atoms induced structural and chemical disorder in α- and β-type SiNy phases. The term ‘chemical disorder’ should be used with caution here. Indeed, the XANES and RMN studies show the

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Page 220 existence of mixed tetrahedra SiCxN4−x. For temperatures below 1400°C, the environment of Si atoms are composed of 3N + 1C (around the same Si atom) or 4N. This allows us to make the hypothesis of a relatively well chemically ordered SiCN structure in which C atoms are homogeneously distributed. The chemical order is then different from those found in crystallised SiC and Si3N4. At 1500°C, a broadening of the peak relative to Si environment measured by RMN is observed. This broadening is attributed to an increase of chemical disorder around the Si atoms: all the environments are detected for Si atoms. This increase of the chemical disorder is acccompanied by an increase of the NSi—C/NSi—N ratio. In short, before 1400°C we expect the existence of a nano-ordered α-SiCN structure composed of SiCN3 and SiN4 tetrahedra and of C—N bonds. At 1500°C, nano-ordered β-SiCN, in which a beginning of demixion occurs, are formed. Gradients of C concentrations appear at the atomic scale in the β-SiCN structure. Locally, an early formation of six member rings of SiC occurs but the number of mixed tetrahedra remain dominant. 5.5.2.4 Nanocrystalline SiCN structural models Building of the models and calculations From the experimental results, the hypothesis of nano-ordered SiCN structures was tested. This could be achieved by the modelisation of nanocrystalline structures involving experimental data (XAS, RMN, XPS and diffraction). In order to simulate the α → β transition, we have decorated the structure of the two polytypes α-Si3N4 and β-Si3N4 with C atoms substituting Si and N atoms in such a manner that mixed tetrahedra must be formed as well as C—N bonds. C atoms substituted to Si form CN4 tetrahedra and C atoms substituted to N atoms form CSi3 sites (Table 5.2.2). Then, a relaxation of the mechanical stress is performed by using an algorithm which allows the potential energy minimisation of the substituted structure. We used the CERIUS software to build the models and to perform potential energy minimisation (see section modelisation). The bond distances in the models were forced to those given by the experimental analysis (1.74 Å for Si—N, 1.89 Å for Si—C and 1.5 Å for C—N) during the relaxation process. This could be done by the addition in the energetic expression of a term which is zero if the bond distances reach the constraint values. We obtained new periodic (low symmetric) SiCN structures in which C atoms are bonded to Si and N atoms. Figure 5.5.8 presents the α-substituted model which is related to 1400°C. The theoretical scattered intensities of the models are calculated using the Debye formula (Debye 1915):

(5.5.6) where I(q) is the total scattered intensity, rnm the distance between the atoms n and m, fn(q) and fm(q) the scattering factor. The sum is taken over all the atomic pairs of the cluster. The theoretical structure factors are deduced and compared with the experimental ones. The calculations were performed using the INSIGHT program.

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Page 221 Table 5.5.2 Construction of the α- and β-substituted models and structural parameters 1400°C (24 h) 1500°C (4 h) 1Construction of the cluster of crystalline Si3N4 α-Si3N4 β-Si3N4 (P 31c; a = b = 7.753 Å, c = 5.618 Å; (P 63/m; a = b = 7.606 Å, c = 2.909 Å; α = β = 90°, γ = 120°) α = β = 90°, γ = 120°) 16 cells (2a × 2b × 4c) 32 cells (2a × 2b × 8c) 448 atoms (1.5nm × 1.5nm × 2 nm) 448 atoms (1.5 nm × 1.5 nm × 2 nm) 2 New description of the structures Space group → P1 Space group → P1 gathering of 2 cells: (1a × 1b × 2c) → (1a × 1b × 1c′) 3 Periodic substitution of Si and N atoms by C atoms (Experiment – HMDS44) (Experiment – HMDS44) 160Si → 36% Si 38.9% 176Si → 39% Si 41.0% 224N → 50% N 42.8% 176N → 39% N 41.2% 64C → 14% C 15.5% 96C → 22% C 15.9% O 2.9% O 2.4% SiCN3 in the majority SiC2N2, SiCN3 SiN4, SiC3N2 in the minority a few CN4 tetrahedra CN4 tetrahedra 4 Relaxation of the geometric stresses dSi—C = 1.89 Å ± 0.02 Å d Si—C = 1.89 Å ± 0.02 Å dSi—N = 1.74 Å ± 0.02 Å d Si—N = 1.74 Å ± 0.02 Å dN—C = 1.5Å ± 0.02 Å d N—C = 1.5 Å ± 0.02 Å N(C)—Si—N(C) from 100° to 120° N(C)—Si—N(C) from 90° to 125° Si—N—Si(C) from 105° to 135° Si—N—Si(C) from 105° to 135° Si(N)—C—Si(N) from 100° to 120° Si(N)—C—Si(N) from 100° to 130° Results of the calculations and discussion NEUTRON The neutron structure factor Sn(q) of the model relative to 1400°C(α-substituted) and for the reference αSi3N4 are given in Figure 5.5.9. Figure 5.5.10 presents the calculated spectra for the model relative to 1500°C (β-substituted) and for the corresponding cluster of β-Si3N4. The hypothesis of nanocrystal of α-Si3N4 can be dismissed. The structures of the experimental spectrum (1400°C) are too different from those of the α-Si3N4 cluster. The general features of the experimental spectra are well reproduced by the models. In particular, the evolution observed in the low q domain (q < 6 Å −1) characterised by the splitting in two components of the FSDP at 1500°C is reproduced. For q > 6 Å−1, the interference functions exhibit small sharp diffraction peaks as expected since the models are small crystallites of α- or β-SiCN. These structures which are not present in the experimental spectra simulate a too strong topological order in the models which are derived from crystal structures. The geometry obtained after relaxation is close to those of crystalline α- and β-Si3N4. The networks of connections between the atoms are the same as in the crystals since the C atoms are only substituted to N and Si atoms and the mixing of the valence states is the same as in silicon nitride. Nevertheless, the haloes which support the small sharp diffraction structures

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Figure 5.5.8 Nanocrystalline α-SiCN model for the intermediate compositions (HMDS43, HMDS44). This structure is relative to 1400°C. are reproduced as well as the width and position of the first sharp diffraction peak which show that the size of the nano-structures to be considered in the calculations are those of the models. To improve the experiment–theory agreement, we simulated the loss of the medium– long-range order in the model (1400° C) by a Fourier Transform (FT) filtering of the high r correlations in the RDF. This corresponds to the disappearance in the interference function of high frequency contributions. We performed the FT on the theoretically calculated S (q), we suppressed the high r contributions (r > 5 Å) and then we made a back FT calculation on the filtered real space contribution (0 → 5 Å). Figure 5.5.9 presents the results for the αmodel. The process smooth high frequencies contributions and the agreement with the experiment at high q values is better. Unfortunately, the FSDP is broadened by this process and the agreement in the low q (q < 4 Å−1) domain is deteriorated. We expect that this procedure will simulate a too drastic increase of the structural disorder beyond 5 Å. The smoothing process simulates an abrupt decrease of the structural order beyond 5 Å around each of the 448 atoms of the model. From these results, we can conclude that the experimental spectrum is intermediate between those of non-smoothed model and smoothed model the sizes of which (1 nm) are close to those of local order models obtained from EXAFS measurements. It is interesting to note that for the β-model (1500°C), all the small diffraction peaks appear on the experimental spectrum of the sample annealed at 1550°C.

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Figure 5.5.9 Calculated neutron interference functions (Sn(q)) for the α-model compared with the experimental spectra (1400°C). The Fourier filtered (FF) calculated spectrum is given as well as the crude one for the cluster of α-Si3N4.

Figure 5.5.10 Calculated neutron interference functions (Sn(q)) for the β-model compared with the experimental spectra (1500°C and 1550°C). The calculated spectrum for the cluster of β-Si3N4 is also given. Figure 5.5.11 presents the calculated and experimental RDF. The evolution of the MRO is reproduced by the calculations. X-RAY The structure factors SX(q) of the models compared to the experimental ones are presented in Figure 5.5.12. The Fourier filtered theoretical spectra are also given (the correlations at r ≥

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Figure 5.5.11 Calculated and experimental neutron RDF. The disappearance of the correlation at 3.9 Å is well reproduced by the modification of the stacking mode (α → β) when going from 1400°C to 1500°C. 4 Å for the α-model and r ≥ 3.8 Å for the β-model are removed). As for the neutron calculated spectra, the Xray calculated spectra exhibit a lot of small sharp structures relative to long-range interatomic correlations in the models, but contrary to neutron, experimental spectra exhibit the signatures of long-range correlations in the high q domain (q > 5Å−1). These correlations are pretty well reproduced by the α-model (1400°C). The experimental spectra are, on the whole q range, intermediate between the theoretical smoothed and theoretical non-smoothed spectra. The differences concerning the low q domain are also reproduced by the calculation: the X-ray spectra do not exhibit a FSDP. The main result of these calculations is the existence of nano-ordered structures α- and β-SiCN in these powders. The structure change when the annealing temperature increases and the modifications of local order are accompanied by ‘phases transition’ which occur in the amorphous state. This result underlines the importance of the existence of the β-SiCN (β-Si3N4 substituted and relaxed) phase in the nucleation of β-SiC crystallite. The nucleation of β-SiC is beginning within the β-SiCN phase at 1500°C. For intermediate compositions (HMDS43 and HMDS44), the C content does not allow the local formation of β-SiC nuclei at 1500°C. More energy is needed to allow the diffusion of C atoms in the β-SiCN structure to reach locally the SiC stoichiometry. At 1500°C, for the rich C powders (HMDS40 and HMDS42), the β-SiCN structure is saturated by C atoms; locally one finds SiC4 tetrahedra and nuclei of β-SiC that are formed. The excess C forms a free C phase. For the N-rich powder

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Figure 5.5.12 (a) Calculated X-ray interference functions (SX(q)) for the α-model compared with the experimental spectra (1400°C), (b) calculated X-ray interference functions (SX(q)) for the β-model compared with the experimental spectra (1500°C). The Fourier filtered (FF) spectrum is also given. (HMDS45), there is not enough C in the structure to prevent the formation of nano-crystallites of β-Si3N4. References Debye, P. (1915) Ann. Physik, 46, 809. Dixmier, J. (1992) J. Phys. I France, 2, 1011. Dixmier, J., Bellissent, R., Balhoul, D. and Goursat, P. (1994) J. Eur. Ceram. Soc., 13, 293. Dürr, J., Schempp, S., Lamparter, P., Bill, J. and Aldinger, F. (1999) Precursor-derived Ceramics (eds J. Bill, F. Wakai and F. Aldinger), Wiley-VCH, p. 224. Hajdu, F. (1971) Acta Crystallogr., A27, 73.

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Page 226 Mayo, S. L., Olafson, B. D. and Goddard III, W. A. (1990) J. Phys. Chem., 94, 8897; (1992) Neutron News, 3 (3), 29–37. Placzeck, G. (1952) Phys. Rev., 86(3), 377. Sazaki, S. (1984) KEK report, anomalous scattering factors for synchrotron radiation users, calculated using Cromer and Liberman’s method, National laboratory for high energy physics. www-lib.KEK.jp/cgi-bin/ img_index?198324022. Waasmaier, D. and Kirfel, A. (1995) Acta Crystallogr., A51, 416. Warren, B. E. (1969) X-ray Diffraction, Addison-Wesley Publishing Company. Waseda, Y. (1984) Lecture notes in physics, novel application of anomalous X-ray scattering for structural characterization of disordered materials, Springer Verlag, Berlin Heidelberg, New York, Tokyo.

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Page 227 6 Conductivity and dielectric properties of Si/C nanopowders Abdelhadi Kassiba 6.1 Introduction Nano-sized powders of silicon carbide (SiC) have recently attracted a great deal of interest mainly oriented towards the improvement of mechanical and thermal properties of nanocomposite materials such as MgO— SiC, Al2O3—SiC (Niihara and Nakahira 1991) and oxide matrices reinforced by SiC fibres (Karlin and Colomban 1998; Monthioux and Delverdier 1996). The nanometric SiC grains are obtained by pyrolysis of polycarbosilane polymers in inert gas (Yajima 1980) or by laser pyrolysis of (SiH4, C2H2) gaseous mixture (Cauchetier et al. 1988). Whatever the used process, the final network contains carbon (C) in excess arranged as isolated clusters or associated with the SiC grains (Charpentier et al. 1999, Monthioux, Chapter 4). Several reports have pointed out the correlation between the structural, mechanical and electrical properties of SiC fibres with the C in excess content (Karlin and Colomban 1998; Mouchon and Colomban 1996; Monthioux and Delverdier 1996). Particularly, the enhanced thermal stability of the ceramics and their improved mechanical features as well as the electrical conductivity of SiC-based ceramics depend on C excess, which prevails in materials annealed at temperatures below 1400°C. Electron paramagnetic resonance (EPR) investigations (Kassiba, Chapter 5.2) on the SiC nanopowders, reveal a high content of delocalised unpaired spins (1019 spin g−1) thermally activated by low energies (~10 meV) (Charpentier et al. 1999). Furthermore, a drastic damping of an EPR λ/2-resonant cavity was found in these materials when annealed above 1400°C. Dielectric losses by conducting paths modify the quality factor of the cavity and the microwave absorption rate. All experimental investigations (Charpentier et al. 1999; Kassiba et al. 2000) have demonstrated that annealing alters irreversibly all the physical properties including the grain morphology, size and composition as well as the crystalline structure and interface features. The latter which are dominated by the C-rich phase at the outermost particle surfaces and the disorder are expected to play an important role in the conduction process and dielectric behaviour. Moreover, we have demonstrated earlier (Kityk et al. 2000), that SiC nanocrystallites embedded within a polymer matrix show promising non-linear optical phenomena. In that case, the interface SiC-polymer is of particular importance on the induced second harmonic generation signal. This can open new areas through the search of original electronic and non-linear optical properties in materials based on SiC nanocrystallites. The following report is dedicated to the analysis of the dielectric behaviour, electrical and transport properties in SiC nanopowders with regard to the structure, grain sizes and interface features. Different experimental approaches are conjugated for a better understanding of the electrical properties. Among them, we have used dielectric relaxation spectroscopy

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Page 228 (DRS), microwave method (MWC), 4-points method and Hall effect measurements. These investigations are carried out as functions of annealing which modulate in large extent the physical properties of the network. Measurements with variable temperatures (150, 570 K) and frequencies (0.1 Hz, 10 GHz) are achieved for a quantitative understanding of the electronic, relaxation and transport phenomena in these SiC nanopowders. 6.2 Material specifications SiC nanopowders are synthesised by CO2 laser pyrolysis of gaseous mixture (SiH4, C2H2) (Cauchetier et al., 1988, Chapter 2). The rate of C in excess is monitored by the gas fluxes while the residence time of the reactants in the laser beam monitors the particle diameters in the as-formed powder. The composition, the SiC crystalline structures and the particle sizes can be modulated to a large extent by annealing under argon up to 1800°C (Charpentier et al. 1999). The mean particle diameter varies from ~30 nm (annealing temperature 1200°C) to the micrometer range (annealing at 1800°C). Transmission electron microscopy (TEM) investigations (Monthioux, Chapter 4) have shown that the grains are constituted by SiC polytypes partially covered by thin layers of C in one batch SiC177 (C in excess 10%) or by an oxide in SiC212 (C in excess 4%). Structural investigations import on the involved SiC polytypes which are essentially 6H and 3C polytypes with annealing-dependent ratios. Amorphous SiC is also present in a minor concentration (≤10%) as well as sp2 C for annealing below 1500°C in SiC212 and remains in SiC177 annealed up to 1800° C. Whatever the C in excess rate (SiC163:1.11, SiC173:1.10 and SiC212:1.04), the batches exhibit similar electrical behaviour at the same annealing stages. So, the following report is limited to the SiC212 batch with the powders referred below as SiC212/1200°C and SiC212/1800°C, following their annealing temperatures at 1200°C and 1800°C. 6.3 Experimental results and interpretation 6.3.1 Dielectric losses and electrical conductivity in the microwave range (MWC) These measurements are of a particular interest since neither ohmic contacts nor a particular sample shape processing is required. A simple diagnostic of the conducting states in the 10 GHz domain can be obtained easily. This requires an EPR spectrometer with a λ/2-resonant cavity in order to evaluate the dielectric losses in the SiC nanopowders. The mechanism is based on a damping of the resonant cavity and lowering its quality factor by conducting materials. The resonant cavity design is devoted to enhance these effects by using a small amount of the SiC powder (~10 mg) located in the cavity region where the microwave electric field is maximum. In all the investigated SiC batches, Ta = 1500°C corresponds to the critical temperature where the dielectric losses are manifested and alters significantly the microwave response of the materials. Figure 6.1 shows the behaviour of the reflected microwave power by the SiC nanopowders annealed between 1200°C and 1800°C. Narrow dips correspond to a minimum of the reflected power realised in the samples annealed below 1500°C. Conversely, flat and broad dips correspond to the materials annealed between 1500° C and 1650°C. This behaviour is consecutive to the occurrence of macroscopic conduction induced by annealing at T ≥ Ta.

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Figure 6.1 Reflected microwave power by SiC212 powder located in the region of an X-band (ω0 = 9.5 GHz) EPR resonant cavity where the microwave electric field is maximum. The dip shapes represented from bottom to top are due to the samples SiC212 annealed at 1200°C, 1300°C, 1400°C, 1500°C, 1550°C, 1600°C, 1650° C, 1700°C and 1800°C. The maximum of dielectric losses (flat dip) is realised in the SiC212/1550°C sample.

Figure 6.2 Electrical conductivity in the microwave range v. the sample temperature in SiC212 nanopowders annealed at 1500° C(▲), 1650°C(●) and 1800°C(■); the lines are a guide for the eye. Quantitative measurements of the electric conductivity by using the process developed by J. G. Theobald in Besançon University (Frand et al. 1995; Kassiba et al. 1999) are shown in Figure 6.2. The conductivity increases with the sample temperature and follows qualitatively an activated law: with [σ0(Sm−1), Eg(meV)] ≈ [1.1, 20], [0.7, 6], [0.9, 9] respectively for annealing at 1500°C, 1650°C and 1800°C.

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Figure 6.3 Electrical conductivity measurements in SiC212 by the conventional 4-points method. The evolution with the temperature shows a thermal activation of the conductivity with energies, 3 meV below 185 K and 30 meV above. Then, it seems that a substantial electrical conductivity is induced by annealing above 1400°C with a thermal activation of the charge carrier motions by low energies 6–20 meV. These values are in agreement with those required to delocalise some paramagnetic species (Kassiba, Chapter 5.2). 6.3.2 Electrical conductivity by the 4-points method The samples are prepared as pressed pellets (diameter 12 mm and 0.5 mm thick). On each pellet, four gold contact points are deposited by CVD technique. Annealing under hydrogen at 300°C during 5 min improves the contact resistors which change from 1000 to 40 Ω and remain ohmic on all the explored temperature range. Figure 6.3 shows the dc conductivity v. the sample temperature in the SiC212/1650°C sample. Above 185 K, a thermally activated conductivity is found, that is, σdc(T) = σ0 exp(−Eg/kT) with Eg ~ 30 meV and a pre-factor σ ≈ 60 Sm−1. The activation energy is in agreement with those obtained from the above microwave characterisations (MWC). However, a discrepancy is found in the absolute conductivity values. The main origins are plausibly due to the shape of the investigated materials, the different porosity degree and the inhomogeneous conducting paths. Furthermore, transport measurements by the Hall effect are carried out in order to determine the nature and concentration of the charge carriers as well as their mobility. n-type conduction, charge carriers concentration N ~ 1019 cm−3 and a mobility μ ~ 1cm2 V−1 s−1 are the main parameters of the electronic transport in these SiC nanomaterials. 6.3.3 Dielectric relaxation spectroscopy (DRS) Theoretical background and phenomenology DRS determines the relative complex permittivity ε*(ω, T) = ε′(ω, T) − iε″(ω, T) v. the alternative applied voltage frequency (ω/2π) and the sample temperature T. The overall

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Page 231 complex conductivity is deduced from the relation: σ*(ω, T) = iωε0ε*(ω, T), where ε0 is the vacuum permittivity. As far as the forthcoming experimental facts are concerned, the real conductivity part can be set in the form of contributions from the steady-state (σdc) conductivity and a frequency-dependent part as follows: However, phenomenological representations are usually employed to account for the frequency behaviour of ε*(ω) in heterogeneous media (Hunt 1998). Among them, the Havriliak Nagami (HN) response (Havriliak and Nagami 1966), which reads:

Δε represents the difference between the low frequency limit, that is, the relaxed permittivity (εs) and the high frequency limit ε∞ with respect to the explored frequency range. The phenomenological parameters α, β monitor the asymptotic behaviour of ε*(ω). Indeed, when ωτ 1; ε″ ~ ω−αβ and for ωτ 1; ε″ ~ωα, τ, represents also a phenomenological relaxation time (Hunt 1998) even if it can be associated to the experimental dielectric loss peak position. Experimental details Novocontrol Broadband Dielectric Spectrometer is well adapted for a class of materials with a conductivity lying between 1 Scm−1 to that of high insulating materials. Measurements of the dielectric function are possible over a wide frequency range (0.1 Hz–1 GHz) by using two experimental configurations. From 0.1 Hz up to 10 MHz, a solartron Si1260 combined with a broadband dielectric converter (BDC) determines the overall sample impedance. In the high frequency limit (1 MHz–1 GHz) an HP4291 radio frequency impedance analyser is used with a precision coaxial line (Kremer and Arndt 1997). For both configurations, the sample cell consists of two round golden plates electrodes filled with the material (pressed pellet with diameter 12 mm and 0.5 mm thick) to form a capacitor. In the high frequency range, the sample capacitor is used as a termination of a golden coaxial line. The corresponding impedance is deduced from the complex reflection factor at the end of the line. The measurements were performed as a function of temperature between 198 and 573 K. A gas stream of nitrogen (N) is used and the measured temperature close to the sample is controlled within the accuracy of ±0.1K. Dielectric behaviour Due to the critical annealing temperature Ta = 1500°C, which alters significantly all the physical properties of the SiC nanopowders, we have examined in detail the conduction and dielectric behaviour of two representative samples, namely SiC212/1400°C and SiC212/1700°C. In the former, the absence of pertinent details on the dielectric functions, above 10 MHz, allows us to limit the measurements to the frequency range 0.1 Hz, 10 MHz. Figure 6.4(a) and (b) shows in a log–log scale, the frequency and thermal evolution of the dielectric function [ε′(ω), ε″(ω)]. The continuous curves are plotted from HN response with the parameters

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Page 232 (α1400 = 1; β1400 = 0.5). α and β are independent of the sample temperature (Ts) as well as ε∞ = 8. Following the frequency range, two variation regimes of ε″(ω) are evidenced. In the low frequency side, the linear variation occurs from the steady-state dc conductivity. The second regime, in the high frequency side, is marked by a dielectric loss peak centred at a frequency νmax = 1/τmax, which varies from ~10 Hz to ~1 MHz with increasing Ts. The real dielectric function ε′(ω) is also marked by a relaxation process with the frequency νmax and by a low frequency dispersion (LFD). The latter behaviour depends on the experimental conditions and testifies the existence of contributions from the contacts between the sample and the electrodes. So, in the adjustment of ε′(ω) by the (HN) equation, we have ignored the (LFD) contribution. The relaxed parameter 140 is one order of magnitude larger than the value in bulk crystalline SiC. This can be due to an interfacial polarisation favoured by the high specific surface of the SiC212/1400°C sample and by the inhomogeneous particle surface composition through the predominance of C in excess.

Figure 6.4 Log–log representation of the (a) real and (b) imaginary dielectric function v. the frequency in the sample annealed at 1400°C. The continuous lines are fits from HN equation. Similarly (c) and (d) represent the dielectric function in the SiC212/1700°C sample.

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Similar investigations were performed on the SiC212/1700°C sample (Figure 6.4(c) and (d)). However, higher frequencies (up to 1 GHz) are needed in order to characterise the dielectric loss peak as in the SiC212/1400° C sample. The phenomenological parameters are (α1700 = 0.5; β1700 = 1) and ε∞ = 11, independent of Ts. The relaxation process is marked by a characteristic frequency (νmax) in the order of 100 MHz and increases is quite lower than towards high frequencies with increasing Ts. The relaxed dielectric constant, in the SiC212/1400°C sample. This is connected with the grain surfaces being striped from the C in excess (Charpentier et al. 1999) and with a coalescence of SiC particles, that is, mean diameter 1600 nm for annealing at 1700°C, limiting the interface polarisation effects which prevail in the SiC212/1400°C sample.

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Figure 6.5 Log–log representation of the conductivity v. the frequency obtained by DRS spectroscopy: (a) SiC212/1400°C and (b) SiC212/1700°C. Electrical conductivity In SiC212/1400°C and SiC212/1700°C samples, the frequency dependence of the conductivity are depicted in Figures 6.5(a) and (b) where the continuous lines are drawn from the (HN) equation. We may note that the behaviour at low frequencies is dominated by σdc up to crossover values about 1 kHz and 10 MHz respectively in the SiC212/1400°C and SiC212/1700° C samples. These values are temperature dependent and follow qualitatively the behaviour of the dielectric loss peak positions. In the high frequency side, no universal behaviour such as σac(ω) α ωp with 0 < p < 1 was identified (Jonscher 1977; Funke 1993). This excludes any speculation on the real nature of the conduction process in these heterogeneous media. The thermal behaviour of the dc conductivity are summarised in Figures 6.6(a) and (b). In the SiC212/1400° C sample, an activation law σdc(Ts) = σ0 exp(−Eg/kTs) with Eg = 500 meV and a pre-factor σ0 = 7Sm−1 are realised above T1 = 330 K. Similarly, in the

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Figure 6.6 Dc conductivity v. 1000/T obtained from DRS spectroscopy measurements. Activated conductivity following Arrhenius law is realised in a limited temperature range. SiC212/1700°C sample, by limiting the temperature range to 250 ≤ Ts ≤ 430 K,σ(Ts) is also accounted by an activated process with an meV and a pre-factor σ0 = 0.3Sm−1. Below 250 K, the thermal variation of the conductivity is characterised by a continuous decrease of the activation energy without any defined variation law such as that relevant in a variable range hopping process (Mott and Davis 1979) or that used in SiC fibres (Chauvet et al. 1992). Above 430 K, as shown in Figure 6.6(b), the conductivity exhibits a rounded shape behaviour (negative temperature coefficient). This process is reversible but it is no longer connected with a metallic character of the SiC212/1700°C sample (Kassiba et al. 2000). The observed behaviour has to be analysed with respect to the conductivity supports in the material, that is, charge carriers diffusion in the crystallite bulk, tunnelling through the interface barriers as well as the material porosity. All these parameters have to be optimised for relevant insights on the particular behaviour around 430 K observed in all the SiC nanomaterials annealed above 1400°C. It is worth noticing the agreement between the absolute conductivity values and the activation energies obtained by DRS and MWC methods even if different sample processing (pellet, powder) are used. With regard to the relaxation process, the parameter νmax varies with the sample temperature as the inverse of the dc conductivity (Figures 6.7(a) and (b)). From the experimental value of νmax and limiting the temperature range to that where an Arrhenius law is realised, we have established the following relations:

and

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Figure 6.7 Representation of the static conductivity v. the relaxation time of the HN response in the samples (a) SiC212/1400°C and (b) SiC212/1700°C. This proportionality is quite general and encountered both in electronic (Hunt 1993; Pelster et al. 1994) and in ion conducting systems (Macedo et al. 1972; Moynihan 1994) as well. No physical supports have been clearly assigned to this conductivity-relaxation dependence as relevant in transport phenomena dominated by surface or bulk effects. On the other hand, EPR investigations (Charpentier et al. 1999) have evidenced charged vacancies, in the SiC crystallites, submitted to a thermal delocalisation by energies 13 meV. Qualitatively, this value can be looked as the binding energy of individual charges to their local sites and constitute a low contribution to the overall activation energy associated with the macroscopic conductivity. The main part comes from the energy barriers at the grain boundaries. The exact mechanism cannot be drawn rigorously but it is worth noting that


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Page 237 the large difference in the activation energies 500 meV and 35 meV respectively in the SiC212/1400°C and SiC212/1700°C samples agrees with that between the specific surfaces (SiC212/1400°C : 35 m2 g−1 and SiC212/1700° C : 1.2m2 g−1). This analysis is supported by the magnitude of the static dielectric constant being enhanced in materials with high specific surfaces due to the interface polarisation effects. Then, surface effects at the grain boundaries monitor the electrical behaviour of these powders and the conductivity-relaxation dependence can be attributed to such effects. 6.4 Conclusion We have used complementary techniques in order to quantify the electrical, dielectric and transport properties in SiC nanopowders. A thermally activated conductivity is pointed out in the materials annealed at T ≥ Ta = 1500°C. At this stage, a removal of the C in excess is operated as well as a coalescence of the SiC particle. These modifications favour the development of conduction channels through the increase of the mean free path of the charge carriers and the lowering of the barrier thickness between the grains. The interface effects seems crucial and explain the conductivity and dielectric behaviour in these media. The interpretation correlates the dielectric properties with the structural and morphological changes by annealing. It is worth noting also the possibility to modulate the electrical behaviour of SiC nanopowders as well as their effective dielectric constant by suitable thermal treatments. References Cauchetier, M., Croix, O. and Luce, M. (1988) Adv. Ceram. Mater., 3(6), 147. Charpentier, S., Kassiba, A., Bulou, A., Monthioux, M. and Cauchetier, M. (1999) Europ. Phys. J: AP, 8, 111. Charpentier, S, Kassiba, A., Emery, J. and Cauchetier, M. (1999) J. Phys. Cond. Matter, 11, 4887. Chauvet, O., Stoto, T. and Zuppiroli, L. (1992) Phys. Rev. B, 46(13), 8139. Frand, G., Bohnke, O., Lacorre, P., Fourquet, J. L., Carré, A., Eid, B., Theobald, J. G. and Gire, A. (1995) J. Solid State Chem., 120, 157. Funke, K. (1993) Prog. in Solid State Chem., 22(2), 111. Havriliak, S. and Nagami, S. (1966) J. Polym. Symp., 14, 89. Jonscher, A. K. (1977) Nature, 267, 673. Karlin, S. and Colomban, Ph. (1998) Composites, 29B, 41. Kassiba, A, Tabellout, M., Charpentier, S., Herlin, N. and Emery, J. R. (2000) Solid State Comm., 115, 389. Kityk, I. V., Makowska, M., Kassiba, A. and Plucinski, K. J. (2000) Opt. Mater., 13, 449. Kremer, F. and Arndt, M. (1997) Dielectric Spectroscopy of Polymeric Materials, Ch. 2 (eds J. P. Runt and J. J. Fitzgerald), American Ceramic Society. Macedo, P. B., Moynihan, C. T. and Bose, R. (1972) Phys. Chem. Glasses, 13(6), 171. Monthioux, M. and Delverdier, O. (1996) J. Eur. Ceram. Soc., 16, 721. Mouchon, E. and Colomban, Ph. (1996) J. Mater. Sci., 31, 323. Mott, N. F. and Davis, E. A. (1979) Electronic Process in Non-Crystalline Materials, 2 edn. Clarendon Press, Oxford. Moynihan, C. T. (1994) J. Non-Cryst. Solids, 1395, 172. Niihara, K. and Nakahira, A. (1991) Ann. Chim. Fr., 16, 479. Pelster, R., Nimtz, G. and Wessling, B. (1994) 49(18), 12718. Yajima, S. (1980) Philos. Trans. R. Soc. London Ser., A 294, 419.

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Page 238 7 Processing and tailoring of Si/C/N-based nanocomposites Paul Goursat, Djamila Bahloul-Hourlier, Benoît Doucey and Martine Mayne 7.1 Introduction Silicon nitride exists in two structural forms with a hexagonal unit cell α-Si3N4 = (space group P31c) (Lang 1977) and β-Si3N4 (space group P63/m) (Grün 1979). The covalent nature of Si—N bonds provides very good mechanical properties, such as high strength and high toughness up to elevated temperatures. The sintering of a powdered sample corresponds to the formation of necks between grains due to the presence of evaporation–condensation and/or diffusion mechanisms with an atomic or molecular redistribution. In the case of Si3N4 the self-diffusion for silicon (Si) or nitrogen (N) is very low in the tridimensional network of SiN covalent bonds, which explains the difficulties encountered to densify this material. Moreover, decomposition occurs at high temperature which blocks the increase in mobility of the species with rising temperature. These difficulties are usually partially solved in other systems by using nanosized powders, but in the case of Si3N4, it is necessary to add a driving force – mechanical and/or gas pressure, sintering aids, etc. – which explains the different types of processes and materials: reaction bonded silicon nitride (RBSN); hot-pressed silicon nitride (HPSN); hot isostatically pressed silicon nitride (HIPSN); gas pressured silicon nitride (GPSN). The additives used for the sintering of silicon nitride (MgO, Al2O3, Y2O3, etc.) react with silica always present at the surface of the grains and lead to the formation of an oxynitrided eutectic phase with the general formula MSiM′ON, where M and M′ are metallic cations. These vitreous phases, which have been studied in quaternary or quinary systems (Figure 7.1), play a very important role during sintering and on the properties at high temperature of the sintered material (Hampshire et al. 1978). Experiments on these glasses and, in particular, the vitreous domain show that the N content changes with the metallic cations and the temperature range (Lewis 1993). Among the main characteristics of these materials it has been shown (Hampshire et al. 1994) that the viscosity increases with N content which explains that crystallisation is slower than for the corresponding oxide glasses and that after sintering, the thermodynamic equilibrium is not reached. 7.2 Sintering of silicon nitride with additives A lot of papers are devoted to the sintering of silicon nitride for modelling, the determination of mechanisms or the development of microstructures (Ziegler et al. 1987). During the heat treatment of the powdered sample (α-Si3N4+additives : Al2O3, Y2O3, etc.), the oxynitrided eutectic flows in the porosity and induces capillary forces, dissolution of fine particles and

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Figure 7.1 YSiAlON system.

Figure 7.2 Shrinkage and α → β phase transformation v. time. a rearrangement process. The shrinkage is usually described by three different overlapping mechanisms (Figure 7.2). 7.2.1 Particle rearrangement The liquid phase wets the α-Si3N4 grains and it localises mainly at the contact points. Under the forces of capillary action, a rapid densification is obtained by grain sliding. In the case of HP or HIP, the densification rate is sharply increased, the applied mechanical pressure being much more higher. The densification rate changes with the particle size, the viscosity and the volume fraction of the liquid. Different modelling emphasise the influence of these parameters. 7.2.2 Dissolution–Diffusion–Reprecipitation The applied pressure induces a preferential dissolution at the contact points of α-Si3N4 crystals in the liquid phase with concentration gradients. A kinetic regime is established with an intergranular migration of the dissolved species and a reprecipitation of β-Si3N4 crystals or the β′-SiAlON solid solution when alumina is used as a sintering additive. The

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Page 240 flow of mobile species explains the redistribution of matter and a macroscopic shrinkage. The comparison of the rate (dτ/dt) of the α → β transformation and the rate of densification (dρ/dt) shows that the ratio

corresponds to a diffusion regime. In the case of hot-pressing (Bowen et al. 1978), the rate of densification is given by the following equation which is derived from the Coble model established for the creep of materials:

(7.1) where P is the applied mechanical pressure; Db the diffusion coefficient in grain boundaries; Ω the volume of the mobile species; W the thickness of grain boundaries; G the diameter of the crystals; K is the Boltzmann constant. 7.2.3 Grain growth When soluble particles with different particle sizes (radius r1, r2) are in the presence of a liquid, the difference in solubility provides the driving force for grain growth which corresponds to the Ostwald ripening phenomena. The grain growth rate (Greenwood 1956) has been formulated as:

(7.2) C∞ is the equilibrium concentration for a planar surface, γ the surface energy, Ω the molecular volume of the mobile species, T the absolute temperature, K the Boltzmann constant, D the diffusion coefficient in the liquid phase, d the mean distance between particles 1 and 2. The grain growth kinetics by liquid phase sintering usually follows the equation: (7.3) where is the mean particle radius, n = 3 the growth exponent, k the growth rate constant and t is time. As a matter of fact, for the sintering of silicon nitride, a lot of parameters are to be considered and the equations are more complex. The presence of acicular crystals of β-Si3N4 in the starting powder hinders the rearrangement mechanism which overlaps the diffusion stage. It is better to use α-Si3N4 with an equiaxed morphology from the kinetics point of view. The content and the size of β-Si3N4 grains in the starting powder also play an important role during the 2nd and 3rd densification stage. β-Si3N4 crystals can act as the nuclei during the reprecipitation process if their size is higher than the size of nuclei issued from the homogeneous precipitation. According to Kang and Han (1995), the growth of β grains along the c axis is due to the different roughness of the liquid–solid interface at an atomic scale of (100) and (001) planes. If the number of β-Si3N4 crystals is low, the free grain growth of large crystals induces a microstructure with a bimodal particle size distribution. It is possible to control and adjust the microstructure of sintered materials with a given size or aspect ratio of β grains by adding β nuclei (Lee et al. 1994) or with a specific post sintering heat treatment (Rossignol 1994).

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Page 241 The viscosity of the liquid phase also influencing the sintering rate is also very important. For the same content of liquid phase, impurities such as oxides of Ca, Mg, Na, Fe, etc. which lower the viscosity, facilitate the rearrangement and the diffusion processes. In return, other elements such as carbon (C) have a deleterious effect. 7.3 Sintering of silicon carbide Silicon carbide is usually densified by solid-state sintering with aluminium, boron and C as sintering additives (Greskowich et al. 1977). Recently it was shown that silicon carbide can be densified at 1800–1900°C with a liquid phase containing Y2O3 and Al2O3 (Kleebe 1992). The β → α-SiC transformation is observed with a growth of elongated α-SiC grains. It is also possible for SiC to design microstructures for specific properties. 7.4 Si3N4/SiC nanocomposites In recent years, much interest has been focused on nanophase materials, which, due to their small grain size, can offer improved properties compared to traditional materials (Siegel 1994). A further advantage of these materials is their ability to be hot-formed because of their fine microstructures (Waikai 1990). The main objective of this work is to elaborate Si3N4/SiC nanocomposites in order to fabricate ceramic parts by hot-forming. The approach consists in introducing SiC (N) nanometric powders as a second phase into a Si3N4 matrix in order to influence the α-Si3N4 to β-Si3N4 phase transformation and the microstructural development of Si3N4/SiC nanocomposites. Usually nanocomposites are ordered in two groups (Niihara 1991): • micro–nano materials: a matrix with microscale crystals and an intragranular or an intergranular nanoscale second phase; • nano–nano materials: a matrix and a second phase with nanoscale crystals. These materials are fabricated according to three different routes. The sintering of mixtures of commercial silicon nitride with additives and a nanopowder of SiC. In the second route the nanoparticles are synthesised in situ during sintering by a reaction between silica and C deposited on silicon nitride crystals (Ishizaki and Yanai 1995). The last one corresponds to the sintering of a mixture of nanopowders: Si3N4, SiC and additives or Si/C/N and additives (Niihara et al. 1990). Till now, the number of papers devoted to the study of the sintering kinetics is very limited, in particular to determine the influence of nanoparticles on the different stages during the sintering of silicon nitride. 7.4.1 Dispersion and slip casting of Si/C/N nanopowders with Si3N4in organic liquids It is well known that the method by which dispersions are prepared is a critical issue in forming high quality green bodies and reliable ceramic parts. The present study concerns the fabrication process of such materials which must be optimised because of the different sizes of powders, and the hard agglomerates in Si/C/N nanopowders. The aim is to obtain a homogeneous distribution of the nanometric second phase in order to prevent the presence of defects (essentially nanopowder agglomerates) decreasing drastically the mechanical properties.

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Page 242 Si3N4 used as the matrix is the commercially available α phase Si3N4 powder (UBE-SNE 10). Si/C/N nanopowders used as second phase are synthesised by laser pyrolysis. Y2O3 (Lonza) and Al2O3 (Rhône Poulenc) are used as sintering aids. It has been observed by X-ray photoelectron spectroscopy (XPS), extended X-ray absorption fine structure (EXAFS) and X-ray near edge structure (XANES) that Si/C/N nanopowders obtained by laser pyrolysis are composed of mixed tetrahedra SiCxN4−x with a small amount of Si—O bonds (Gheorghiu et al. 1992; Ténégal et al. 1996). The surface of such powders, analysed by IR spectroscopy, contains Si—OH and Si—H bonds (Cauchetier et al. 1991). Preparation of the green bodies Two successive stages are necessary to prepare green bodies: the mixing of the different powders and the forming process. In order to avoid the reagglomeration phenomenon induced by dry mixing, the mixing process in liquid medium is well adapted. Moreover, the use of organic liquids is preferable to prevent hydrolysis of silicon nitride based powders. Four different organic liquids are tested: ethanol, acetone, acetophenone and hexane. Each powder (Si3N4, Si/C/N) is mixed in each liquid (concentration = 50 g/l), and the evaluation of the most suitable liquid is realised by sedimentation tests. A mechanical device and a dispersing agent are chosen in order to, respectively, break down agglomerates and stabilise the suspensions. Two mechanical devices are selected: the ‘turbula’ (system Schutz) allowing powder mixing by helicoidal movement, and the ultrasonic probe (Vibra Cell 72412, Bioblock Scientific) involving the breaking of hard agglomerates. Two dispersing agents (ICI Chemicals) having different nature are tested: the basic PS2 surfactant soluble in alcohol and the acid PS4 surfactant soluble in ketones. Each powder (Si3N4 or Si/C/N) is mixed using ‘turbula’ or ‘turbula’ + tultrasonic probe in ethanol + PS2 or acetophenone + PS4. The evaluation of the most suitable mechanical device and dispersing agent is achieved by sedimentation tests. Viscosity measurements coupled with sedimentation tests allow one to determine the optimum concentration of the well-adapted dispersing agent. After the homogenisation and the dispersion of suspensions, the liquid must be removed in order to obtain green bodies. The forming process selected is slip casting in a plaster mould because it allows to preserve the dispersed state of the suspension. Organic liquid medium Whatever the powder (Si3N4 or Si/C/N) and the organic liquid (ethanol, acetone, acetophenone, hexane), flocculation occurs in the suspensions. Sedimentation tests consisted of evaluating the relative height of sedimentation (hrelative = hsediment/htotal). In the case of flocculated suspensions, a low relative height indicates a compact stacking of powders due to a beneficial effect of the liquid on particle dispersion. The evolution of relative height of sedimentation in each liquid is presented for the two silicon nitride based powders (Figure 7.3). For the two powders, an increase of the relative height is observed as follows: ethanol, acetone, acetophenone, hexane. Such an evolution corresponds to the classification of the more adapted liquid to the less adapted. The suitability of each liquid for powders can be explained by acid–base interactions between the surface sites of the particles and liquid molecules. The surface of Si3N4 or Si/C/N powders is essentially composed of acid Si OH bonds and some basic Si2—NH and N—H bonds.

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Figure 7.3 Evolution of the relative sedimentation height for each powder in each liquid. The interactions liquid/powder result mainly from the attraction of basic liquid by acid Si—OH groups. Thus, the more basic liquid (acetophenone) should be the best medium for the dispersion of powders. Meanwhile, the adsorption of large molecules, such as acetophenone, on all Si OH sites is difficult because of the rigidity of Si—NH—Si bonds and also of the repulsion between the lone electron pair of N and the π orbitals of benzene group. That’s why the linear ethanol molecule, even though it is less basic than acetophenone, is easily adsorbed on Si3N4 and Si/C/N powders. Mechanical device for mixing and dispersion, and dispersing agent Whatever the liquid, flocculation occurs in the suspensions. So the use of a well-adapted mechanical device to disperse powders and of a dispersant to stabilise the suspensions is necessary. Sedimentation tests reveal that basic PS2 surfactant in ethanol does not have an effective action on powders, because for various dispersant contents, the suspensions remain flocculated. This is probably due to the low number of acid Si—OH sites available for the adsorption of basic PS2 molecules. Effectively, these sites should be occupated by the basic ethanol molecules. The suspensions dispersed with PS4 (5% for Si3N4, 10% for Si/C/N, wt% of the dry matter content) are deflocculated, indicating the beneficial effect of this surfactant. The evolution of the sediment height with time for Si3N4 and Si/C/N suspensions mixed with ‘turbula’ or ‘turbula’ + ultrasonic probe (Figure 7.4) indicates the positive effect of ultrasonic probe, suggesting that this device is necessary to break down agglomerates. The suitability of PS4 dispersant is linked to a good compatibility between acid surfactant molecules and basic surface sites of particles, non-occupated by the acetophenone molecules. Forming process For slip casting in plaster mould, suspensions should have high dry matter content (≥50 vol%) in order to eliminate all the liquid by capillarity. The studied slurry is composed

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Figure 7.4 Evolution of sedimentation height with time. T: Turbula, US: Ultrasonic Probe; ◆ Si3N4, ☐ Si/C/N (T), × Si3N4 (T + US), ○ Si/C/N(T + US). of a mixture containing of 71 wt% of Si3N4, 20 wt% of Si/C/N, 6 wt% of Y2O3, 3 wt% of Al2O3 and acetophenone. The dry matter content is 60 wt%. The dispersant content adapted to this slurry is determined by both sedimentation tests and viscosity measurements. For a 2.5 wt% of PS4, sedimentation tests indicate a sediment height near to zero during three days. This result gives evidence that the suspension is stabilised by 2.5 wt% of PS4 surfactant. Viscosity measurements as a function of dispersant content (Figure 7.5) reveal a minimum of viscosity around 2.5 wt% of PS4 which is in agreement with sedimentation tests. Slip casting of this slurry takes place over one or two days, and the density of the green bodies obtained is high (≈57% of the theoretical density) linked to a compact stacking of powder particles resulting from the well-dispersed suspension. This study has shown that compact green bodies can be produced in an organic liquid from a mixture of two silicon nitride based powders with different granulometries. Although the acetophenone is not the most convenient medium for the mixing process of these powders, the acetophenone/PS4 surfactant seems to be a good chemical couple to disperse suspensions. Thus, stable and well-dispersed slurries containing nanometric powders are obtained using an ultrasonic probe to break down hard agglomerates of nanopowders. Slip casting of such slurries leads to well-compacted green bodies associated with the homogeneous and dispersed state of the suspensions. 7.4.2 Sintering of micro–nanocomposites Powders (Si3N4, (UBE SN E10), Si/C/N) and additives (Y2O3 = 6 wt%, Al2O3 = 3 wt%) are deagglomerated in liquid media. After drying, the green bodies are hot-pressed (35 MPa) in the presence of N (1 atm) (Table 7.1). The theoretical density of the composites is calculated from the density of each constituent (Si3N4, SiC, glassphase) and by using the rule of mixtures.

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Figure 7.5 Viscosity as a function of dispersant content. Table 7.1 References of micro–nanocomposites Compounds (wt%) Theoretical SiC content (wt%) Grades Si3N4 SiCN29 SiCN35 SiC173 91 — — — 0 Si3N4 71 20 — — 1.7 Si3N4/SiCN29 (20) 79.9 — 11.1 — 5 Si3N4/SiCN35 (11.1) 68.8 — 22.2 — 10 Si3N4/SiCN35 (22.2) 46.6 — 44.4 — 20 Si3N4/SiCN35 (44.4) 71 — — 20 20 Si3N4/SiC (20) 51 — — 40 40 Si3N4/SiC (40) Sintering kinetics with a linear increase of temperature The shrinkage curves for the different compositions with the same content of nanopowders of SiC show a sigmoidal shape (Figure 7.6(a)). The degree of densification at 1700°C is close to 98%. On this figure is also represented the curve concerning the shrinkage of Si3N4.In order to evidence the different behaviours, the derived curves giving the rate of densification are shown in Figure 7.6(b). For monoliths and nanocomposites containing a SiCN second phase, the beginning of the shrinkage corresponds to the formation of the oxynitrided eutectic in the YSiAlON system at about 1350°C. The content, the composition and the properties, particularly the viscosity of the liquid phase change strongly with temperature. Up to 1470°C with a high viscosity, the rate of densification corresponding to a rearrangement process and a dissolution of small particles is low. The shift to high temperatures for the Si3N4/SiC

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Figure 7.6 Isobar densification curves. (a) Relative density, (b) rate of densification. ○ Si3N4, ▲ Si3N4/SiCN29 (20), × Si3N4/SiCN35(22, 2), □ Si3N4/SiC(20). mixtures is explained by the SiC nanopowders. In this case, free C can react during the heat treatment with oxygen at the surface of Si3N4 grains and/or the additives. This reaction delays the formation of the liquid phase and modifies its quantity or composition and therefore the viscosity (Hampshire et al. 1978). Furthermore, SiC particles and agglomerates being insoluble in the liquid phase, the dissolution stage is suppressed. At 1470°C, the sharp increase in the rate of densification is explained by the dissolution of Si3N4 and SiCN particles which initiates the α → β phase transformation. These results show that at the file:///C|/Users/steppenwolf/Documents/educatio...ks/to%20burn/nanocomposites/files/page_246.html (1 of 2) [20/01/2010 10:05:35 AM]

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Page 247 opposite situation above 1550°C. The contributions of the rearrangement and dissolution processes to the overall densification change with the size of the crystals and the thickness of the intergranular vitreous phase (see equations 7.1 and 7.2). For nanopowders, at the same liquid phase content the large specific surface area and the presence of agglomerates induce a decrease of the thickness of grain boundaries and the flow of the species which explains the lower rate of densification at the beginning of sintering for Si3N4/SiCN composites compared to Si3N4 monoliths. At higher temperatures (T > 1550°C), the increase for composites is explained by a faster dissolution of SiCN nanoparticles (see equation 7.2). One notices that the temperature corresponding to the end of densification depends on the SiC content in the final material. The rate of densification during the dissolution– reprecipitation stage is governed by the thickness of grain boundaries (see equation 7.1). The thickness diminishes when the content of intergranular SiC particles increases. The flow of mobile species is impeded which explains the shift to higher temperatures. The same phenomenon is observed for Si3N4/SiC composites. Moreover, the starting SiC powder contains an excess of free C which reacts with oxygen from silica or amorphous silicon oxynitride located at the surface of silicon nitride grains or from additives and modifies the viscosity and the liquid phase content. These experiments show clearly that the source of SiC nanoparticles plays an important role on the dissolution–reprecipitation process and consequently on the rate of densification. The SiCN nanoparticles are more soluble in the oxynitrided liquid phase than SiC nanoparticles that are practically inert in this temperature range (Kleebe 1992). Densification kinetics during isothermal treatments The densification curves (Figure 7.7(a)) for the various grades exhibit the same shape with a monotonous decrease of the rate of densification with time (Figure 7.7(b)). The differences are mainly due to the composition and the content of the second nanometric phase. These results are consistent with the previous experiments. At 1530°C, the oxynitrided eutectic is formed which explains the high rate of densification for the rearrangement stage. The rate increases with the content of SiCN nanoparticles. The dissolution of these particles starts at the beginning of the rearrangement when the pressure is applied. The overall rate is faster than for Si3N4 alone. On the contrary, for Si3N4/SiC composites the shrinkage is low due to the inertness of SiC nanoparticles towards the eutectic. Moreover, SiC precipitates hinder the diffusion of species in grain boundaries which limits the densification whatever the starting nanopowder used. The densification for composites is uncompleted after 2 h at 1530°C under a pressure of 35 MPa. Composites microstructure It is known that it is possible to modify or adjust the microstructure of ceramic materials by introducing an inert second phase but little work is published on the Si3N4/SiC micro– nanosystem. Due to the fact that a lot of parameters influence the microstructure of silicon nitride, it is necessary to study carefully all the phenomena. First, the α → β phase transformation for Si3N4 and the crystallisation of β-Si3N4 from SiCN nanopowders are analysed, then the behaviour of mixtures is studied in order to determine the interactions and to localise SiC nanoprecipitates.

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Figure 7.7 Isothermal densification curves. (a) Relative density, (b) rate of densification. ○ Si3N4, ▲ Si3N4/ SiCN29(20), × Si3N4/SiCN35(22, 2), □ Si3N4/SiC(20). SEM OBSERVATIONS After sintering, the samples are polished and plasma etched (CF4/O2). Then, they are analysed in a plane perpendicular to the hot-pressing direction. Si3N4/SiCN29–Si3N4/SiCN35 COMPOSITES A fine microstructure composed of equiaxed grains, with a mean diameter close to 0.2 μm is observed when samples are heat treated at 1550°C. Then the aspect ratio increases with file:///C|/Users/steppenwolf/Documents/educatio...ks/to%20burn/nanocomposites/files/page_248.html (1 of 2) [20/01/2010 10:05:36 AM]

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Figure 7.8 Si3N4/SiCN29(20) micro–nanocomposite sintered (a) 1550°C, (b) 1650°C.

Figure 7.9 Si3N4/SiC35(22.2) micro–nanocomposite sintered at 1650°C. the sintering temperature (Figure 7.8). The comparison of Si3N4/SiCN29 and Si3N4/SiCN35 microstructures (Figure 7.9) indicates that the growth of β-Si3N4 crystals is hindered by higher SiC contents, but the agglomerates of nanograins are present. For these two grades the grain size is smaller than in the case of monolithic Si3N4 densified with the same experimental conditions and with the same additives (Figure 7.10). Si3N4/SiC173 COMPOSITES For these composites the influence of SiC nanograins on the development of the microstructure and the growth of β-Si3N4 crystals is detected, but with these nanopowders large agglomerates

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Figure 7.10 Monolithic Si3N4 sintered at 1650°C.

Figure 7.11 Si3N4/SiC micro–nanocomposite sintered at 1650°C. are revealed by a chemical attack (Figure 7.11). The presence of agglomerates is mainly due to the insolubility of SiC in the eutectic phase in the temperature range used for sintering. PHASE TRANSFORMATIONS The different powders have been heat treated according to the following conditions: T: treated at 1530°C, 2 h, under N2, 1 atm; TP: T + 35 MPa pressure; TPA: TP + additives(Y2O3 6 wt%, Al2O3 wt%). The α-,β-Si3N4 contents were determined by XRD experiments. The results are summarised in Table 7.2. The values for the α/β ratios after T and TP heat treatments are nearly equivalent and without additives no phase transformation is observed. On the contrary, when Y2O3 and Al2O3

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Page 251 Table 7.2 Influence of heat-treatment conditions on the densification and the rate of α → β transformation. T: treated at 1530°C/ 2 h/N2; TP: treated at 1530°C/2 h/N2/35 MPa; TPA: treated at 1530° C/2 h/N2/35 MPa/(+6%Y2O3 + 4%Al2O3) Starting α/β (wt%) α/β ratio after heat treatment (wt%) Relative density (%) TPA T TP TPA T TP TPA Si3N4UBE 100/0 100/0 70/30 0 0 30 / SiCN29 49/51 49/51 0/100 0 0 100 90 SiCN35 79/21 78/22 20/80 — — 74 63

Figure 7.12 β-Si3N4 content v. temperature. ○ Si3N4, ▲ Si3N4/SiCN29(20), × Si3N4/SiCN35(22, 2), ☐ Si3N4/ SiC(20). are added, crystals or SiCN particles are dissolved in the YSiAlONC eutectic and β-Si3N4 reprecipitates. The rate of the α-β transformation is higher for SiCN nanopowders than for Si3N4. Moreover, the densification decreases when SiC content is raised. In the case of powder mixtures (Si3N4, SiCN, additives), samples are hot-pressed under N at different temperatures during 2 h. The ratios α + β are compared with previous results corresponding to the individual constituents (Figure 7.12 and Table 7.3). The α → β transformation being thermally activated, the β-Si3N4 content increases with the sintering temperature for all the materials. When the same heat treatment conditions are used, the curves on Figure 7.13 show an increase of the α → β transformation rate with SiC content. For sintering temperatures lower than 1600°C and whatever the SiCN nanopowder content the nanocomposites contain more β-Si3N4 than the monoliths. Above 1600°C, the β-Si3N4 content is always higher in nanocomposites than for monoliths but the gap is less pronounced.

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Page 252 Table 7.3 α-Si3N4 content after sintering Materials Sintering temperatures (°C) Experimental results (%) α/α + β Theoretical ratio α/α + β Monolith Si3N4 1550 70 1600 30 1650 15 Si3N4/SiCN29(20) 1550 43 57 1600 27 24 1650 12 11 Si3N4/SiCN35(11.1) 1550 52 67 1600 22 28 1650 10 14 Si3N4/SiCN35(22.2) 1550 42 64 1600 24 26 1650 7 13 Si3N4/SiCN35(44.4) 1550 70 55 1600 15 21 1650 0 11 Si3N4/SiC(20) 1550 68 70 1600 15 30 1650 0 15 Si3N4/SiC(40) 1550 87 70 1600 53 30 1650 0 15

Figure 7.13 β-Si3N4/τ v. SiC content (T = 1650°C). These changes can be explained by the more important contribution to the phase transformation by nanopowders at low temperatures compared to that of microscale Si3N4. A very interesting result is obtained in the case of samples containing SiC nanopowders which are practically insoluble in the eutectic. An increase of the rate of the

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Page 253 α → β-Si3N4 transformation is clearly obtained and the transformation is completed at 1650°C. Conclusion The overall experiments show that the morphology, the composition and the content of the second phase play an important role on the α → β-Si3N4 transformation. The transformation rate changes with the SiC content in the SiCN starting powder. This effect is more pronounced with SiC nanopowders. Similar observations are mentioned in the literature, nevertheless, they are fragmentary and much debated. Sasaki et al. (1992) suggest that SiC nanoparticles act on the precipitation process as heterogeneous nuclei. However, this explanation is questioned by Pan et al. (1996) who show that for Si3N4/SiC nanocomposites SiC nanoprecipitates are not nuclei but only included by β-Si3N4 during the grain growth. The present work which concerns mixtures with SiCN or SiC shows that the nanopowders influence the germination and reprecipitation stage. Si3N4 nuclei are formed at the surface of SiC nanoparticles which lower the energetic barrier. It corresponds to a heterogeneous nucleation induced by SiC nanoparticles. Moreover, it has been shown (Kleebe 1992) that SiC is weakly soluble at very high temperature in the presence of YSiAlONC liquid phase. This very limited solubility is in favour of the reprecipitation of β-Si3N4 or β-SiAlON on SiC particles. The influence of SiC is more effective when its content increases from 10% to 40%, but the number of nuclei is too high for the transformation, and the main part is localised at grain boundaries. 7.4.3 Sintering of nanocomposites The study of the rheology of these nanopowders in different organic liquids and the nature of dispersing agents is in progress. Different grades of nanopowders, synthetised by laser pyrolysis are mixed with oxides (Y2O3, Ø 1 μm, 99.9%, Rhône-Poulenc; Al2O3, Ø 50 nm, 99.9%, Marketech International) in ethanol. The starting compositions and their references are given in Table 7.4. The powders are mixed and deagglomerated by milling in a polyethylene jar with Si3N4 balls. The mixtures are dried at 200°C and hot-pressed (35 MPa) in a atmosphere at various Table 7.4 References and starting compositions Grade Starting composition (wt%) Theoretical composition (wt%) Nanopowder Y2O3 Al2O3 SiCNO phase Y2O3 Al2O3 HSAl09 HSAl09 100 — — 97 — 3 HMDS35 + 3A HMDS35 97 — 3 97 — 3 SiCN35 + 6Y3A SiCN35 91 6 3 91 6 3 SiCN29 + 6Y3A SiCN29 91 6 3 91 6 3 HMDS35 + 6Y3A HMDS35 91 6 3 91 6 3 HSAl07 + 6Y HSAl07 94 6 — 90.2 6 3.8 HSAl15 + 6Y HSAl15 94 6 — 90.5 6 3.5 HSAl09 + 6Y HSAl09 94 6 — 91.2 6 2.8 SiCN35 + 9Y4.5A SiCN35 86.5 9 4.5 86.5 9 4.5 SiCN29 + 9Y4.5A SiCN29 86.5 9 4.5 86.5 9 4.5 HSAl09 + 6Y3A HSAl09 91 6 3 88.3 6 5.7

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Page 254 temperatures. After sintering, the different phases are determined by X-ray diffraction and the chemical composition is obtained by elemental analysis. The rate of densification is calculated by using the rule of mixtures and the theoretical density of the crystalline phases. Up to now, the number of papers devoted to sintering of silicon carbonitrided powders is very limited compared to papers on the synthesis of nanopowders. In particular, the route of sintering aids addition, in situ or by mixing powders may influence the development of the microstructure. The thermal stability studies of different grades of nanopowders have shown the influence of the precursors used and the chemical composition. In order to fabricate dense nanocomposites, it is necessary to study sintering kinetics of various nanopowders and to establish relationships between the characteristics of the powders and the sintering behaviour. Influence of the synthesis process NANOPOWDERSFROM LIQUID PRECURSORS The mixtures are densified with a linear increase of temperature or during isothermal treatments. The shrinkage curves on Figure 7.14 show that powders with a high N content (C/N = 0.2) exhibit a better ability for sintering than carbon-rich powders (C/N = 2.5). The presence of additives containing aluminium (HSAl09, HMDS35 + 3A) enhances the densification rate due to a possible oxynitride liquid phase. In the case of a high N content combined with the addition of Al introduced during the synthesis (HSAl09) a positive effect is observed compared to the grade HMDS35+3A where Al is added by a mechanical mixing. The rate of densification is higher for the HSAl09 powder than for HMDS35 + 3A at the same temperature. However, the shrinkage is blocked at around 70% as it is shown by the

Figure 7.14 Isobar densification curves. Relative density ● HMDS35 + 3A, ▼ HMDS66, ■ HSAl07, × HSAl09.

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Figure 7.15 Isobar densification curves. Relative density. ● HMDS35+6Y3A, ○ HSAl09+6Y, × HSAl09. Table 7.5 Chemical composition of the mixtures HMDS35 + 6%Y2O3 + 3%Al2O3 and HSAl09 + 6%Y2O3 Grade Chemical composition (wt%) Si C N O Al Y 91% HMDS35 46.0 5.4 35.0 4.6 — — 6% Y2O3 — — — 1.3 — 4.7 3% Al2O3 — — — 1.4 1.6 — HMDS35 + 6Y3A 46 5.4 35.0 7.3 1.6 4.7 94% HSAl09 46.6 6.4 33.8 5.7 1.5 — 6% Y2O3 — — — 1.3 — 4.7 HSAl09 + 6Y 46.6 6.4 33.8 7.0 1.5 4.7 isothermal experiment carried out at 1650°C on the HSAl09 powder (Figure 7.15). A sharp increase of the shrinkage is observed at the beginning of the isotherm, then the rate of densification decreases progressively. This behaviour could be related to the decomposition of the nanopowder (see thermal stability) which hinders the sintering. In order to restrict this decomposition, 6 wt% of Y2O3 is incorporated in the starting mixture (HMDS35 + 6Y3A, HSAl09 + 6Y) to have the same additive content (Table 7.5). Isobar curves reported in Figure 7.16(a) show clearly that an increase of the additive content favour the wetting of the surface by the liquid phase which stabilises the nanoparticles. This phenomenon explains that the beginning of the shrinkage starts at a lower temperature and the relative density is higher for all the grades. One notices that whatever the additive contents (HSAl, HSAl + 6Y), the rate of densification (Figure 7.16(b)) depends on the percentage of

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Figure 7.16 Isobar densification curves. (a) Relative density. (b) rate of densification. ▲ HSAl07+ 6Y, ◇ HSAl09 + 6Y, ☐ HSAl15 + 6Y, ● HMDS35 + 6Y3A. C in the starting nanopowders. When the C/N ratio increases, the rate of densification and the final density are lowered. Shrinkage curves with a sigmoidal shape are obtained only in the case of N rich powders. The comparison of the different powders confirms the beneficial approach of prealloyed powders when Al is introduced in situ during the synthesis (Figure 7.16). The phase contents (α-Si3N4, β-Si3N4, Si2N2O) of the nanocomposites are reported in Table 7.6. The incorporation of additives promotes the α- to β-Si3N4 transformation but also file:///C|/Users/steppenwolf/Documents/educatio...ks/to%20burn/nanocomposites/files/page_256.html (1 of 2) [20/01/2010 10:05:42 AM]

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the precipitation of silicon oxynitride.

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< previous page Page 257 Table 7.6 Crystallised phases Grade HSAl09 HSAl09 + 6Y HSAl09 + 6Y3A

Si2N2O + ++ +++

Si3N4 +++ ++ +

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Figure 7.17 Isobar densification curves. Relative density. ▼ SiCN35 + 6Y3A, ∇ SiCN35 + 6Y3A, ■ SiCN29 + 9Y4, 5A, ☐ SiCN29 + 6Y3A. NANOPOWDERS FROM GASEOUS PRECURSORS Different grades of SiCN powders mixed with the same additives are hot-pressed with a linear increase of temperature. The kinetic curves are reproduced in Figure 7.17. A better compaction at room temperature of these powders is observed, compared with the previous grades; moreover, the relative density changes with the gas used during the synthesis of the powder. This difference could be related to the nature and the quantity of the adsorbed species on the surface of the particles as it has been demonstrated by TGA–MS(see thermal stability). These species influence the behaviour of the nanoparticles during the mixing in ethanol. Carbon is detrimental to the sintering processes whatever the additives percentage. When C/N ratio increases, densification is delayed. The same relative density (80%) is reached at 1650°C for SiCN29 + 6Y3A and at 1740°C for SiCN35 + 6Y3A. XRD patterns evidence the presence of β-Si3N4 and an amount of β-SiC which increases with the C/N atomic ratio in the starting powders. Comparative study of micro–nano and nano–nanocomposites In order to compare the sintering behaviour and to determine the experimental conditions to obtain fully dense composites, different micro or nanopowders were used. The mixtures

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Page 258 (SCN29+6Y3A, HMDS35+6Y3A, HSAl09+6Y, Si3N4/SiCN35+6Y3A) with the same additive content were heat treated. The C/N ratio equal to 0.2 represents a compromise to obtain a composite with a final composition of Si3N4/8 wt% SiC which exhibits a ductile behaviour at high temperature. SiC particles stabilise the microstructure during a transitory stage for hot forming. Then this ratio allows an acicular grain growth of βSi3N4 crystals during a post forming treatment to obtain a good toughness. The shrinkage curves evidence the difficulties to have a high compactness with nanopowders (Figure 7.18 (a)). Nevertheless, the rate of densification (Figure 7.18(b)) increases rapidly and fully dense composites are obtained at lower temperatures than for the micro–nano mixture. Isothermal curves confirm the ability for sintering of these fine powders. On the contrary,

Figure 7.18 Isobar densification curves. (a) Relative density, (b) rate of densification. ● HMDS35 + 6Y3A, ◇ HSAl09 + 6Y, ■ SiCN29 + 6Y3A, ☐ Si3N4/Si/C/N + 6Y3A. file:///C|/Users/steppenwolf/Documents/educatio...ks/to%20burn/nanocomposites/files/page_258.html (1 of 2) [20/01/2010 10:05:43 AM]

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Page 259 micro–nanocomposites reach 90% of the theoretical density after 2 h at 1550°C under 35 MPa (Figure 7.19). The α → β-Si3N4 transformation ratio was determined by XRD for the starting powders and for the sintered composites. Si3N4 content in the SiCN powders was calculated using the rule of mixtures and by taking into account the chemical analysis results (Table 7.7). This ratio is equal to 14% for micro–nanocomposites and 100% for the nano–nanocomposites (Isotherm 1550°C, 2 h, N2). For the two mixtures, the beginning of the shrinkage is detected at 1460°C which illustrates the beneficial effect of the size of the powders. For the nano– nanocomposites with a densification higher than 96% after 90 mn 1550°C, the shape of the curves changes with the starting powder. The SiCN29 + 6Y3A grade, nanopowders issued from gaseous precursors, exhibits a higher rate of densification than the powders issued from liquid precursors (HMDS35 and HSAl09).

Figure 7.19 Isobar densification curves (1650°C). Relative density ● HMDS35 + 6Y3A, ◇ HSAl09 + 6Y, ■ SiCN29 + 6Y3A, ☐ Si3N4/Si/C/N + 6Y3A. Table 7.7 α → β-Si3N4 transformation rate α/α + β (wt%) Transformation rate Si3N4–α → β(%) Powders Si3N4 UBE 100 SiCN29 (87% Si3N4) 47 SiCN35*(45% Si3N4) 78 SiCN29 + 6Y3A nano/nano Mixture 37 100 Sintered material 0 Si3N4/SiCN35 + 6Y3A micro/nano Mixture 98 14 Sintered material 84

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Figure 7.20 SiCN29 nanocomposite sintered at 1600°C. (a) SEM, (b) TEM. Microstructure Si3N4/SiC NANOCOMPOSITES Observations of SiCN nano–nanocomposites show a uniform microstructure with equiaxed grains. The mean diameter of the grains is about 150 nm and the maximum length of β-Si3N4 crystals is lower than 450 nm (Figure 7.20). TEM micrographs confirm these observations. Small particles (Ø < 3 nm) with stacking faults correspond to SiC nanoprecipitates which are randomly dispersed within β-Si3N4 grains or at grain boundaries. The grain size for micro–nanocomposites (Figure 7.21) is approximately twice greater than for nano–nanocomposites. The microstructure refinement and the distribution of SiC precipitates is explained by the size of the starting powders. The dissolution is faster in the case of nanopowders and the supersaturation of the liquid phase is reached rapidly, which induces a more homogeneous nucleation process. Moreover SiC nanoprecipitates formed can act as nuclei. In consequence, the number of precipitate seeds increases and crystal growth is limited. Si3N4/Si2N2O NANOCOMPOSITES In the case of composites issued from HMDS or HSAl powders, the grain size is wider. Equiaxed grains with 200 nm in diameter are observed but also acicular crystals which are distributed all around agglomerates of small grains. The development of this bimodal distribution of grains could be explained from results obtained by Emoto et al. (1998), on the sintering of (Si3N4-Mg2Al4Si5O18) powders. The precipitation of Si2N2O nuclei in the first stage is due to the high oxygen content in the liquid phase. The growth of these crystals is favoured by a rapid migration of the species linked to the low viscosity of the eutectic. This process is followed by compositional changes around Si2N2O crystals which induce in a second stage the precipitation of β-Si3N4. As a result, fine Si3N4 crystals with equiaxed morphology are confined by acicular Si2N2O grains.

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Figure 7.21 Si3N4/SiCN29 micro–nanocomposite sintered at 1600°C. Discussion According to previous works (Bellosi 1995) on the sintering of nanopowders, the compaction at room temperature is lower than 40%. This value which characterises fine powders is explained by hard agglomerates with particles bonded by Van der Waals forces but also by the nature of chemical species on the surface of the particles. When a bimodal distribution is used as in the case of the micro–nano system, the compactness is favoured. This difference explains that at the beginning of the shrinkage the rate is higher for nano–nanocomposites. The rearrangement process and the solubility of the nano particles which are inversely proportional to the size are faster. As a result of these phenomena, the α → β-Si3N4 transformation is completed only for the nano–nanocomposites hot-pressed for 2 h at 1550°C. These results emphasise the influence of the process to introduce additives for the sintering of silicon nitride materials. Kinetic curves for prealloyed powders with the same additive content are faster which shows that the distribution of aluminium at an atomic scale facilitates the rearrangement process by a simultaneous wetting of the particles. On the contrary, with the classical route by mixing powders the oxynitride eutectic is localised in a limited number of grain boundaries which necessitates its migration and the impregnation of the porosity. It has been demonstrated that the sintering of silicon nitride is closely linked to the volume of liquid phase. A minimum quantity is necessary to impede its decomposition. The same behaviour is observed with nanopowders. The too low content of aluminium introduced in situ during the synthesis does not allow it to reach the theoretical density. Two processes are competing when temperature increases. The densification starts due to the rearrangement and dissolution–diffusion–reprecipitation mechanisms but also the degradation of SiCNO grains with a release of volatile species. Above 1650°C, the decomposition exceeds the rate of the shrinkage. For the behaviour at high temperature of these powders, an important parameter is the C content. The two different grades obtained by pyrolysis of liquid or gaseous precursors contain a high C content. The presence of C delays the shrinkage but the role of C differs according to its nature, free C or bounded to Si. In the case of HMDS and HSAl powders, it is mainly

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Page 262 free C which impedes the densification. As observed by Doug et al. (1997) and explained in the study of thermal stability, at high temperature an in-situ carboreduction reaction takes place in the powder and carbon monoxide is formed. The release of CO modifies the oxygen content of the liquid phase and induces an increase in its viscosity. In consequence, these changes are followed by a higher temperature for the beginning of the shrinkage and by a lower densification rate. Concerning C bounded to Si its influence on the sintering behaviour of powders issued from gaseous precursors is less important than free C. SiC nanoprecipitates are progressively formed but they are practically inert towards the liquid phase in this range of temperature. The densification of silicon carbide with Al2O3 and Y2O3 additives is only efficient above 1750°C. This is the reason why SiC nanoparticles formed by the dissolution of SiCNO phase impede the intergranular diffusion of the species and the oversell rate of densification. These studies show also the great reactivity of nanopowders which changes with the nature of the precursors. When powders issued from liquid precursors are hot-pressed, Si2N2Ois formed during sintering. This oxynitride precipitates in the YSiAlON system only when the eutectic is saturated by oxygen. This element comes from additives but also from reactions with the environment during the mixing process of the powder in a liquid media. Recently, Wada et al. (1999) have demonstrated that the oxygen content of carbonitrided nanopowders increases when a slurry is prepared with ethanol. In the case of the HMDS nanopowders, the amorphous state and the different surface groups could explain the specific reactions with ethanol. The experiments on the thermal stability have evidenced the presence of NH and NH2 groups bounded on the surface of the particles. It is known that these groups are easily hydrolysed or can react with ethanol which induces the formation of Si—O—Si bonds. The quantity of oxygen fixed at the surface varies with the surface groups which depend on the pyrolysis conditions and the nature of the precursors. 7.5 Conclusion These kinetics studies in the sintering of Si3N4/SiC composites from various powders evidence the role of the SiC content on the starting and the development of the microstructure. When Si/C/N nanopowders are added to a Si3N4 micropowder, they generate an increase in the densification rate due to the faster solubility of nanograins in the liquid phase. The densification kinetics are also closely linked to the grain structure and the chemical composition. To fabricate fully densified composites, it is necessary to use crystallised powders which are stable and more able to be sintered. The shrinkage curves show that better results are obtained with powders issued from gaseous precursors than from liquids. For the different grades of powders the rate of densification decreases with a higher C/N atomic ratio. The incorporation of SiC nanoparticles hinders the shrinkage due to their inertness towards the eutectic. Moreover, free C present at the surface of the grains or as particles in the powders reacts with oxygen from additives which delays the formation of the liquid phase. On the contrary, SiC nanoparticles quicken the α/βSi3N4 transformation and form nuclei for β-Si3N4 grains which induces a heterogeneous nucleation and a fine microstructure. The route to incorporate additives is also an important factor. The distribution of aluminium at an atomic scale favours the migration of the liquid, the wetting of the surface of the grains and the rearrangement process. These experiments validate the interest of prealloyed powders and their ability to fabricate Si3N4/ SiC nanocomposites.

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Page 263 References Bellosi, A. (1995) Mater. Sci. Forum, 195, 79. Bowen, L. J., Weston, R. J., Carruthers, T. G. and Brook, R. J. (1978) J. Mater Sci., 13, 350. Cauchetier, M., Croix, O., Luce, M., Baraton, M.-I., Merle, T. and Quintard, P. (1991) J. Eur. Ceram. Soc., 8, 215. Doug, S., Jiang, D., Tan, S. and Guo, J. (1997) J. Mater. Sci. Lett., 16, 1080. Emoto, H., Mitomo, M., Wang, C. M., Hirosturu, H. and Inaha, T. (1998) J. Eur. Ceram. Soc., 18, 527. Gheorghiu, A., Sénémaud, C., Roulet, H., Dufour, G., Moreno, T., Bodeur, S., Reynaud, C., Cauchetier, M. and Luce, M. (1992) J. Appl. Phys., 71 (9), 4118. Greenwood, G. W. (1956) Acta Metall., 4, 241. Greskowich, C., Prochazka, S. and Rosolowski, J. H. (1977) Sintering behaviour of covalently bonded materials, Nitrogen Ceramics (ed. F. L. Riley), Noordhoff, Leyden. Grün, R. (1979) Acta Crystallogr., B35, 800. Hampshire, S., Nestor, E., Flynn, R., Besson, J. L., Rouxel, T., Lemercier, H., Goursat, P., Sebaï, M., Thompson, D. P. and Liddell, K. (1994) J. Eur. Ceram. Soc., 14, 261. Hampshire, S., Park, H. K., Thompson, D. P. and Jack, K. H. (1978) Nature, 274, 880. Ishizaki, K. and Yanai, T. (1995) Silic. Indus., 7–8, 215. Kang, S. J. and Han, S. M. (1995) MRS Bull., 2, 33. Kleebe, H. J. (1992) J. Am. Ceram. Soc., 10, 151. Lang, J. (1977) Nitrogen Ceramics (ed. F. L. Riley), Noordhoff, Leyden. Lee, R. R., Chen, C. J. and Lin, J. T. (1994) Ceram. Trans., 42, 221. Lewis, M. H. (1993) Trans. Tech. Publ., 89–91, 334. Niihara, K. (1991) Ceram. Sci. Jpn, 99, 10, 974. Niihara, K., Suganuma, K., Nakatiera, A. and Izaki, K. (1990) J. Mater. Sci. Lett., 9, 598. Pan, X., Gu, H., Van Weeren, R., Dauforth, S. C., Cannon, R. and Rühle, M. (1996) J. Am. Ceram. Soc., 79, 2313. Rossignol, F., Goursat, P., Besson, J. L. and Lespade P. (1994) J. Eur. Ceram. Soc., 13, 299. Sasaki, G., Nakase, H., Sugaruma, K., Fujita, T. and Niihara, K. (1992) J. Ceram. Soc. Jpn, 100, 536. Siegel, R. W. (1994) Springer Ser. Mater. Sci. Phys. New Mater., 27, 65. Ténégal, F., Flank, A.-M. and Herlin, N. (1996) Phys. Rev. B, 54, 17, 12029. Wada, S., Kondo, Y., Sudo, E. and Maki, Y., (1999) J. Ceram. Soc. Jpn, 107, 611. Wakai, F., Kodama, Y., Sakaguchi, S., Murayama, N., Izaki, K. and Niihara, K. (1990) Nature, 344, 421. Wang, C. M., Emoto, H. and Mitomo, M. (1998) J. Am. Ceram. Soc., 81, 5, 1125. Ziegler, G., Heinrich, J. and Wötting, G. (1987) J. Mater. Sci., 22, 3041.

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Page 264 8 Mechanical properties Jean-Louis Besson The composites used to assess the mechanical properties are mainly those fabricated using Si/C/N nanoparticles synthesised from gaseous precursors as the source of SiC (grade SiCN35). They are referred to as SiC micro–nanocomposite when the Si/C/N powder is added to Si3N4 powder and Si/C/N nanocomposite when the Si/C/N powder is the only source of both Si3N4 and SiC. In the assessment of the compressive creep behaviour, composites elaborated with nanoparticles coming from the decomposition of hexamethyldisilazane (grade HMDS35) or that of a liquid mixture of hexamethyldisilazane and aluminium isopropoxyde (grade HSAl09) are also briefly considered under the respective references: HMDS and HSAl nanocomposites. 8.1 Room temperature properties 8.1.1 Introduction Niihara and coworkers (Niihara 1991) have observed a remarkable improvement in strength and toughness by incorporating nanosized particles of SiC in various ceramics matrix (Al2O3, MgO, Si3N4). In every case, most of the finest SiC particles (typically less than 200 nm) were predominantly dispersed within the matrix grains and some larger particles at the grain boundaries. From HRTEM investigations, it was concluded that no impurity phases were present at the interfaces between the SiC particles and the matrix grains so that the particles were directly bonded to the matrix. Moreover, the fracture mode was changed from a mixed intergranular–intragranular mode to a completely transgranular mode. Though different toughening and grain-refinement mechanisms were proposed for the different systems, the basic idea was that the intragranular nanosized particles play a major role in the increase in toughness and strength up to high temperatures, and this for two main reasons: first, the development, generated by the large thermal expansion mismatch, of high tensile hoop stresses around the particles within the matrix grains, leading to crack deflection; second, the direct bonding of the SiC particles with the matrix even in the presence of grain boundary phases coming from sintering aids or impurities. In the case of a Si3N4 matrix (Niihara 1991), the fracture strength increases from 850 MPa to a peak value of 1550 MPa and the fracture toughness from 4.5 to 7.5MPam1/2. The peak values were both obtained for 25 wt % SiC. Similar trends were observed by Niihara’s group (Sawaguchi et al. 1991; Sasaki et al. 1992; Hirano and Niihara 1995) for Si3N4/SiC nanocomposites processed by various routes, though the peak position differed and the improvements were less pronounced. The degradation of the properties at high SiC content was claimed not

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Page 265 to be caused by the dispersed nanoparticles but related to flaws coming from aggregation of SiC particles (Sasaki et al. 1992). In Al2O3/SiC nanocomposites, neither Zhao et al. (1993) nor Borsa et al. (1997) or Xu et al. (1997) found an influence of the intragranular SiC nanoparticles on the intrinsic material toughness. In the case of Si3N4 matrix composites, Pezzoti and Sakai (1994), using a material densified without any sintering agent, also excluded any strengthening effect due to the presence of intergranular SiC particles. Their results are consistent with earlier work by Greskowich and Palm (1980) who found that in Si3N4/SiC composites where the sintering aids (BeSiN2 and SiO2) dissolve in the β-Si3N4 lattice, the toughness was independent of the SiC volume fraction up to 23 vol%. Contrary to Niihara and co-workers, most of the other researchers (Rendtel et al. 1995; Sajgalik et al. 1996; Sato et al. 1995; Liu et al. 1997; Hermann et al. 1998) have reported at best no improvement of either toughness or strength and more generally a degradation. As it seems to be now established that the intragranular location of SiC nanoparticles does not enhance the ‘intrinsic’ fracture properties of these socalled ‘nanocomposites’, changes in fracture toughness and fracture strength are to be related to the efficiency with which the ‘classical’ energy dissipating mechanisms (such as crack deflection, crack bridging or closing forces) are acting in the material under study, taking into account its actual microstructure. 8.1.2 Fracture strength Fracture strength, σr, was evaluated by 3-point bending test using rectangular bars (4 × 3 × 25 mm3). The bars were bevelled parallel to their length in order to eliminate edge flaws. The tensile surfaces of the specimens were perpendicular to the hot-pressing direction and polished to a 3-μm diamond finish. The span length was 20 mm and the cross-head speed 0.2 mm/min.

(8.1.1) where F is the ultimate load, w the width, h the height and x the distance between the section of rupture and the nearest fulcrum point. Changes in fracture strength with the sintering temperature are shown in Figure 8.1. For a sintering temperature of 1550°C, the strength increases with the SiC content in spite of the lower densification. The strength of the monolith increases sharply after sintering at 1600°C and remains stable for higher temperatures. The same trend is observed for the 5 wt% SiC micro–nanocomposite whereas in the case of the 10 wt% SiC micro–nanocomposite, the strength increases slowly up to a sintering temperature of 1650°C. As the densification of the monolith is complete after hot-pressing at 1550°C, the change in strength between 1550°C and 1600°C cannot be related to a change in porosity, and, as it remains constant for higher sintering temperatures, it cannot be associated with the acicular growth of the β-Si3N4 grains. The remaining explanation is based on the existence of residual stresses. X-ray diffraction (XRD) studies have shown that after sintering at 1550°C, the α-and β-Si3N4 grains are under compression, which implies that the intergranular glass phase experiences local tensile stresses (Rossignol 1995). These internal stresses decrease as the α-toβ-Si3N4 transformation goes on and stabilises when the α content becomes lower than about 50%. In the case of the 5 wt% SiC micro–nanocomposite, the α content after sintering at 1550°C is ~55% instead of 70% for the monolith and the strength is higher, in spite of a lower

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Figure 8.1 Fracture strength v. sintering temperature: ▲ monolith; ◆ 5% SiC micro–nanocomposite; ● 10% SiC micro–nanocomposite; ■ Si/C/N nanocomposite as-sintered; annealed 5 h at 1600°C. density, due to the lower residual tensile stress in the glass phase. The increase in strength when the sintering temperature is raised to 1600°C can be attributed partly to an increase in densification (from 97.5 to 99% TD) and partly to the progress in the α-toβ-Si3N4 transformation. As the α content in the 10 wt% SiC micro–nanocomposite is already lower than 50% after sintering at 1550°C, the increase in strength is solely due to the increase in densification from 83 to 99%. In the case of the Si/C/N nanocomposite, the strength is rather low (470 ± 30 MPa) for the as-sintered material (1600°C, 1 h), but it increases to 570 ± 30 MPa after a 5 h heat treatment at the sintering temperature and becomes similar to that of the 10 wt% SiC micro– nanocomposite. 8.1.3 Toughness Fracture toughness, K1C was estimated by microhardness indentation method in a Vickers test using Evans and Charles’ formula (Evans and Charles 1976). (8.1.2) where P is the indentation load (N) and c the half median crack length (m) and K1C is expressed in MPa m1/2. The toughness of the monolith increases twofold in the sintering temperature range whereas the increase in toughness of the micro–nanocomposites remains moderate

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Figure 8.2 Toughness v. sintering temperature: ▲ monolith; ◆ 5% SiC micro–nanocomposite; ● 10% SiC micro–nanocomposite; ■ Si/C/N nanocomposite as-sintered. (Figure 8.2). The toughness of the as-sintered Si/C/N nanocomposite is similar to that of the micro– nanocomposite though a higher fraction of the SiC precipitates is found inside the Si3N4 grains. When similar microstructures are considered (full density and equiaxed Si3N4 grains, i.e. sintering temperatures of 1550°C, 1600°C and 1650°C for the monolith, the 5 wt% SiC and the 10 wt% SiC micro–nanocomposites, respectively) the toughness is similar (Figure 8.3). However, if the comparison is made for the minimum sintering temperature leading to full density whatever the material (1600°C), the toughness of the monolith is much higher than those of the micro–nanocomposites that enlightens the dominating role of the aspect ratio of the Si3N4 grains. The toughness of the monolith is enhanced by the development of elongated βSi3N4 (the aspect ratio rises from 3.5 to 6) as usually observed for self-reinforced silicon nitride. In contrast, the presence of a large amount of intergranular SiC nanoprecipitates, though accelerating the α-to β-Si3N4 transformation, inhibits the acicular growth of the β-Si3N4 grains that remain quite equiaxed in shape and this saturation in aspect ratio prevents an increase in fracture toughness. In the case of the Si/C/N nanocomposite, post-sintering heat treatments at temperatures superior to 1650°C, lead to a more and more acicular shape for the β-Si3N4 grains and the toughness increases from 3.6 ± 0.3 to 5.4 ± 0.4 MPa m1/2 (Table 8.1). An additional increase in toughness is obtained after a further heat treatment at 1400°C leading to a partial crystallisation of the amorphous intergranular phase. So, it is confirmed that the main factor is the aspect ratio of the β-Si3N4 grains and that the intragranular SiC particles play a minor role, if any.

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Figure 8.3 Fracture toughness at room temperature:—◆—after sintering at 1600°C; - - -■- - - for similar microstructures. Table 8.1 Room temperature toughness of the SiC nanocomposite after various post-sintering heat treatments Heat treatment Toughness (MPa m1/2) Temperature (°C) Time (h) As-sintered 3.6 ± 0.3 1600 5 3.5 ± 0.1 1650 5 3.8 ± 0.5 1700 5 4.6 ± 0.4 1750 5 5.4 ± 0.4 1750 + 1400 5+3 6.9 ± 0.7 8.2 High-temperature properties 8.2.1 Introduction High-temperature properties of silicon nitride based ceramics, though influenced by the size and morphology of the Si3N4 grains, are primarily controlled by the nature and chemistry of the intergranular phase. After sintering, the intergranular phase is amorphous but can be converted more or less completely into crystalline phases by a post-sintering heat treatment. However, there always remain thin intergranular films. The most profound influence on the properties comes from the viscosity of these grain boundary films. For instance, MgO additive

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Page 269 leads to a low viscosity intergranular phase whereas Y2O3 converts easily to N-melilite (Y2Si3N4O3). The viscosity of the residual glass phase must be high enough to prevent cavitation and grain separation, but low enough to accommodate grain boundary sliding through diffusion of matter. The rheology of the intergranular phase can also be modified by a small amount of impurity such as fluorine, which has a detrimental effect on the high-temperature strength (Tanaka et al. 1994; Rendtel et al. 1995). Then, in order to assess the actual role of SiC nanoparticles on the high-temperature behaviour, it is necessary to compare the behaviour of the composites with that of their monolithic counterpart densified with the same additives. Moreover, the choice of the densification aids will be different depending on the high-temperature properties which are desired, high strength and high creep resistance or high ductility and even superplastic behaviour: in the former case, yttria alone will be used whereas it will be a mixture of yttria and alumina in the latter case. So, Niihara and co-workers (1990) have obtained an improvement in the flexural strength by incorporating 32 wt% SiC nanoparticles in a silicon nitride densified with 8 wt% yttria. The 4-point bending strength was still 840 MPa at 1500°C whereas that of the corresponding monolith had decreased to 500 MPa at 1400°C. This improvement was attributed to the fact that the SiC nanosized particles in the grain boundaries were directly bonded to the Si3N4 adjacent grains, without any impurity phase. The same grade of composite was also tested in tensile creep and it was noted that, contrary to the monolith that suffered severe cavitation induced by grain boundary sliding, cavity formation was rarely observed in the composite (Hirano et al. 1996). From the study of the bending creep of silicon nitride doped by nanosized SiC particles from different origins and sintered with Y2O3 aid only, Rendtel et al. (1995, 1998) also concluded that the addition of SiC nanoparticles is a very effective means of improving creep resistance of silicon nitride ceramics. On the contrary, if it is a high ductility that is desired, Si3N4/SiC nanocomposites densified with Y2O3 or MgO are not suitable as they fracture in a brittle manner with less than 2% elongation whereas the grades sintered with Y2O3 plus Al2O3 exhibit a high ductility which can reach more than 100% elongation (Rouxel et al. 1992). 8.2.2 Tensile ductility For the present materials, the high-temperature tensile behaviour was studied in the 1595–1640°C temperature range under nitrogen atmosphere. Dog-bone shape specimens with 10 mm long and 3 mm diameter cylindrical gauges were used. The gauges were polished finally with 0.3 μm diamond paste. Since large elongations were involved, true strain, εt and true stress, σt, were calculated on the basis of a constant volume throughout the deformation: (8.2.1) (8.2.2) where ε is the nominal strain (or elongation), l/l0, with l0 the initial gauge length and σ the nominal stress. The experiments were conducted at constant strain rate by monitoring the cross-head speed, , from the measurement of the elongation, ε:

(8.2.3)

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Page 270 Table 8.2 Theoretical SiC content, sintering temperature (Ts), mean grain size, aspect ratio of β-Si3N4 grains and α/β-Si3N4 ratio Material SiC (wt%) Ts (°C) Mean grain size (nm) Aspect ratio β-Si3N4 Si3N4 SiC α/β-Si3N4 Monolith 1550 250–300 70/30 1600 350a 30/70 3–5 1700 400a 0/100 5–7 Micro–nanocomposite 10 1600 400 30–40 30/70 Low 20 1650 100–400a 30–40 0/100 3–4 Si/C/N nanocomposite 8 1600 150–450


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Page 271 Figure 8.4 shows the tensile curve and the change in hardening rate of the 70/30 monolith tested at 1595°C under a strain rate of 5 × 10−5 s−1. Four regions can be identified on the tensile curve. The first region corresponds to the elastic response and a transient stage. The second and the third regions are associated with regimes of permanent and uniform deformation and the fourth regime to a stage of cavitation damage. The hardening rate in region 2 was smaller than in region 3. It was concluded that deformation in region 2 was essentially controlled by the αphase whereas it was controlled by the acicular β-grains in region 3. The transition between the two regions occurs for a volume fraction of α-phase of about 30%. The low hardening rate in region 2 was attributed to the equiaxed shape of the α-Si3N4 grains and to the transient liquid phase which accompanies the α- to βphase transformation. In region 3, the strong hardening, with a quasi-linear hardening rate, was related to the increase of the cross-linking degree in the acicular β-grain network and, to a lesser extent, to grain growth, that makes more difficult grain boundary sliding which is the major mechanism of superplasticity (Langdon 1994). The tensile curves for the composites and their corresponding monoliths are shown in Figure 8.5. The constant strain rate was 5 × 10−5 s−1 and the temperature 1615°C. The 10 wt% SiC composite was also tested at 1595°C and 1640°C. The experimental data are collected in Table 8.3. The deformed specimens had uniform elongation all along their gauge sections and the rupture always happened within the gauge. The surfaces were smooth and free of cracks but were covered with a thin whitish layer, which could be due to volatilisation of SiO gas from the YSiAlON liquid phase and/or to a slight active oxidation by oxygen atoms present in small quantity in the nitrogen gas. The resulting weight loss increased with the test duration but was at worst 0.6%.

Figure 8.5 Tensile curves under a strain rate of 5 × 10−5 s−1: M 30/70 and M 0/100 monoliths and 20% SiC micro–nanocomposite at 1615°C; 10% SiC micro–nanocomposite 1595°C, 1615°C, 1640°C.

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Page 272 Table 8.3 Experimental data from the tensile curves Material Temperature (°C) t(s−1) εrupture (%) εt σmax (MPa) dσt/dεt (MPa) 70/30 monolith 1615 5 × 10−5 66.5 0.510 35 55/83 30/70 monolith 1615 5 × 10−5 61.2 0.480 55 110/123 0/100 monolith 1615 5 × 10−5 24.6 0.220 37 130 10 wt% SiC composite 1595 5 × 10−5 5.4 0.053 35 10 wt% SiC composite 1615 5 × 10−5 10.0 0.095 38 133 10 wt% SiC composite 1640 5 × 10−5 17.4 0.160 29 80 20 wt% SiC composite 1615 5 × 10−5 3.3 0.033 30 10 wt% SiC composite 1595 10−5 37.7 0.320 23 46 The only grade which shows up the four regions identified in the reference is the 30/70 monolith. In this case, the deformation at rupture is high (ε ~ 60%) and a slight decrease in stress occurs before fracture. However, region 2 is barely marked due to an initial α-Si3N4 volume content close to that for which the transition between regions 2 and 3 occurs. The fracture surface was rough on the whole section suggesting that rupture is brought about by cavitation damage. Neither region 2 nor region 4 are observed for the 0/100 monolith which had a starting microstructure consisting entirely of acicular β-Si3N4 grains and for which rupture was governed by slow crack growth. The presence of an increasing fraction of β-Si3N4 grains with higher and higher aspect ratios explains the increase in hardening rate in region 3 from 80 to 130 MPa and the decrease in elongation at rupture from 66.5 to 24.6%. The tensile behaviour of the 10 wt% SiC composite contrasts strongly with that of its monolithic counterpart. The deformation at rupture is very low and the fracture surface shows a well-demarcated rough region, associated to slow crack growth, and a smooth region of fast and essentially brittle fracture regime. This underscores the primary role of the SiC nanoparticles and of their distribution. Three arguments can be invoked to explain the role of the SiC particles (Mayne et al. 1988). First, since a large amount of glass phase is being trapped in the SiC grain pockets, the intergranular films are thinner than in the monolith. Second, sliding of the β-Si3N4 grains is hampered by the SiC particles located along the grain boundaries. Finally, the largest Si3N4 grains support the stress and the SiC particles gathered together in the pockets are essentially unstressed and do not participate in the deformation. This is accentuated by an increase in the SiC content. On the contrary, an increase in the test temperature, lowering the viscosity of the intergranular glass phase and possibly increasing its amount, facilitates grain boundary sliding, reduces the hardening rate and increases the deformation at rupture. In all cases, failure occurs at the maximum stress after a deformation of less than 25%. In Figure 8.6, the tensile curves at 1595°C of the 10 wt% SiC composites tested under different strain rates (5 × 10−5 and 10−5 s−1) are compared. The behaviour of the specimen tested at the lower strain rate is strongly different from that of the specimen tested at the higher strain rate. The elongation at rupture is seven times greater, the maximum stress one-third lower and the rupture happens after a significant lowering of the stress. A pseudo-stationary hardening stage is well developed, with a low hardening rate (dσ/ dε ~ 46 MPa). The fracture surface is rough on the whole section and cavitation damage shows up on polished sections. The volume fraction of cavities, derived from density and Young’s modulus measurements, was about 1.5%. This change in behaviour with strain rate contrasts with the behaviour

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Figure 8.6 Tensile curves of the 10% SiC micro–nanocomposite at 1595°C: 5 × 10−5 s−1; 10−5 s−1. (Reprinted from Mayne, M., Rouxel, T., Bahloul-Hourlier, D. and Besson, J.-L. (1998) J. Eur. Ceram. Soc., 18, Si3N4/SiCN nanocomposites: tensile ductility and rupture behaviour, 1985–1994 with permission from Elsevier Science.) of the monolithic ceramic where the hardening was strain-enhanced and essentially time-independent (Rouxel et al. 1997). Again, an explanation involving SiC nanoprecipitates has to be put forward. With regard to the temperature, it is speculated that the α- to β-phase transformation goes on during the tests and that grain growth occurs with more and more acicular grains as the transformation proceeds. Moreover, the acicular grains must tend to align towards the tensile axis. Indeed, this trend had been observed in the case of the monoliths though grain growth remained limited. The presence of a deformation texture was also corroborated by X-ray diffraction results which suggested the alignment of the c-axis of the elongated β-Si3N4 towards the tensile axis. In the case of the 10 wt% SiC composite tested under a strain rate of 5 × 10−5 s−1, the change in microstructure remained very limited due to the short duration of the test. Differences between the deformed and undeformed microstructures become substantial in the case of the composite tested under the strain rate of 10−5 s−1 for which the α- to β-Si3N4 phase transformation is complete and the majority of the grains elongated and aligned towards the tensile axis. However, the length of the acicular grains (1.2–1.9 μm) remains significantly lower than in the monolith tested during a similar duration (~2.5 μm), that confirms the role of the SiC nanoparticles in the limitation of grain growth. From these differences in the final microstructure ensuing from the lower strain rate, it is suggested that: • When the strain rate is high, the Si3N4 grains interlock and the diffusion is too slow to release the stresses that appear at the buttressing points, that results in a poor ductility and a rupture essentially brittle.

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Page 274 • On the contrary, under a sufficiently low strain rate, enough time is provided to release the local stresses at the interlocking points and the acicular β-grains align towards the tensile axis. The stress is then transferred to the SiC nanoprecipitates and sliding becomes active between the rounded SiC particles inside the pockets. The fact that in the latter case, the hardening rate is comparable to that in region 2 of the 70/30 monolith, where deformation is controlled by sliding of equiaxed grains, supports this hypothesis. In conclusion, under the higher strain rate, deformation is controlled by grain boundary sliding of the microsized Si3N4 grains whereas the SiC nanoprecipitates contribute to the deformation under the lower strain rate. 8.2.3 Creep behaviour The creep behaviour was investigated during compressive tests performed in air between 1200°C and 1450° C with nominal stresses ranging from 40 to 180 MPa. The specimens were parallelepipedic bars (3 × 3 × 8mm3). In the previous studies, as recalled above, different grades of Si3N4/SiC nanocomposites were developed according to whether an enhanced creep resistance or a high ductility were aimed at: grades for high creep resistance were densified with Y2O3 as the only sintering additive whereas a mixture of Y2O3 and Al2O3 was used for grades with high ductility. The present study focuses on grades sintered with yttria and alumina as additives. As in the case of tensile ductility, the creep behaviour was compared and discussed with reference to that of the monolith sintered at 1550°C. After testing in air at 1350°C, the monolithic specimens were covered with a 100 μm thick oxide scale, white in colour, and showed underneath a zone (400 μm in depth) depleted in yttrium. In contrast, the composite samples were black and the zone depleted in yttrium only about 100 μm thick. This underlines a better oxidation resistance of the composites. The compressive creep tests being conducted under constant load instead of constant stress, as soon as deformation becomes noticeable, true strain, true stress and true strain rate were calculated using equations (8.2.1), (8.2.2) and (8.2.3) which stand for the frictionless case (Debschütz et al. 1993). Though, in compression, εt, σt and t are negative, their absolute values were considered in drawing the curves. Typical creep curves obtained at 1350°C under a nominal stress of 180 MPa are shown in Figure 8.7 for the monolith and the composites. For the SiC micro–nanocomposites, the creep resistance is higher than that of the monolith and increases with the SiC content. On the contrary, the Si/C/N nanocomposite has a lower creep resistance. In the as-sintered state, it fails in 33 h after a brief tertiary stage whereas it sustains the 50 h test duration after annealing for 5 h at the sintering temperature (1600°C). The HMDS and HSAl materials have a very low creep resistance and fail in 2 and 5 h, respectively. Change of true strain rate v. time is displayed in Figure 8.8 for the monolith, the 10% SiC micro– nanocomposite and the Si/C/N nanocomposite tested at 1350°C under 90 MPa. For the micro– nanocomposite, the strain rate decreases rapidly during the first 15 h (primary creep) and then tends slowly towards a nearly constant value. After 35 h, the residual decrease is essentially due to the slow decrease in the true stress, (secondary creep). The primary stage is shorter for the Si/C/N nanocomposite and a steadystate creep is reached in about 5 h. In the case of the monolith, after the sharp initial decrease, the creep rate goes through a maximum around the tenth hour and, next, diminishes slowly (pseudo-stationary creep). In this latter

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Figure 8.7 Creep curves at 1350°C under 180 MPa: (1) monolith; (2) 5% SiC micro– nanocomposite; (3) 10% SiC micro–nanocomposite; (4) 20% SiC micro– nanocomposite; (5) Si/C/N nanocomposite, as-sintered; (6) Si/ C/N nanocomposite, heat treated 5 h at 1600°C; (7) HMDS; (8) HSAl.

Figure 8.8 True strain rate v. time at 1350°C and 180 MPa: ▲ Monolith; ● 10 wt% SiC micro– nanocomposite; ■ Si/C/N nanocomposite. Open symbols: data corrected for constant stress.

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Page 276 case, the decrease remains too pronounced to be solely attributed to the decrease in the true stress. No signs of cavitation were evidenced in the materials by transmission electron microscopy (TEM) investigation that is consistent with the absence of a tertiary regime on the creep curves. The most distinctive deformation-induced features were strain whorls located in grain boundaries with thin glassy film, and, in the case of the monolith, an increased number of acicular grains. In both cases, once the oxidised scale had been removed, no other crystalline phases were detected by XRD in addition to β-SiC and the Si3N4 polymorphs. The α/β-Si3N4 ratio in the composite samples remained unchanged during creep tests up to 1350°C, diminished slightly at 1400°C and the α-phase was no longer detected after creep at 1450°C. For the monolith, the α/β ratio decreased slowly up to 1300°C and then rapidly for higher temperatures. The data were analysed using the Dorn–Boltzmann relation: (8.2.4) where A is a constant, d the mean grain size, p the grain size exponent, σ the stress, n the stress exponent, Ea the apparent activation energy, R the gas constant and T the absolute temperature. Actual stresses, calculated according to equation (8.2.2), were used to determine the stress exponents. From the best fit straight lines in the ln /ln σ diagram, values of n = 0.78, 1.05 and 0.94 were obtained for the monolith, the 10 wt% SiC micro–nanocomposite and the Si/C/N nanocomposite, respectively (Figure 8.9). The apparent activation energies were derived in the ln /T−1 diagram (Figure 8.10) with strain rates normalised for a stress of 180 MPa using the stress exponents previously determined. Values of 514, 590 and 580 kJ/mol were found for the monolith, the 10 wt% SiC micro–nanocomposite and the Si/C/N nanocomposite, respectively. Activation energies

Figure 8.9 True strain rate v. stress at 1350°C: ▲ Monolith; ● 10 wt% SiC micro–nanocomposite; ■ Si/C/N nanocomposite; ☐ data corrected for constant stress.

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Figure 8.10 True strain rate v. reciprocal absolute temperature: ▲ Monolith; ● 10% SiC micro– nanocomposite; ■ Si/C/N nanocomposite. reported in the literature for silicon nitride densified with yttria and alumina range from 400 to 800 kJ/mol. For Si3N4/SiC nanocomposites, the only values are those reported by Rendtel et al. (1995). They range from 975 to 1085 kJ/mol with the exception of a value as high as 1435 kJ/mol. It is well established that the creep of Si3N4 ceramics containing an intergranular phase is strongly influenced by it (Lange et al. 1980a). The primary stage is dominated by the viscoelastic response of the glass itself (Lange et al. 1980b; Arons and Tien 1980) and diffusional creep in the stationary stage, characterised by a stress exponent ~1, is enhanced by the viscous grain boundary phase that provides paths of high diffusive conductance (Stocker and Ashby 1973). The redistribution and the exudation to the surface of a noticeable fraction of the glass phase (Clarke 1985; Besson et al. 1986) results in an improved packing of the grains. The contribution due to the viscoelastic response of the glass phase vanishes in about 20 h. Then, grain boundary sliding becomes restricted by grain-to-grain contacts at asperities along grain boundaries. These contacts induce localised elastic deformations that are evidenced by the ‘strain whorls’ observed in crept samples cooled under load (Lange et al. 1980b). Differential chemical potentials that arise due to stress gradients provide the driving force for a solution – migration – precipitation process to relax the localised stresses. The increase in the number of grain-to-grain contacts leads to a pseudo-stationary stage with a slowly decreasing creep rate until the hardening effect is balanced by the recovery effect due to diffusion. A detailed discussion of the creep behaviour of the materials tested in the present study has already been published (Besson et al. 1998; Doucey 1999). The conclusions are that this scenario applies to the composites. In the case of the monolith, the maximum in the creep rate,

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Page 278 which is not observed in the composites where the α to β transformation does not proceed during the test, is very likely to be associated to this transformation. The transformation is accompanied by an increase in the quantity of liquid phase and the precipitation of small-sized β-Si3N4 grains that favours an increase in the creep rate. Later, the strain hardening due to the acicular growth of the β-Si3N4 grains has the upper hand and the creep rate decreases. Eventually, grain-to-grain contacts form and the behaviour is then similar to the one observed in low α-Si3N4 containing silicon nitride ceramics. Since the stress exponents and the apparent activation energies of the monolith and of the composites are basically the same, the difference in creep resistance must derive from differences in the microstructure. Several features can be put forward to explain, the better resistance of the SiC micro– nanocomposite: • the intergranular phase in the composite is more refractory than in the monolith as evidenced by the change in Young’s moduli with temperature (Figure 8.11); • the SiC grains present in the intergranular films obstruct the easy paths of diffusion; • due to their larger size, the Si3N4 grains form a rigid network that bears the stresses so that the SiC grains gathered together in pockets are virtually stress-free and do not contribute to the deformation; • because of the trapping of a large amount of glass phase in these stress-free SiC grain pockets, the intergranular films in the composite are thinner than in the monolith. On the contrary, due to its very fine and quasi-equiaxed microstructure, the Si/C/N nanocomposite exhibits a high ductility, even higher than that of the monolithic counterpart.

Figure 8.11 Change in Young’s modulus v. temperature: ▲ monolith; ● 10 wt% SiC micro– nanocomposite. (Reprinted from Besson, J. L., Mayne M., Bahloul-Houlier D. and Goursat P. (1998) J. Euro. Ceram. Soc., 18, Si3N4/SiCN nanocomposites: influence of SiC nanoprecipitates on the creep behaviour, 1893–1904, with permission from Elsevier Science.)

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Page 279 As it has been demonstrated (Rossignol et al. 1995; Rouxel et al. 1997) that the monolithic silicon nitride can be tailored to exhibit a very high ductility, it is anticipated that this material, that presents a creep rate four times higher than that of the monolith, should behave superplastically at a lower temperature and a higher strain rate. 8.2.4 Influence of heat treatments A 5 h post-sintering heat treatment at 1600°C on the monolith develops an acicular β-Si3N4 structure with a high aspect ratio whereas the same treatment performed on the 10 wt% SiC micro–nanocomposite has little effect, that enlightens the role of the intergranular SiC nanoprecipitates in hindering grain growth. As a consequence, the creep resistance of the monolith was improved and became similar to that of the composite which was not affected (Figure 8.12). A heat treatment for a longer time and/or at a higher temperature is required for the SiC nanoprecipitates to be incorporated in the Si3N4 matrix (Sajgalik et al. 1996; Pan et al. 1996) and permit the development of an acicular structure enhancing the creep resistance. It is the case for the Si/C/N nanocomposite for which the creep resistance increases, as the toughness does, after heat treatments at temperatures superior to 1600°C. As for the toughness, the higher the temperature, the higher the creep resistance. As for the toughness too, an additional treatment at 1400°C further enhances the creep resistance which becomes similar to that of the heat-treated monolith or that of the SiC micro–nanocomposite.

Figure 8.12 Creep curves at 1350°C under 180 MPa: ▲ monolith; ● 10 wt% SiC micro– nanocomposite; ■ Si/ C/N nanocomposite; full symbols: as-sintered; open symbols: monolith and 10 wt% SiC micro– nanocomposite after heat treatment at 1600°C for 5 h. Si/C/N nanocomposite after heat treatment at 1600°C for 5 h and at 1400°C for 3 h.

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Page 280 8.3 Conclusion The role of SiC nanoparticles on the strength and the toughness of Si3N4/SiC composites is not yet clear. The increases obtained by Niihara’s group were not observed by other researchers. Concerning the ductility, the grades densified with yttria as the only densification aid have a better creep resistance than the ones sintered with both yttria and alumina whether they contain SiC nanoparticles or not. However, the amount, and above all, the distribution of the SiC particles strongly affect the high-temperature behaviour. The main effect is the pinning of the grain boundaries by the intergranular precipitates that hinders the acicular growth of the β-Si3N4 grains during sintering. This can result, for well-tailored conditions (in particular, a low sintering temperature), in a very fine, equiaxed microstructure with a high ductility suitable for hot-forming. Afterwards, post-forming heat treatments, at higher temperatures, may be carried out during which the SiC particles become included into the silicon nitride grains (Pan et al. 1996). Then, the grain growth of the β-Si3N4 grains is no longer inhibited and an acicular microstructure develops leading to an improved creep resistance. References Arons, R. M. and Tien, J. K. (1980) J. Mater. Sci., 15, 2046. Besson, J. L., Streicher, E., Chartier, T. and Goursat, P. (1986) J. Mater. Sci. Lett., 5, 803. Besson, J. L., Mayne, M., Bahloul-Hourlier, D. and Goursat, P. (1998) J. Eur. Ceram. Soc. Special issue Ceram. Matrix Compos., 18, 1893. Borsa, C. E., Jones, N. M. R., Brook, R. J. and Todd, R. I. (1997) J. Eur. Ceram. Soc., 17, 865. Clarke, D. R. (1985) J. Mater. Sci., 20, 1321. Debschütz, K. D., Caspers, B., Scheider, G. A. and Petzow, G. (1993) J. Am. Ceram. Soc., 76, 2468. Doucey, B. (1999) PhD thesis, Université de Limoges, France; http://thesenet.abes.fr. Evans, A. G. and Charles, E. A. (1976) J. Am. Ceram. Soc., 59, 371. Greskovich, C. and Palm, J. A. (1980) J. Am. Ceram. Soc., 63, 597. Hermann, M., Schuber, C., Rendtel, A. and Hübner, H. (1998) J. Am. Ceram. Soc., 81, 1095. Hirano, T. and Niihara, K. (1995) Mater. Lett., 22, 249. Hirano, T., Niihara, K., Ohji, T. and Wakai, F. (1996) J. Mater. Sci. Lett., 15, 505. Langdon, T. G. (1994) Mater. Sci. Eng., A174, 225. Lange, F. F., Davis, B. I. and Clarke, D. R. (1980a) J. Mater. Sci., 15, 601. Lange, F. F., Clarke, D. R. and Davis, B. I. (1980b) J. Mater. Sci., 15, 611. Liu, Q., Gao, L., Yan, D. S., Mandal, H. and Thompson, D. P. (1997) J. Eur. Ceram. Soc., 17, 593. Mayne, M., Rouxel, T., Bahloul-Hourlier, D. and Besson, J. L. (1998) J. Eur. Ceram. Soc., 18, 1985. Niihara, K. (1991) J. Ceram. Soc. Jpn., 99, 974. Niihara, K., Izaki, K. and Kawakami, T. (1990) J. Mater. Sci. Lett., 10, 112. Pan, X., Mayer, J., Rühle, M. and Niihara, K. (1996) J. Am. Ceram. Soc., 79, 585. Pezzoti, G. and Sakai, M. (1994) J. Am. Ceram. Soc., 77, 3039. Rendtel, A., Hübner, H. and Herrmann, M. (1995) Fourth Euro Ceramics, vol. 4, (ed. A. Bellosi, Gruppo Editoriale Faenza Editrice), Italy, p. 225. Rendtel, A., Hübner, H., Herrmann, M. and Schubert, C. (1998) J. Am. Ceram. Soc., 81, 1109. Rossignol, F., Rouxel, T., Besson, J. L., Goursat, P. and Lespade, P. (1995) J. Phys. III, 5, 127. Rouxel, T., Wakai, F. and Izaki, K. (1992) J. Am. Ceram. Soc., 75, 2363. Rouxel, T., Rossignol, F., Besson, J. L. and Goursat, P. (1997) J. Mater. Res., 12, 480. Sajgalik, P., Dusza, J., Hofer, F., et al. (1996) J. Mater. Sci. Lett., 15, 72. Sasaki, G., Nakase, H., Suganuma, K., Fujita, T. and Niihara, K. (1992) J. Ceram. Soc. Jpn., 100, 536. Sato, S., Chu, M. C., Kim, J. H., Kobayashi, Y. and Ando, K. (1995) J. Ceram. Soc. Jpn., 103, 676.

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Page 281 Sawaguchi, A., Toda, K. and Niihara, K. (1991) J. Am. Ceram. Soc., 74, 1142. Stocker, R. L. and Ashby, M. F. (1973) Rev. Geophys. Space Phys., 11, 391. Tanaka, I., Igashira, K., Kleebe, H. J. and Rühle, M. (1994) J. Am. Ceram. Soc., 77, 275. Xu. Y., Zangvil, A. and Kerber, A. (1997) J. Eur. Ceram. Soc., 17, 921. Zhao, J., Stearns, L. C., Harmer, M. P., Chan, H. M., Miller, G. A. and Cook, R. F. (1993) J. Am. Ceram. Soc. , 76, 503.

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Page 282 9 Theoretical investigations of Si/C/N-based alloys Fabio Finocchi 9.1 Introduction Numerical simulations are a powerful tool for the study of the physical and chemical properties of various systems, as shown by the increasing number of scientific publications in which this technique is used.* The flexibility of the model representing the physical system makes it possible to perform several numerical experiments, as well as to assess the influence of the various parameters on the final results. On one hand, the information obtained can complement the experimental data and help the experimentalist to draw microscopic models which will be able to explain the observed behaviour. On the other hand, numerical experiments can be used as benchmarks for approximated analytical theories on model systems. Last but not least, the tremendous advances in computing power make numerical simulation an easily accessible technique and open the way to the simulation of bigger systems with a higher degree of complexity and closer to real materials. This fact, in conjunction with the progresses that have been achieved in the construction of reliable theoretical models, has largely contributed to the computer-aided design of novel materials. However, the contribution of numerical simulations to the investigation of the mechanical, structural and electronic properties of the Si—C—N based system is presently rather thin. Three main factors have so far limited the applications: (i) the gap between the relevant experimental time scales and the simulation time; (ii) the size of the system which mimics the real sample; and (iii) the structural and chemical complexity of the alloys. Regarding (i), the microscopic structure of the SiCN nanopowders is sensitive to preparation conditions. They are generally obtained through laser synthesis from organic precursors (Cauchetier et al., Chapter 2). The relative abundance of the constituents, and other thermodynamic parameters (such as the temperature), has been used to design nanopowders with specific characteristics. The time scales involved in such processes are by many orders of magnitude bigger than those that are presently affordable in simulations, which usually range from some picoseconds to a few tens of nanoseconds. Therefore, in order to bypass * In molecular dynamics simulations, the individual trajectories of particles (atoms, molecules) are followed over a period of time, making full allowance for their mutual interactions. The thermodynamic properties are computed as time averages over the trajectories. Int he MonteCarlo method, a sequence of random moves in the configuration space of the system of particles is made, according to a statistical weight. It is chosen in such a way that the probability of each configuration be asymptotically equal to its statistical weight in some ensemble (canonical, macro-canonical, etc.). Some aspects are common to both molecular dynamics and Monte Carlo simulations: the choice of the inter-particle potential energy, the treatment of finite-size effects, and the selection of the starting configuration. For more details, the reader can look to the book by Allen and Tildesley (1987), while an account of recent developments can be found in vol. 106 of the Faraday Discussions, in the book edited by Meyer and Pontikis (1991), and in the review by Baskes (1999).

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Page 283 the stage of the simulation of the full history of the real material, a good hint on its structure at the atomic scale is often crucial. When amorphous materials are considered, such a hint is not generally trivial. A practical alternative is to generate the metastable phase through a rapid quench from the melt. The quench rate appears as the most relevant parameter which influences the disordered structure obtained at low temperature. It is impossible to cite all the literature devoted to the subject. However, in some cases, the use of a quench rate that is many orders of magnitude larger than its experimental counterpart, had not prevented the results of the simulation to be comparable to the experimental data (Servalli and Colombo 1993; Wang et al. 1993; Vollmayr et al. 1996; Rosato et al. 1999). Regarding (ii), the simulation of bulk phases often makes use of periodic boundary conditions, which eliminate surface effects (Figure 9.1). In such a case, the limiting parameter is the number of atoms contained in the unit cell. The simulation of heterogeneous nanophases, or nanoassembled materials, may need of the order of a million atoms per unit cell, which is rarely affordable. Ab initio simulations (Car and Parrinello 1985, 1988; Oguchi and Sasaki 1991), in which the interatomic forces are derived from the knowledge of the electronic structure, are presently unable to treat more than some hundreds of atoms. One way out is to make the problem simpler, by choosing a model interatomic potential that is devised to fit the most important physical properties of the system under study. Since a model potential may be hardly transferable to many different physical context, one has to pay attention to its reliability in the actual conditions of the simulation. A meaningful example is provided by the simulation of the various phases of amorphous hydrogenated carbon, which is characterised

Figure 9.1 Sketch of a two-dimensional periodic system. The particles are free to cross any one of the box edge. Note that the knowledge of the positions and velocities of the particles inside one cell is sufficient to know the dynamic of all the other particles, by making use of the periodic boundary conditions. They are said to be the periodic images of the particles inside the simulation box.

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Page 284 by qualitative variations of the C—C and H—C bonding that deeply affects the macroscopic properties of the sample (Robertson 1991). The two previous issues are especially delicate points when the system is formed by distinct atomic species, each having a specific chemical behaviour. First, as the number of distinct atomic species increases, the number of bonding configurations as well as the number of fitting parameters in the interatomic potential grows up rapidly. Second, each bond is characterised by specific physical properties, such as the charge transfer, the bond angle, etc., which can depend on the nth neighbour configuration (n = 2, 3, …) in a nontrivial way. It is thus more difficult to assess the transferability and the reliability of the model interatomic potential energy as the complexity of the system grows. Third, the landscape of the corresponding thermodynamic potentials is generally rather complex, because of the presence of many local minima, eventually separated by huge barriers. The dynamic simulation of phase separation, or of frustrated systems, are examples in which the complexity of the problem results from the competition between atoms with different chemical behaviours. As a consequence, the simulation of disordered alloys is particularly challenging. Presently, the use of ab initio based molecular dynamics simulations does not permit the treatment of more than some hundreds of atoms, and to follow their dynamics for more than few tens of picoseconds. On the other hand, the fitting of an effective model interatomic potential and its validation, through transferability tests, may require a formidable effort. All the previous considerations apply to the case of systems very close to thermodynamic equilibrium. When one is interested in kinetics, specific techniques, able to cope with rare events, such as the overcoming of a free energy barrier, should be used. They are in general quite complex and more time consuming than the usual methods for simulations of (meta-)stable systems. For a critical discussion of these methods, the interested reader can look to the review written by Ciccotti in the book edited by Meyer and Pontikis (1991). 9.2 Simulations of Si/C and Si/N disordered alloys Both amorphous silicon carbide (α-Si1−xCx) and silicon nitride (a-SiNy) are materials which have attracted much attention. From a technological point of view, they are useful as hard coatings, resistant to corrosion and chemically inert. α-SiC has also been used as p-type window layer in amorphous-silicon based solar cells. At odds with their crystalline counterparts, they can exist in many different stoichiometries. It is therefore relevant to understand the influence of the stoichiometry on the mechanical and electronic properties of these amorphous alloys. For tiny C (N) concentrations, the proposed microscopic models consider the C (N) atoms as impurities hosted in the α-Si network (Wooten and Weaire 1987; for a critical discussion see also Mousseau and Lewis 1997). Such a picture is consistent with most of the measured structural properties. On the other hand, for compositions close to those of the crystals (C/Si = 1; N/Si = 4/3), the amorphous networks have been built from the corresponding ordered structures through the deformation of bond angles and the lack of long-range order. Indeed, the Si—C(Si—N) bond lengths are remarkably similar in both amorphous and crystalline phases. Apart from above mentioned values of C/Si (N/Si) concentrations, it is a priori not evident in any known reference system, it is useful to understand the topology and the connectivity of the network of the amorphous alloy that could exist. Thus, much room is left to numerical simulations, to propose structural models that may be able to account for the experimental data collected for the α-Si1−xCx and a-SiNy alloys.

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Page 285 9.2.1 Chemical order Before starting to discuss the results obtained from the numerical simulation, however, it is important to define the concept of chemical order. Since the SiC and Si3N4 crystals (the actual phase is not important for the sake of the following discussion, provided that it is close enough in energy to the most stable phase at normal conditions) are stable, the formation of heteronuclear Si—C (Si—N) bonds is always energetically favoured on the formation of Si—Si and C—C (N—N) bonds within two separated pure phases. Therefore, an amorphous alloy is said to be totally chemically ordered when the number of heteronuclear bonds are as large as possible. In disordered materials, the coordination number of the species A with the species B is usually defined as the integral of the A—B pair distribution function up to the first minimum (see Figure 9.3). Such a definition is adopted in the following discussions. However, one should pay attention to the presence of dilated A—B bonds, as well as to slow structural relaxation, which reflect in a non-zero value of the first minimum. Strictly speaking, only a characterisation of bonds in terms of the electronic distribution (Bader 1991) is nonambiguous. Although the concept of chemical order is very clear and intuitive, the choice of the order parameter is not unique and may hinder some important physical aspects of the question (Cargill and Spaepen 1981). For this purpose, let us define the following quantities, for a binary A—B alloy. NA, (NB): number of atoms of the A (B) species. ZA (ZB): average coordination number of the A (B) species. NAA, NAB, NBB: number of A—A, A— B, B—B bonds, respectively. Obviously, these quantities are not independent. They are linked by the following equations: Let us express, in the case of A1−xBx, the degree of chemical order in terms of the above-defined quantities. In the case of full disorder, given an A-type atom, the a priori probability to form an AA bond is equal to that of forming an AB bond. Therefore:

At odds, in the case of perfect chemical order, a type A atom always forms an AB bond whenever type B atoms are available. Therefore: and where x0 is the critical concentration, equal to x0 = ZB/(ZA + ZB) In the case of a-Si1−xCx, a currently used order parameter is τ = NCC/NSiC. One can see that for a fully disordered system τ = (ZC/ZSi)x/(1 − x), which is in general smaller than x/(1 − x), since Si has a larger average coordination number than C. For a perfectly ordered systems, τ = 0 (x < x0); (ZC/ZSi)x/(1 − x) − 1 (x > x0). Apart from the fact that τ is ill defined for very high C concentration (x → 1), one can see that τ does not purely depend on the C concentration x, but also on the ratio of the average coordination numbers of C

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Figure 9.2 Order parameter τ = NCC/NCSi computed for two virtual Si1−xCx systems, having all the C and Si coordination numbers equal to ZC = 3.5 and ZSi = 4.2, or ZC = 4.0 and ZSi = 4.0, respectively. Both cases are representative of real amorphous samples. The cases of full disorder and perfect chemical order are plotted. Note that the curve corresponding to the disordered alloy for ZC< ZSi may be close to that of the completely ordered Si1−xCx system with both C and Si atoms fourfold coordinated. and Si. This is not a purely academic question, since an appreciable amount of graphitic-like bonding implies ZC< ZSi · Even when other order parameters, which are free from the singularities at x = 0 or x = 1, are adopted, such as ζ = NCC/(NCC + 2NSiC + NSiSi), the degree of order is invariably a function of both C concentration and C average coordination number, compared to that of Si. Also the critical value of the C concentration (x0), at which some homonuclear bonds appear even in the case of perfect chemical order, is a function of the average C and Si coordination numbers (Figure 9.2). The previous discussion is noteworthy, since many experimental data were discussed on the basis of a simplified model for the chemical order, in which ZC = ZSi = 4, even if sp2 C bonding was actually detected in the amorphous sample. Moreover, in the case of strong departure from the chemical order, the crystalline, tetrahedral phase of SiC cannot be used as a starting model for the microscopic structure of amorphous Si—C alloys. Therefore, in the following, we prefer to discuss the observed trends for the chemical order, keeping in mind that the precise values of the order parameter cannot be always considered as a definitive measure of the degree of chemical order in the absence of an independent estimate of both C and Si average coordination numbers. 9.2.2 Amorphous silicon carbide Kelires (1991) carried out a Monte Carlo simulation of amorphous SiC, by adopting an effective interatomic potential energy introduced by Tersoff (1989). The model potential

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Page 287 for SiC is derived from a two- and three-body interaction originally designed for elemental semiconductors, and then modified in order to take into account the chemical specificity of C and Si through environmentdependent parameters. Kelires found a moderate trend towards chemical order (τ = 0.5 at C/Si = 1). A non-negligible fraction of C atoms are threefold coordinated. In the following first-principles molecular dynamics simulation, in which the interatomic forces are derived on quantum mechanical grounds, Finocchi et al. (1992, 1993) obtained α-SiC at stoichiometric composition and density close to that of the β crystalline phase, through a quench from the melt. The final sample, at room temperature, has a complex structure which can be classified neither as chemically ordered nor as completely random. Although the Si—Si, Si—C and C—C equilibrium bond lengths (characterised by the first peaks in the pair correlation functions) are very close to those in the corresponding crystals, they are statistically distributed over a different range of distance, depending on the specific bonding. For instance, the width of first peaks of the C—C, Si—C and Si—Si pair correlation functions increase in the order, showing a greater disorder on the Si sub-lattice than on that of carbon (C). The resulting total radial distribution function, being a sum of the various contributions, may be difficult to be interpreted. As a consequence, only the access to the partial correlation functions by means of complementary and species-sensitive techniques, allows one to disentangle the specific contributions. Although the simulated system may strongly differ from the real amorphous samples, we think that such an ambiguity is characteristic of semi-conducting alloys with (i) a large degree of disorder and (ii) competition between different type of bonding (e.g. sp2/sp3 for C). In a following paper, Tersoff (1994) focused on the influence of the enthalpy of formation of Si—C bonds from Si Si and C—C bonds (ΔHSi—C), on the degree of chemical order. In the crystal, ΔHSi—C is equal to −16 kJ/mol/bond, according to the experimental data at normal conditions. Tersoff performed Monte Carlo simulations of amorphous Si1−xCx, using different interatomic potentials, which correspond to distinct ΔHSi— C. In that way, he studied the order parameter τ as a function of ΔHSi—C. He found that both ΔHSi—C and atomic size difference between Si and C drive the system towards the chemical order. However, the steric factor is much more effective at driving ordering in the crystal than in the amorphous alloy, since in the latter the microscopic stress resulting from the formation of homonuclear bonds can be more efficiently released than in the long-range ordered phase. All these arguments are valid at not too high temperatures, when the entropy contribution, which usually favours disorder, can be neglected. Further ab initio molecular dynamics simulations (Finocchi and Galli 1994) show that the inclusion of hydrogen in the α-SiC network, at small concentrations, favours tetrahedral bonded carbons. Since a trend of sp2-like carbons to segregate has been observed in the ab initio simulations – as a result of the strong C—C bonding – the authors concluded that the inclusion of H promotes chemical order, towards values of the order parameter (τ = 0.5) consistent with most experimental data (Figure 9.3). A recent molecular dynamics simulation (Mura et al. 1998) studied the influence of the quenching rate and the C/Si ratio on the final chemical order. The use of a modified version of the Tersoff effective potentials (instead of the more accurate, but time-consuming first-principle-based interatomic potential energy) allows the authors to carry out much longer simulations. Apart from quantitative differences from the results of the ab initio simulations, they also found that the Si sub-lattice can be easily distorted, a fact which attenuates the preference for a chemical order, which is far from being complete in the network. Some of the previous results can be discussed on the basis of simple arguments based upon the thermodynamic data, which, in the case of silicon carbide, drive the metastable

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Figure 9.3 Pair distribution functions for α-SiC : H, as computed by Finocchi and Galli (1994). They are defined as gij (r) = V/(NiNj)nij (r)/(4πr2r) where nij denotes the time-averaged number of atomic species j in a shell between r − r/2 and r + r/2 centred on a type i atom. V is the volume of the simulation box and Ni the number of atoms of type i inside V . From the gij , the total pair distribution function is obtained as: G(r) =∑ij Ni Nj gij (r)/(∑i Ni)2. For the equations connecting gij (r) with the X-ray, neutron diffracted intensities (Finocchi and Galli 1994; Waseda 1980). amorphous phases towards the chemical order. However, three factors conspire to maximise disorder: i the entropy contribution, which is especially important at high temperatures; ii the presence of a remarkable disorder on the Si sub-lattice; indeed, the absolute value of the energy of formation of two stretched Si—C bonds from a C—C and a stretched Si Si bond, may be reduced in this case; iii the presence of sp2-like carbon, which reinforces C—C bonding, lowers the average C coordination number, and distorts the tetrahedral network. Although the energetic terms should prevail for small C concentrations, and drive the α-SiC towards chemical order, the C-rich systems might exhibit a complex behaviour, which depends on the macroscopic thermodynamic parameters (density, temperature, H contents, etc.) as well as on the preparation conditions and on the history of the sample. 9.2.3 Amorphous silicon nitride The enthalpy of formation of crystalline silicon nitride from Si and molecular nitrogen is equal to −828 kJ/ mole (O’Hara et al. 1997). The degree of chemical order expected for amorphous

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Page 289 Si3N4 and SiC on the basis of purely energetic arguments can thus be compared by considering the enthalpy of formation of one heteronuclear Si—N(Si—C) bond, obtained by normalising the enthalpy of formation of the crystalline compounds from the respective elemental phases, by the number of heteronuclear bonds. This amounts to −0.70 eV per Si—N bond, and to −0.16 eV per Si—C bond. The absolute values of the two enthalpies correspond to very different temperatures, of about 8000 and 2000 K, respectively. This crude estimation shows that, according to thermodynamics, a-Si3N4 should be almost perfectly chemically ordered. Moreover, the segregation of a Si-rich phase might be accompanied by the formation of highly diffusing N dimers, resulting in a continuous N loss. In agreement with this framework, all the simulations so far performed on a-Si3N4 indicate a high degree of chemical order, with small amounts of homonuclear bonds. Umesaki et al. (1992) carried out molecular dynamic runs on amorphous Si3N4, using a simple Busing-type potential, which consists of a point charge Coulomb interaction, plus a short-range repulsive term. The parameters were fitted to reproduce the α and β crystalline phase of silicon nitride. The density was fixed at 2.6 g/cm3, and 420 or 1260 particles were used in the simulation. The structure of the simulated samples was characterised in terms of the partial pair distribution functions. The computed Si N mean bond length is equal to 1.74 Å, and the Si N first coordination number is slightly lower than 4 (3.95). The second neighbour distances (the corresponding coordination numbers are given in brackets) are 2.65 Å (8.3) and 3.05 Å (7.1), for N—N and Si—Si, respectively. All these numbers are in very good agreement with the X-ray and neutron diffraction experiments. The authors conclude that the short-range order of amorphous Si3N4 is very similar to that of the crystals. However, the presence of some dangling bonds, and the distortions of the dihedral angles between the SiN4 tetrahedra result in the loss of the medium-range order with respect to the crystal. Recently, de Brito Mota et al. (1998, 1999) simulated a-SiNx systems (for 0 ≤ x ≤ 1.5) through Monte Carlo runs in the constant-temperature constant-pressure (NPT) ensemble. They used interatomic potentials of the Tersoff type, which were fitted to crystalline silicon nitride, as well as to molecules containing Si, N and H. The inclusion of N in the a-Si network favours denser metastable phases, since the equilibrium density grows from 2.31 g/cm3 at x = 0 to about 3.5 g/cm3 at x = 1.5. This trend compares well with the available experimental results (Guraya et al. 1990). It is a consequence of the formation of Si N bonds, at the expense of the longer Si Si bonds, as N is incorporated in the amorphous sample. Correlated with such a behaviour, the angular disorder around Si sites grows as a function of the N concentration x. a-SiNx can be considered as well chemically ordered, although a non-negligible number of Si Si bonds are detected (τ = 0.15 for x = 1.33, which corresponds to the stoichiometric composition), in accord with the interpretation of the X-ray photoemission spectra for SiNx powders (Sénémaud et al. 1993). The inclusion of hydrogen reduces the number of dangling bonds, releases the microscopic stress and improves the conductivity properties. However, no major structural changes are recorded in the Si—Si and the Si—N pair distribution functions, between the hydrogenated and the unhydrogenated simulated systems. They are shown in Figure 9.4, for SiN1.33. Given the relative success of the empirical potentials in describing a-Si3N4, Vashista and collaborators (Vashista et al. 1995; Kalia et al. 1997) performed large-scale molecular dynamics simulations on disordered silicon nitride phases, adopting a slightly more complicated interatomic potential which, in addition to the Coulomb charge–charge interaction, and to the short-range repulsion, contains a charge–dipole interaction. In a first simulation (Vashista et al. 1995), they prepared the amorphous samples (consisting of 1344 atoms) by quench from the melt, at different macroscopic densities. As the density decreases, the

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Figure 9.4 Total pair correlation functions for amorphous SiNx. Copyright (1998) by the American Physical Society for x = 0.5, 1.0, 1.33 and 1.5. The experimental results for a-SiN1.33 (Guraya et al. 1990) are shown in the inset for comparison. (Reprinted with permission from de Brito Mota, F., Justo, J. F. and Fazzio, A. (1998) Phys. Rev., B58, 8323–8328.) network connectivity is reduced and floppy modes appear in the vibrational spectra. The consequences for the mechanical and thermal properties of a-Si3N4 are discussed. In a following study, they assembled preformed Si3N4 clusters to form a nanophase Si3N4 system, containing about a million atoms. The system was characterised as composed by (1) interfacial regions with a large fraction of undercoordinated Si, which separate (2) bulk-like domains from (3) the pores (three-phase model). The dependence of the computed elastic moduli on porosity and grain size, as well as the behaviour of the material under crack propagation, can be understood on the basis of the three-phase structural model. In particular, the nanophase system is able to sustain an order-of-magnitude larger external strain than crystalline Si3N4. 9.3 Simulations of ternary Si/C/N systems As previously discussed in the Introduction (Section 9.1), the simulation of disordered ternary systems is a formidable task, which has limited so far the application of atomic-scale simulation techniques to the study of amorphous or poly-crystalline Si—C—N systems. In the following, we mainly review the recent theoretical advances in the field of crystalline Si—C—N systems, and in the study of disordered (amorphous or nanopowders) alloys.

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Page 291 9.3.1 Theoretical studies of C defects in Si3N4and N defects in SiC In this section, we give a brief account of the theoretical work that has been carried out to study the structural and electronic characteristics of the inclusion of either C or N in the crystalline network of the crystal formed by the two other species. In particular, we account for the most relevant modification of the structure and the electronic properties resulting from the inclusion of the substitutional impurity at very low concentrations, as a first step towards the understanding of the formation of ternary Si—C—N compounds. They are not necessarily the most stable point defects in these compounds. For example, in SiC, the SiC antisite defects is more stable than both NSi and NC (Bernholc et al. 1989). The reader who is interested to the more general issue of point defects in semiconductors can look to the available reviews (see, for instance, Lannoo 1992). Fukumoto (1997) studied the inclusion of N in 3C—SiC within the density functional theory (DFT) in the local density approximation (LDA) for the exchange and correlation effects (for a review on the method, see Jones and Gunnarsson 1989). Super-cells containing 64 atoms in total were used, and the relaxation of the atomic positions were performed through damped dynamics. However, in those calculations, the volume of the super-cell is not varied as a function of the position of the impurity, so that a spurious, residual pressure might somehow affect the energetic, since finite N concentrations are considered. Fukumoto considered only substitutional impurities, that is, NSi and NC. NC gives rise to a shallow donor state, with electron density widely spread over the entire crystal, and a maximum close to off-bond sites around Si. NSi, on the other hand, forms a deep level, characterised by a very localised electron distribution around the substitutional N. He also studied the formation energy of the defects, as a function of the Si chemical potential, and found that NC is always favoured over NSi by about 7 eV, in the whole physical range of Si chemical potentials compatible with the existence of stoichiometric silicon carbide. Amadon and Finocchi (1999) recently studied, within the same theoretical framework (DFT in the LDA) CSiβSi3N4, at different C concentrations (1 or 2 carbons in a supercell containing 84 atoms). Both the atomic and the external degrees of freedom (i.e. the lattice parameters) are consistently relaxed for all the structures considered. For the lowest CSi concentrations, the C—N bonds are stretched with respect to those computed in β-C3N4. When doubling the CSi concentration, scattered configurations of CSi are favoured over geometries with second neighbours CSi. The resulting microscopic stress is long ranged and lowers the cohesive energy of the host crystal. As a consequence, the formation of CSi defects in β-Si3N4 from diamond and silicon nitride crystals, and molecular nitrogen is thermodynamically unfavoured, showing that the progressive substitution of Si with C in silicon nitride is not a practical way towards the synthesis of ternary alloys. 9.3.2 Crystalline Si/C/N compounds Si—C—N ordered systems have only recently been obtained. They have been initially prepared in nanocrystalline phases by Riedel et al. (1997) through the reaction of silicon tetra-chloride with trimethylsylcarbo-diimide (CH3)3—Si—N=C=N—Si (CH3)3, followed by calcination. The resulting amorphous Si—C—N powder crystallises at 400°C to give SiC2N4. The latter, when heated up to 900°C, decomposes in crystalline Si2CN4, C2N2 and N2. These phases were studied by in situ X-ray diffraction. Si2CN4 shows a layered structure, formed by distorted SiN4 tetrahedra as in Si3N4, which are linked through carbodiimide groups. It has an orthorhombic unit cell (lattice parameters: a = 5.44 Å, b= 13.58 Å and c = 4.81 Å),

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Figure 9.5 Structure of Aaba2 Si2CN4. N atoms are dark grey, Si light grey and C white. The three directions, parallel to the lattice vectors, are drawn. Note that in the [100] and in the [001] directions, the structure is replicated, in order to get a better feeling of the connectivity of the crystal. with space group Aaba2 (Figure 9.5). However, the internal parameters defining the precise atomic positions were only roughly determined. The structure of SiC2N4 was obtained too, but with less accuracy. The stability of these structures has been theoretically investigated by Kroll et al. (1999) through density functional calculations. They exploited the functional analogy between the carbodiimide and the O groups to build up various structural model compatible with the X-ray diffraction data. The cubic high-temperature β phase of SiC2N4 consists of a Si atom at the centre of a tetrahedron formed by the four NCN groups. However the authors, thinking that the cubic phase may only reflect an average geometry, searched for other less symmetric structures. Two tetragonal structures resulted to have a higher cohesive energy than the cubic phase. The most stable one can be described as two interpenetrating Si—N=C=N—Si networks. It has a c/a ratio equal to 2 and a space group P4322. Such a structure is only 4 KJ/mol more stable than the ideal cubic one, an energy difference which is close to the intrinsic precision of the theoretical method employed. However, the P4322 structure was proposed as a good candidate for the low-temperature α phase of SiC2N4, not only for energetic reasons. Indeed, the energy plotted as a function of the bonding angle of N shows a shallow minimum around 160° for the tetragonal structures, at odds with its 180° value in the cubic structure. The tetragonal structure has thus more flexibility, with two consequences: its computed bulk modulus (about 8 GPa) is much smaller than that of the cubic phase (108 GPa);

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Page 293 at higher temperatures, the tetragonal structure might appear as cubic on average, due to the easy deformation of the bonding and dihedral angles. As far as the crystalline Si2CN4 system is concerned, two theoretical studies appeared almost simultaneously: that by Kroll et al. (1999), and another one by Amadon and Finocchi (1999). Both of them found that the orthorhombic structure with space group Aaba2 proposed by Riedel et al. (1997) is a mechanically stable phase. The computed lattice parameters agree within less than 2% with those extracted from the X-ray diffraction data: a = 5.51 Å, b = 13.75 Å, c = 4.83 Å (Amadon and Finocchi 1999); a = 5.45 Å, b = 13.81 Å, c = 4.82 Å (Kroll et al. 1999); a = 5.44 Å, b = 13.58 Å, c = 4.81 Å (Riedel et al. 1997). Apart from a small discrepancy about the b/a ratio, both calculations found a larger equilibrium volume (by 2–3%) than the experiment. This is in contrast with the outcome of most density functional calculations adopting the LDA, which usually gives a trend to over-binding and to slightly underestimated lattice parameters – by 1–2% in elemental and compound semiconductors (Jones and Gunnarsson 1989). Due to the presence of very soft modes, essentially linked to the easy deformations of the Si—N=C angles, a non-monotone behaviour of the phonon frequencies v. the volume cannot be excluded, which might give a non-trivial value of the thermal expansion coefficient. However, for the moment, we are still at the level of speculations. An important consequence of the coexistence of more rigid Si—N layers linked with deformable N=C=N sticks, into a highly non-isotropic structure, is the relevance of the relaxation of the internal degrees of freedom (i.e. the atomic positions inside the unit cell) for the calculation of the mechanical properties of these phases. From the computed elastic constants, Amadon and Finocchi calculated a bulk modulus equal to 110 GPa. Experimentally, Si2CN4 decomposes at high temperatures in silicon carbide, silicon nitride, plus molecular nitrogen (reaction 1). Both calculations give a positive reaction enthalpy at 0 K, ranging from 215 to 285 kJ/ mole. However, another possible reaction is possible, giving solid C, silicon nitride and molecular N as products (reaction 2). The path from the products in reaction 1 to those in reaction 2 implies a negative enthalpy, both experimentally and theoretically, of the order of about −280 kJ/mole. This raises the question of the stability of Si2CN4 with respect to reaction 2. While Amadon and Finocchi found a reaction enthalpy in the range −20 to +10 kJ/mol, the latter was estimated at −39 kJ/mole by Kroll et al., who concluded for the thermodynamic instability of Si2CN4. However, Amadon and Finocchi were less conclusive on this point, since (i) the reaction enthalpies are likely to be too low by few tens of kJ/mole in the case of reaction 2, as a result of the use of the LDA, and (ii) the presence of high barriers and entropy contributions should favour the decomposition of Si2CN4 in silicon carbide, silicon nitride, plus molecular N, as in reaction 1. Crystalline Si—C—N thin films were recently grown by microwave plasma-enhanced chemical vapour deposition on the (100) surface of Si (Chen et al. 1996), and then studied by X-ray absorption spectroscopy (Chang et al. 1998). The X-ray diffraction data were interpreted on the basis on a crystalline structure analogous to that of the α phase of Si3N4, from which it differs for the values of the lattice parameters (those of α-Si3N4 are given in brackets): a = b = 6.908 Å (7.753 Å; c = 5.262 Å (5.618 Å). The XANES spectra of the thin Si—C—N films, taken at the C, N and Si K edge, showed marked similarities with those of the β phase of SiC, the α phase of Si3N4, and the combination of the two latter phases, respectively. Moreover, from the EXAFS oscillations at the Si K edge, Chang et al. derived an average Si—N coordination equal to 3 (with a Si—N bond length equal to 1.73 Å) and Si—C mean coordination number equal to 1 (Si—C bond length: 1.88 Å). These facts suggest that the local bonding environments of C and N in the ternary thin films are rather similar

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Page 294 to those in the binary crystals formed with Si (β-SiC and α-Si3N4). Si, at variance, shows a dominant Si—N bonding, but the number of Si—C bonds is not negligible. All the experimental results were complemented by theoretical calculations in the muffin-tin approximation for the scattering potentials provided by the ions. From these results, the authors suggested a structural model based on mixed C—Si—N3 tetrahedral units in which C—N bonds should represent a very small fraction of the total number of bonds. These observations raise a few interesting issues. First, the thin films are said to be crystalline, but their composition was not precisely estimated. As a rough indication, under the hypothesis that C and N atoms are fourfold and threefold coordinated to Si, respectively, and that Si atoms form mixed C—Si—N3 tetrahedra, one obtains a C/Si ratio equal to 1/4, and a N/Si ratio equal to 1. Thus, the progressive inclusion of C atoms in an hypothetical α-Si3N4 network, should be made by substitution of N with C, with a much smaller proportion of Ni replaced by Si atoms, which depends on the precise concentration of the three components. As a consequence of the difference of Si—N and Si—C bond lengths, a microscopic stress at the atomic scale would result, which may strongly influence the ground state properties of the crystals, such as its stability, mechanical behaviour and electronic structure, as in the case of CSi Si3N4 and NC—SiC, which have been discussed in Section 9.2. This observation is also coherent with recent experimental results reported for amorphous CSixNy thin films (x ~ 0.2, y ~ 0.3) which were grown on (100) oriented Si substrates, by laser deposition (Thangen et al. 1999). The films are hard and plastic, and show an internal compressive stress of the order of a few GPa. A reversed approach has been undertaken by some theorists, who searched for super hard Si—C—N structures (for a review, see Badzian and Badzian 1997), obtained by progressive replacement of Si by C in βSi3N4. On the basis of an empirical model, Cohen (1985) suggested that the structures characterised by tetrahedral units with short (and strong) covalent bonds are good candidates for materials with a high bulk modulus. Among them, the hypothetical solid β-C3N4, iso-structural to the β phase of Si3N4, should have a very large bulk modulus, comparable to that of diamond. Afterwards, some calculations (Liu and Cohen 1989; Liu and Wentzcovitch 1994) studied, on first-principle grounds, the stability and the mechanical properties of C3N4 solids. In particular, through a variable-shape ab initio molecular dynamics simulations, Liu and Wentzcovitch (1994) confirmed the very low compressibility of the β phase (space group P63/m). Such a crystal would have a short C—N bond length (1.45 Å) and a bulk modulus equal to 437 GPa, just smaller than that of diamond. However, the authors found a more stable rhombohedral structure (space group R3m), about 0.15 eV/unit formula lower in energy than the β-C3N4 phase. The rhombohedral phase is more similar to graphite than to diamond, with a C—N bond length equal to 1.36 Å on average, and a quite low bulk modulus (51 GPa). Both phases are thermodynamically unstable vis-à-vis of the decomposition in diamond plus molecular N (the enthalpy of reaction ranging from −75 to −62 kJ/mol). However, as in the case of diamond, they found both phases mechanically stable, and did not exclude that they might be synthesised. These seminal studies motivated the search for hypothetical Si—C—N tetrahedrally bound compounds, obtained through the replacement of Si with C in Si3N4. Among them, Wang et al. (1998) proposed two model structures of β-Si2CN4 and β-SiC2N4 on the basis of density functional calculations in the full potential linear augmented plane wave method. They calculated the equilibrium structural parameters and the bulk modulus of the hypothetical structures, under symmetry restrictions. However, these calculations did not provide any information on the thermodynamic and mechanical stability of these compounds. Then, Lowther (1999) compared the β and the cubic phases of hypothetical SiC2N4 with C3N4. He

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Page 295 found that the electron distribution along the C—N bonds in SiC2N4 is rather similar to its counterpart in βC3N4, which suggests a remarkable stability of the strong C—N bonding in various chemical environments. Moreover, he found the cubic (P4n2) phase more stable than the β phase by a very tiny energy. Once again, the comparison of the cohesive energies of these ternary phases with respect to elemental or binary crystals was not done. 9.3.3 Si/C/N nanopowders and amorphous alloys It is noteworthy that a fit to the X-ray or neutron diffraction spectra performed on Si—C—N nanopowders at different temperatures gives results similar to the binary crystals, regarding the Si—N and Si—C bond lengths (see Chapter 5). In analogy with the results by Chang et al. on the Si—C—N crystalline thin films, Ténégal et al. (1996) found that the local bonding structure of amorphous silicon carbonitride powders, as probed by EXAFS, may be described in term of mixed C—Si—N3 tetrahedrons, but, at the same time, evidence for C—N bonds connecting the tetrahedrons was also found. The same samples were further analysed by X-ray and neutron diffraction. The experimental spectra were interpreted (Ténégal et al. 1998) on the basis of a medium-range order based on different stacking of the Si/C/N layers, reminiscent of either the α or β phase of crystalline silicon nitride. Also in that case, the Si3N4 crystalline phases were used as input to build up structural models through: (i) replacement of Si or N atoms by carbons and (ii) a local optimisation of the structure by using empirical force fields. One can note that all the models derived through the selective replacement of C atoms in a silicon nitride crystalline network have the virtue of fitting the experimental data remarkably, thereby providing an atomic scale picture of the configurations of ternary Si—C—N systems. However, they are not complemented by any thermodynamic data, or by the calculation of the electronic properties for the corresponding atomic arrangement. As a consequence, they are neither able to explain the Si—N, C—N and Si—C bonding properties within the disordered network, nor understand the thermodynamic evolution of the Si—C—N systems. Moreover, they have not been showed to be coherent with the measured electronic structure of the ternary amorphous or crystalline Si—C—N systems. The comparison between photoemission experiments (Driss-Khodja et al. 1996) and the electronic structure computed for the proposed structural models would be thus significant to assess their validity. To our knowledge, Matsunaga et al. (1999) were the first to simulate truly amorphous Si3CzN4−z ceramics. They employed a Tersoff potential to describe the C—C, Si—S, Si—N and Si—C interactions. Both C—N and N —N interatomic potentials are purely repulsive. The justification of such a simplified model was given on the basis of the high mobility of N2 dimers – which do not contribute to the bonding in the condensed phase – as far as the N—N interaction is concerned. The C—N attractive interaction was not taken into account since a decomposition of C—N bonds is observed during pyrolysis above 1273 K (Seher et al. 1994). The molecular dynamics simulations were performed in the canonical (NVT) ensemble, by generating the amorphous configurations through a rapid quench from a liquid-like structure at ~4000 K. The statistic was taken on the sample equilibrated at T ~ 1270 K for about 50 ps. Several C concentrations z were simulated. The results were characterised in term of partial pair distribution functions. For Si3C1.5N2.5, both N and C atoms are undercoordinated (ZN ~ 2.4; ZC ~ 3.1), at odds with Si (ZSi ~ 4.2), which are bonded to 2 C atoms, 1.4 N atoms and 0.8 Si atoms, on average. As more C is added, more and more Si—C bonds are created at the expense of Si—N bonds.

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Page 296 The overall picture of the structure of the amorphous sample cannot in any case be characterised as purely consisting of mixed C—Si—N3 and C2—Si—N2 tetrahedra, since both the presence of Si Si bonds, and the low coordination number of N and C reveal a bigger complexity. However, the computed total pair distribution function for Si3C1.5N2.5 (Figure 9.6), summed on the different bonding configurations, fairly compares to that measured by X-ray diffraction (Uhlig et al. 1996) for Si3C1.8N2.7.

Figure 9.6 (a) Computed total distribution function for amorphous Si3C1.5N2.5, by Matsunaga et al. (1999). (Reproduced from kind permission of Ceramic Society of Japan.) (b) The comparison with the experimental one (Uhlig et al. 1996) is also shown. The first peak marked in the figure corresponds to both Si—N and Si— C bonds. The second one is due to Si—Si bonds, while the third peak comes from C—C, N—N and N—C second neighbours. file:///C|/Users/steppenwolf/Documents/educatio...ks/to%20burn/nanocomposites/files/page_296.html (1 of 2) [20/01/2010 10:06:06 AM]

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Page 297 Such a simulation raises the interesting issues of the presence of (minority) Si—Si and C—N bonding, as well as of C and N dangling bonds, eventually saturated by hydrogen, which can be studied further. Also, the role of the formation of micro-voids may help the release of the internal stress and the measured mechanical properties. In conclusion, the available theoretical models give a still incomplete picture of Si—C—N systems. Apart from the theoretical and experimental investigations on carbodiimide-based Si—C—N crystals, some structural models are available, which fit the experimental data, without explaining why the proposed structures are (meta)stable. Some other theoretical investigations provide precise information on the energetic and the mechanical stability of hypothetical compounds, which have not yet been experimentally obtained. In very few cases the authors showed the way to a practical synthesis of these hypothetical compounds. A joint theoretical and experimental effort should be done in order to bridge the gap between the measured properties of these ternary systems and their explanation in terms of a unifying model. References Allen, M. P. and Tildesley, D. J. (1987) Computer Simulations of Liquids. Clarendon Press, Oxford. Amadon, B. and Finocchi, F. (1999) Eur. Phys. J., B11, 207. Baskes, M. I. (1999) Curr. Opin. Solid State Mater. Sci., 4, 273. Bader, R. W. F. (1991) Chem. Rev., 91, 983. Badzian, A. and Badzian, T. (1997) Int. J. Refractory Metals Hard Mater., 15,3. Bernholc, J., Antonelli, A., Wang, C., et al. (1989) Mater. Sci. Forum, 38/41, 713. Car, R. and Parrinello, M. (1985) Phys. Rev. Lett., 55, 2471. Car, R. and Parrinello, M. (1988) in Simple Molecular Systems at Very High Densities (eds A. Polian, P. Loubeyre and N. Boccara), Plenum Press, New York. Cargill III, G. S. and Spaepen, F. (1981) J. Non-Cryst. Solids, 43, 91. Chen, L. C., Yang, C. Y., Bhusari, D. M., et al. (1996) Diamond Relat. Mater., 5, 514. Chang, Y. K., Hsieh, H. H., Pong, W. F., et al. (1998) Phys. Rev., B58, 9018. Cohen, M. L. (1985) Phys. Rev., B32, 7988. de Brito Mota, F., Justo, J. F. and Fazzio, A. (1998) Phys. Rev., B58, 8323. de Brito Mota, F., Justo, J. F. and Fazzio, A. (1999) J. Appl. Phys., 86, 1843. Driss-Khodja, M., Gheorghiu, A., Dufour, G., et al. (1996) Phys. Rev., B55, 4287. Faraday Discussion (1997) Solid State Chemistry: New Opportunities From Computer Simulations, 106. Finocchi, F., Galli, G., Bertoni, C. M., et al. (1992) Phys. Rev. Lett., 68, 3044. Finocchi, F., Galli, G., Bertoni, C. M., et al. (1993) Physica, B185, 379. Finocchi, F. and Galli, G. (1994) Phys. Rev., B50, 7393. Fukumoto, A. (1997) Phys. Stat. Sol., 202, 125. Guraya, M. M., Ascolani, H., Zampieri, G., et al. (1990) Phys. Rev., B42, 5677. Jones, R. and Gunnarsson, O. (1989) Rev. Mod. Phys., 61, 689. Kalia, R. K., Nakano, A., Tsuruta, K., et al. (1997) Phys. Rev. Lett., 78, 689, 2144. Kelires, P. C. (1991) Europhys. Lett., 14, 43. Kroll, P., Riedel, R. and Hoffman, R. (1999) Phys. Rev., B60, 3126. Liu, A. Y. and Cohen, M. L. (1989) Science, 245, 841. Liu, A. Y. and Wentzcovitch, R. M. (1994) Phys. Rev., B50, 10362. Lannoo, M. (1992) in Handbook on Semiconductors, vol. I. (ed. T. S. Moss), Elsevier, Amsterdam. Lowther, J. E. (1998) Phys. Rev., B57, 5724. Lowther, J. E. (1999) Phys. Rev., B60, 11943. Matsunaga, K., Iwamoto, Y., Fisher, C. A. J., et al. (1999) J. Ceram. Soc. Jpn, 107, 1025.

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Page 298 Meyer, M. and Pontikis, V. (eds), Proceedings of the NATO Advanced Study Institute on Computer Simulations in Material Science, Aussois, France (1991). Series E: Applied Science – vol. 205. Kluwer Academic Publishers (Dordrecht/Boston/London). Mousseau, N. and Lewis, L. J. (1997) Phys. Rev., B56, 9641. Mura, D., Colombo, L., Bertoncini, R., et al. (1998) Phys. Rev., B58, 10357. Oguch, T. and Sasaki, T. (1991) Prog. Theor. Phys., 103, 93. O’Hara, P. A. G., Tomasziewicz, I. and Seifert, H. J. (1997) J. Mater. Res., 12, 3203. Riedel, R. A., Greiner, G., Miehe, et al. (1997) Angew. Chem., 36, 603. Robertson, J. (1991) J. Non-Cryst. Solids, 137/138, 825. Rosato, V., Celino, M. and Colombo, L. (1998) Comp. Mater. Sci., 10, 67. Seher, M., Bill, J., Aldinger, F., et al. (1994) J. Cryst. Growth, 137, 452. Sénémaud, C., Driss-Khodja, M., Gheorghiu, A., et al. (1993) J. Appl. Phys., 74, 5042. Servalli, G. and Colombo, L. (1993) Europhys. Lett., 22, 107. Ténégal, F., Flank, A. M. and Herlin, N. (1996) Phys. Rev., B54, 12029. Ténégal, F., Bouchet, B., Bellissent, R., et al. (1998) Philos. Mag. A, 78, 803. Tersoff, J. (1989) Phys. Rev., B39, 5566. Tersoff, J. (1994) Phys. Rev., B49, 16349. Thangen, T., Lippold, G., Riede, V., et al. (1999) Thin Solid Films, 348, 103. Uhlig, H., Frieß, M., Durr, J., et al. (1996) Z. Naturforsh., 51a, 1179. Umesaki, N., Hirosaki, N. and Hirao, K. (1992) J. Non-Cryst. Solids, 150, 120. Vashista, P., Kalia, R. K. and Ebbsjo, I. (1995) Phys. Rev. Lett., 75, 858. Vollmayr, K., Kob, W. and Binder, K. (1996) Phys. Rev., B54, 15808. Wang, C. -Z., Ho, K. M. and Chan, C. T. (1993) Phys. Rev. Lett., 70, 611. Wang, C. -Z., Wang, E. -G. and Dai, Q. (1998) J. Appl. Phys., 83, 1975. Waseda, Y. (1980) The Structure of Non-crystalline Materials. Mc-Graw Hill, New York. Wooten, F. and Weaire, D. (1987) Solid State Phys., 40,1.

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Page 299 Index ab initio molecular dynamic simulations (amorphous silicon carbide) 287 acicular grain growth 258 activation energy 235, 237 additives 238 aerosol 8 agglomerates 246 27Al MAS–NMR 117 27Al MAS–NMR of Si/C/N/Al/Y 132, 133, 135, 136 27Al MAS–NMR of reference samples 134 27Al NMR gibbsite 117 27Al NMR saponite 117 27Al NMR spin flip angle 115 Al and Y additives 22 Al2O3242 aluminium 27Al NMR spectroscopy 115 aluminium isopropoxyde 264 amorphous ceramics by thermolysis of polymer precursors 2 analytical method informations 4 annealed SiC 58, 59, 60 annealing treatment 192 Arrhenius law 235 atomic chemical composition CxSiyNz 127 atomic NMR and chemical compositions CxSiyNz 130 13C MAS–NMR of Si/C/N/Al/Y 130, 132 carboreduction 50 carbothermal nitridation 51 charge carriers 230 chemical and structural changes 38 chemical composition 12 chemical order 285 chemical shift anisotropy (CSA) 112 C–N bond 167, 168, 192, 217 C/N ratio 38, 177, 192 composite 264 composites microstructure 247 conduction band (CB) 154 conductivity 227, 230, 231, 232, 235, 236 core level binding energy (BE) 154 core levels 155 creep resistance 269 cross measurement interest 4, 5 cross-polarisation (CP) 112 Curie law 141, 146, 150 dangling bonds (DBs) 147 decomposition 26, 42 defect concentration 139, 141 demixion 220 densification kinetics 247 dielectric function 230, 232 dielectric loss 227, 228, 230, 232–4 dielectric properties 227 dispersing agent 242 dissolution–diffusion–reprecipitation 239 ductility 269 file:///C|/Users/steppenwolf/Documents/educatio...ks/to%20burn/nanocomposites/files/page_299.html (1 of 2) [20/01/2010 10:06:07 AM]

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EELS electron energy loss spectroscopy 79, 81 effect of NH394 electrical conductivity 234 electrical properties 227 electron paramagnetic resonance (EPR) 139–41, 143–4, 146, 148 electronic transport 230 EPR spectrometer 228 evolution of structure 33 evolved species 42 exhausted molecules 45 extended X-ray absorption fine structure (EXAFS) 163, 182, 187

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Page 300 extended X-ray absorption fine structure (EXAFS) modelisation 198 first sharp diffraction peak (FSDP) 215 flame 9 flame temperature 9 fluorescence yield (FY) 153 forming process 243 fracture strength 265 free carbon 128, 135 free carbon effect 131, 134 gaseous precursor 8, 54, 264 grain growth 240 g-tensor 151 Hall effect 228, 230 hardening rate 271 haxamethyldisilazane 8, 264 heterogeneous nucleation 253 high power decoupling (HPD) 112 high temperature properties (silicon nitride based ceramics) 268 HMDS materials 274 hot-pressed 244, 251, 253 HSAl materials 274 hyperfine diagonal tensor 140 hyperfine interactions 150 hyperfine lines 141 147 hyperfine tensor 145 hyperfine triplet 151 inclusion of H promotes chemical order (amorphous silicon carbide) 287 influence of hot pressing on 29Si and 27Al MAS–NMR spectra 137 infrared laser pyrolysis 6, 19 in-situ carboreduction reaction 262 interfacial barriers 235 interfacial polarisation 232, 233, 237 laser synthesis 6 laser-driven reaction 9 liquid precursor 8, 72 low size dispersion 23 magnetic field 144 mass change 27 mechanical properties 264 micro–nano composites 2, 241 microstructure 260 microwave 141 microwave method 228 microwave power 229 milling 253 mixed tetrahedra 166, 192 mixing process 242 mobility 230 model interatomic potential energy 284 modelisation of nanocrystalline structures 220 molecular mechanic 200 file:///C|/Users/steppenwolf/Documents/educatio...ks/to%20burn/nanocomposites/files/page_300.html (1 of 2) [20/01/2010 10:06:08 AM]

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morphological change 39 multiple scattering 185, 200 23Na MAS–NMR 116 nano–nano composites 2, 241 nanoparticles 264 nanopowders 6 nano-sized powders 227 neutron diffraction 211 nitridation 27 nitriding process 151, 152 numerical simulation limitations (gap, size, alloys) 282 octamer silicic trimethyl ester (Q8M8) 112, 114 order parameter 285, 286 organometallic precursors 20 oxynitride phase 46 oxynitrided eutectic phase 238 paramagnetic centres 139, 140, 143, 145–7, 150 paramagnetic defects 144, 145, 146, 151, 152 paramagnetic defects in Si/C/N nanopowders 139, 140 paramagnetic species 148, 151 particle rearrangement 239 Pauli 147 permittivity 230 α → β phase transformation 239, 246 α → β phase transformation rate 251 phase transformations 250 α polytypes silicon carbides 118, 120 post-sintering heat treatment 279 powder-making processes 3 prealloyed powders 256, 261 preceramic nanopowders 41 quality factor 228 quench rate 283 rate of densification 256 reactivity 41

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Page 301 reactivity of nanopowders 262 relaxation 228, 231–3, 235, 236 residence time 9 residual functionalities 53 resonance 139, 144, 147 resonant cavity 228 room temperature properties of ceramics 264 second phase 253 sedimentation tests 242 SEM observations 248 Shrinkage curves 245, 254 29Si chemical shift for crystallised SiCxN4−x 121 29Si chemical shift for silicon carbides 118 29Si chemical shift for silicon nitrides 120 29Si MAS–NMR duty cycle influence on Si/C/N/Al/Y spectrum 129 29Si MAS–NMR of Si/C/N/Al/Y 132, 136 29Si MAS–NMR silicon carbide spectra 117 Si/C powders 156 α-SiCN 220, 222 β-SiCN 220, 224 Si/C/N nanopowders 41 Si/C/N powders 154, 164, 179 Si/C/N/Al/Y nanopowders 172, 205 Si/C/N/Al/Y/O 31, 38, 39 Si/C/N/Al-based powders 83, 94 Si/C-based precursors 54 Si/N powder 164 α-Si3N4157, 186, 219 α → β Si3N4 transformation 259 β-Si3N4186 Si3N495 Si3N4/Si/C/N composites 100, 101, 104 Si3N4/SiC composites 96 β-SiAlON 134 SiAlON 105 Si-base powders 8 β-SiC 186 β-SiC silicon carbide 118, 120 SiC 6, 246 SiC nanopowders 139 SiC phase 57 SiC precipitates 260 SiC/Si3N4 phase differentiation 94 SiC-based powders 64, 72 SiCN 6 SiCxN4−x 158 SiCxN4−x tetrahedra 196, 219 SiCxNyOz phase 46 SiH48 Si–K edge 187 Si–K edge XANES 196 Silicon carbide 241 silicon nitride (α and β) 124 silicon nitride 238 simulation of amorphous Si3CzN4−z 295 simulation of amorphous silicon carbide 286 file:///C|/Users/steppenwolf/Documents/educatio...ks/to%20burn/nanocomposites/files/page_301.html (1 of 2) [20/01/2010 10:06:08 AM]

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simulation of amorphous silicon nitride 288 simulation of crystalline Si/C/N 291 simulation of ternary Si/C/N 290, 291, 295 SiN3C tetrahedra 200 sintered composites 95 sintering 238 sintering aids 238 slip casting 242 soft X-ray spectroscopy (SXS) 153–5, 160, 163 specific surface area 18 spin hamiltonian 140, 144 spin susceptibility 143, 145 spin-orbit coupling 140 strength 264, 269 superplastic behaviour 269 susceptibility 141, 150 TEM operating conditions 108 TEM operating procedure 106 TEM sample preparation (powders) 105 TEM sample preparation (sintered materials) 106 tensile ductility 269 TG/MS analysis 42 TGA/MS study 42 theoretical structure factors 220 thermal behaviour 40 thermal conversion 44 thermal stability 41 thermal treatment 25–7 total electron yield (TEY) 153 toughness 264, 266 transamination/condensation reactions 43, 44 transport properties 227 unpaired electron 141, 144, 145, 151 unpaired spin 139, 143, 150 vacancies 141, 145 valence band 159 valence bond (VB) 153 X-band 141, 142, 144, 146, 148, 229 X-EDSX-photon energy dispersive spectroscopy 56, 88, 109

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Page 302 X-ray absorption near edge structure (XANES) 153, 154, 185 X-ray absorption spectroscopy 161, 162, 176, 179 X-ray diffraction 22, 211 X-ray emission spectroscopy (XES) 153, 154 X-ray photoemission spectroscopy (XPS) 153, 154 Y2O3242

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Nanostructured silicon-based powders and composites

Nanostructured Silicon-based Powders and Composites

Nanostructured Silicon-based Powders and Composites

Nanostructured Superconductors

Nanostructured materials and nanotechnology

Nanostructured Materials and Nanotechnology

Composites

Nano-Architectured and Nanostructured Materials

Characterization of Powders and Aerosols

Analytical Methods for Food and Dairy Powders

Green Composites: Polymer Composites and the Environment

Nanostructured Materials

Nanostructured Biomaterials

Nanostructured Superconductors

Nanostructured Materials

Nanostructured Materials

Nanostructured Oxides

Copolymers, Polyblends, and Composites

Cascade Separation of Powders

Advanced Catalysts and Nanostructured Materials

Nanostructured Thin Films and Surfaces

Fibers and Composites

Adsorption by Powders and Porous Solids

Laminar Composites

Polymer Composites

Nanostructured Conductive Polymers

Nanostructured Materials in Electrochemistry

Bulk Nanostructured Materials

Chemical Physics of Nanostructured Semiconductors

Nanostructured silicon-based powders and composites