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The Springer Series in Materials Science covers the complete spectrum of materials physics, including fundamental principles, physical properties, materials theory and design. Recognizing the increasing importance of materials science in future device technologies, the book titles in this series reflect the state-of-the-art in understanding and controlling the structure and properties of all important classes of materials.

51 Microscopic and Electronic Structure of Point Defects in Semiconductors and Insulators Determination and Interpretation of Paramagnetic Hyperfine Interaction Editors: J. M. Spaeth and H. Overhof

52 Polymer Films with Embedded Metal Nanoparticles By A. Heilmann

53 Nanocrystalline Ceran1ics Synthesis and Structure By M. Winterer

54 Electronic Structure and Magnetism of Complex Materials Editors: D.J. Singh and A. Dimitrios

55 Quasicrystals An Introduction to Structure, Physical Properties and Applications Editors: J.-B. Suck, M. Schreiber, P. Haussler

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57 Radiation Effects in Advanced Semiconductor Materials and Devices By C. Claeys and E. Simoen

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Markus Winterer

Nanocrystalline Ceramics

Synthesis and Structure

With 1 71 Figures

i Springer

Dr. Markus Winterer TU Darmstadt, Institute of Materials Science, Petersenstr. 23, 64287 Darmstadt, Germany

Series Editors:

Prof. R.M. Osgood, Jr. Microelectronics Science Laboratory Department of Electrical Engineering Columbia University Seeley W. Mudd Building NewYork, NY 10027, USA

Prof. Dr. Jiirgen Parisi Universitiit Oldenburg Fachbereich Physik Abt. Energie- und Halbleiterforschung Carl-von-Ossietzky-Str. 9-11 26129 Oldenburg, Germany

ISSN 0933-033x

ISBN 978-3-642-07784-5 Cataloging-in-Publication data applied for Die Deutsche Bibliothek - CIP-Einheitsaufnahme Winterer, Markus:

Prof. Robert Hull University of Virginia Dept. of Materials Science and Engineering Thornton Hall Charlottsville, VA 22903-2442, USA

Nanocrystalline ceramics : synthesis and structure I Markus Winterer.-

(Springer series in materials science; 53) (Physics and astronomy online library)

ISBN 978-3-642-07784-5 ISBN 978-3-662-04976-1 (eBook) DOI 10.1007/978-3-662-04976-1

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

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©Springer-Verlag Berlin Heidelberg 2002 Originally published by Springer-Verlag Berlin Heidelberg New York in 2002 Softcover reprint of the hardcover I st edition 2002 The use of general descriptive names, registered names, trademarks, etc. in this publication does not imply, even in the absence of a specific statement, that such names are exempt from the relevant protective laws and regulations and therefore free for general use.

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Dedicated to my brothers Christoph and Andreas and my parents Meta and Martin

Foreword

In the last two decades the synthesis and the investigations of new materials

(called nanostructured or nanocrystalline materials) by tailoring the atomic and/or

chemical microstrucure on a nanometer scale has become one of the most rapidly

growing areas of Materials Science. The numerous studies published so far have

remarkably improved our basic undertanding of the structure as well as of the

properties of these substances. Most nanostructured materials are far away from

thermodynamic equilibrium. Hence, the results of studies of their structure are of

limited general significance unless the structure, chemical composition, defect ar-

rangement etc. of the specific nanostructured material investigated is well charac-

terized. In the past numerous controversial results and discussions may have been

avoided by paying more attention to the non-equilibrium nature of nanostructured

materials. Dr. Winterer's work is a pioneering and remarkable example of a thor-

ough study of this kind.

Prof. Dr. rer. nat. Herbert Gleiter

Director of the Institute of Nanotechnology, Karlsruhe Research Center,

Karlsruhe, Germany, October 200 1

This thesis is a comprehensive document and certainly one of the best I have

read on nanostructured materials. Overall, I got the clear impression that every as-pect of the work described represents frontier work. For example, the consistency

between the EXAFS/RMC results and the structural evolution during sintering

obtained by more conventional analytical methods was quite impressive.

Regarding the topic that is closest to my own scientific interest, I was intrigued

to see how much progress has been made in the chemical synthesis (CVS) of

nanoceramic materials. This is an exceptional piece of work- perhaps the most

authoritative in the field. I was also impressed by the work on Al20rdoped Zr02 compositions. Until reading this work, I had no idea that the CVS process was

such a non-equilibrium process, enabling remarkable supersaturations of Al20 3 in

the Zr02 host to be realized. What made this work particularly fascinating to me

VIII Foreword

was the one-to-one correspondence between Winterer's findings and ours, using a rapid solidification method. What is clear now is that these two quite different

processing methods, i.e. rapid condensation from the vapor state and rapid solidi-

fication from the melt, are complementary. I hope that the author continues his re-

search along this line, since I believe that he is opening the door to a major field of

study in nanostructured materials.

The other area that is well documented in the treatise is the general area of gas

phase synthesis of nanoceramic powders. The author demonstrates a complete un-derstanding of the mechanisms and kinetics involved in controlled thermal de-

composition of metal organic precursors. A particularly nice touch is the control of

decomposition of two precursors to produce: (1) homogeneous nanoparticles, (2)

coated nanoparticles, and (3) phase-separated nanoparticles. The characterization

of these materials by High-Resolution TEM is particularly well done. With regard

to sintering of these 'designer' nanoparticles, it is interesting to note the effective

control of grain growth during sintering using coated nanoparticles. The issues

related to segregation are also dealt with very nicely.

The evolution of the structures from clusters to bulk materials in both SiC- and

Zr02-base systems is well described and to the point. The discussion builds on previous work in the field and the author has missed nothing of importance in the

literature.

If there is any shortcoming in this comprehensive treatise, it is the absence of discussion of scale-up issues. Clearly, there are problems associated with the hot wall reactor process when looked at from a practical viewpoint, which the author

does not discuss. However, this is essentially an academic work, so the author

certainly can be excused for not dealing with practical implications. In summary, this is an excellent treatise on the synthesis, processing, charac-

terization, and modeling of nanostructured ceramic materials. I hope that it gets the attention that it deserves.

Prof. Dr. Bernard H. Kear

Department of Ceramic and Materials Engineering, Rutgers University,

Piscataway, NJ, USA,

October 2001

Preface and Acknowledgements

This book was originally prepared as my habilitation thesis. It is a synopsis of major results of eight years of scientific research at the Institute of Materials Sci-ence at the Darmstadt University of Technology. It is my hope that readers of this book get interested in nanocrystalline ceramics and inspired for their own re-search.

The research presented in this book and ongoing work would not have been possible without the hard work of my students and collaborators. Therefore, it is a great pleasure to thank all who have contributed in various ways: • Prof. Horst Hahn for his continuous support and encouragement to work on

nanoceramics and for the opportunity and freedom to develop my own re-search;

• Prof. Vladimir Srdic' for his synthesis and sintering studies of zirconia during his time as a Humboldt fellow in our group, Dr. Robert Nitsche who was my first PhD student and started with IGC synthesis, sintering and EXAFS spec-troscopy of zirconia, Dr. Sylke Klein for her systematic work on the synthesis of silicon carbide (she started with CVS), Dipl. Ing. Stefan Kobel who was my first diploma student in Darmstadt, Dipl. lng. Stefan Seifried who did pioneer-ing work on granular films and set up the LPDS, Dipl. Ing. In-Kyum Lee who set up the AMS, Dipl. lng. Andreas Benker for his work on yttria stabilized zir-conia, Dipl. lng. Michael Schallehn for the plasma-coating experiments, Frank Sauberlich for his work on strontium titanate, Dr. Andreas Moller for his con-tribution to powder dispersability and surface potential, all other members of the 'Dunne Schichten' group for the good atmosphere and our guests who pro-vided an international flair and open-minded discussions, especially the Hum-boldt fellow Dr. Subramshu Bhattacharya (liT Madras) and Dr. Sarbari Bhat-tacharya who were kind enough to correct my English;

• Dr. Gerhard Miehe, Dr. Thomas Weirich, Dr. Veronique Buschmann, and Dr. Matthias Rodewald (Structural Analysis Division) who contributed the TEM images, Dr. Peter Hoffmann (Chemical Analysis Division) for his generous gift of furnaces and a glove box, and Dipl. Ing. Martin Heck for X-ray fluorescence measurements;

X Preface and Acknowledgements

• Dr. Tomas Diaz de la Rubia and Dr. Maria-Jose Caturla (LLNL) for introduc-ing me to MD simulations, Prof. Roth, Dirk Lindackers and Christian Jansen

(Uni Duisburg) for the generous help with the AMS, Prof. Eckert (Uni Mun-ster) for solid-state NMR measurements, Prof. Priya Vashishta, Prof. Kalia and

Dr. Alok Chatterjee (University of Lousiana) for MD simulations on nanocrys-

talline silicon carbide, and Prof. Sotiris Pratsinis (ETH Zurich) for his kind en-

couragement and advice to improve the CVS model.

External support was essential for the measurement of the EXAFS spectra and the collection of the neutron scattering data and is gratefully acknowledged: Dr. Chun Loong (ANL), Dr. Robert McGreevy and Dr. Bob Delaplane (NFL Stud-

svik), Dr. Udo Keiderling (BENSC/HMI), Dr. Ronald Frahm, Dr. Larc Troger,

and Dipl. Phys. Klaus Attenkofer (HASYLAB/DESY), Dr. Menno Oversluizen,

Dr. Andy Dent and Dr. Fred Mosselmans (DRL), Dr. Frentrup and his students (Humboldt Universtitat Berlin, BESSY), and all other personel at those facilities

for their help with the experiments.

Especially, I want to thank Jiirgen Schreeck our technician and Jochen Korzer

and his colleagues in the machine shop who were of indispensable help in building

and constructing a lot of equipment and fast help in case of improvisations. Additionally, I gratefully acknowledge the generous support by the following

funding agencies: the German Science Foundation (DFG), the Alexander von

Humboldt Foundation (AvH), the Federal Ministry of Education and Research (BMBF) and the Large Scale Facility activity of the Training and Mobility of Re-searchers programme of the European Commission.

Last but not least, I want to thank Prof. Gleiter and Prof. Kear for their very kind forewords.

Markus Winterer, Darmstadt

May2002

Abbreviations, Acronyms and Symbols

a a A at Aiik, r) AK AMS

ASB

a. Aw b b B

b. b(2(J) BET Bi BJH

b. c Cz

Ci

ci

cp cP

Cv eve CVD CVP CVR cvs d

mean molecular velocity (p. 39) root mean square displacement (p. 160) scattering angle (p. 156) lattice constant (p. 27) surface area of a single aerosol particle (p. 48) total surface area of an aerosol (p. 47) surface area of a monomer (p. 48) EXAFS amplitude function for a single atom pair (p. 165) absorption factor (p. 159) aerosol mass spectrometer (p. 229)

alurninium-s-butoxide (p. 17) surface area of a completely coalesced (sinterered) particle (p. 48) surface area of reactor wall (p. 43) lattice constant XRD constant background parameter, sample (p. 161) XRD full width at half maximum, sample (p. 98)

XRD full width at half maximum, standard (p. 21, 98) XRD background intensity (p. 159) Brunauer-Emmett-Teller nitrogen adsorption isotherm (p. 24) temperature factor in Rietveld analysis (p. 160) Barret-Joyner-Halenda analysis (p. 25)

XRD constant background parameter, standard (p. 22, 161) lattice constant second order cumulant (p. 168) atomic fraction (p. 157) concentration (p. 38) molar heat capacity at constant pressure (p. 40) specific heat capacity at constant pressure (p. 40) molar heat capacity at constant volume (p. 40) Chemical Vapor Condensation (p. 9) Chemical Vapor Deposition (p. 9) Chemical Vapor Precipitation (p. 9) Chemical Vapor Reaction (p. 9) Chemical Vapor Synthesis (p. 9)

(crystallite) diameter (p. 161)

XII Abbreviations, Acronyms and Symbols

D D

D d* do dBET

dg Dgb

Dt Dm Ds Ds, Dsur dXRD

E Eo Ea Ev EDX EPR EXAFS F(q), FK feff7 f,Jm Fi (k) FULLPROF G g GDE gu(r), g(r)

g(0-0)

g(Zr-0)

g(Zr-Zr)

goo(r)

gz,o(r)

gz,zlr) h, k, l HRSEM HSP IGC k k K

ko

diffusion coefficient (p. 39) fractal dimension (p. 26)

particle diffusion coefficient (p. 49) critical cluster diameter (p. 37) initial grain size (p. 95) particle size determined from the BET specific surface area (p. 24) geometric mean diameter (p. 20) grain boundary diffusion coefficient (p. 50) lattice diffusion coefficient (p. 95) mass fractal dimension (p. 26, 29, 50) surface fractal dimension (p. 26, 29) surface diffusion coefficient (p. 81, 95) grain size determined from XRD line broadening (p. 21) X-ray poton energy (p. 162) threshold energy (p. 162) activation enthalpy (p. 41) activation energy for diffusion (p. 82) Energy Dispersive X-ray analysis (p. 100) Electron Paramagnetic Resonance (p. 149) Extended X-ray Absorption Fine Structure (p. 161) structure factor (p. 157, 160) software for the ab initio computation of EXAFS spectra (p. 166) atomic scattering factors (p. 157) magnitude of the backscattering amplitude (p. 162) software for Rietveld refinements (p. 161) Gaussian profile function (p. 160) transition parameter (p. 49) General Dynamic Equation (p. 36)

partial pair distribution function (p. 157)

0-0 partial pair distribution function (p. 186)

Zr-0 partial pair distribution function (p. 185)

Zr-Zr partial pair distribution function (p. 186)

0-0 partial pair distribution function (p. 168)

Zr-0 partial pair distribution function (p. 168)

Zr-Zr partial pair distribution function (pp. 178) Miller indices (p. 160) High Resolution Scanning Electron Microscopy (p. 100) hard sphere potential (p. 168) Inert Gas Condensation (p. 36) photo electron wave vector (p. 162) reaction constant (p. 41) shape factor in Scherror eqation (p. 98) preexponential constant (p. 41)

k8 ,k Kn L l L

LK m M,M; MCGR

MJ N N n N N,

Nu Nmc NMR

NP Nu NVT p, P2 P3 Pa PCS PDF Pe,Peh, Pem PK Pr q QEXAFS

r R R R Re REFLEXAFS

RJ rm RMC RMCA rmcxas s

Abbreviations, Acronyms and Symbols XIII

Boltzmann constant Knudsen number (p. 42) Lorentzian profile function (p. 160) particle mean free path (p. 49) typical dimension (p. 42) Lorentz factor (p. 159) mass molar mass software for pair distribution analysis of structure factors (p. 157)

temperature factor (p. 160) degree of agglomeration (p. 24) coordination number (p. 172) grain size exponent (p. 93) number number density of monomers (p. 85) coordination numbers (p. 167) number density of monomers in clusters (p. 85) Nuclear Magnetic Resonance (p. 153) number density of precursor molecules (p. 47) Nusselt number (p. 44) ensemble at constant number, volume and temperature (p. 167) partial pressure of momomers (p. 82) second moment (p. 168) third moment (p. 197) acceptance probability (p. 166) photon correlation spectroscopy (p. 100) pair distribution function (p. 168) Peclet number (p. 42) preferred orientation function (p. 165) Prandtl number (p. 42) scattering vector (p. 27, 156) Quick EXAFS (p. 169)

radius, coordination distance (p. 172) gas constant radius of gyration (p. 27) reliability factor (p. 167) Reynolds number (p. 42) Reflection EXAFS (p. 169) coordination distance (p. 162) atom position (p. 156) Reverse Monte Carlo (p. 158) software to analyse structure factors by RMC modeling (p. 157) software to analyse EXAFS spectra (p. 166) scale factor (p. 159)

XIV Abbreviations, Acronyms and Symbols

s So So2 SANS Sc SEM S; Sij(q)

T

tic TEM TMS Tw u U, V,and W

v

VTMS• VHe

w xafs XAFS XANES X;

xj, yj, Zj

XRD Z; Zu,Zij ZTB

Zw ~

AEo ~FG ~G*

M ~RG ~RH

~/3

specific surface area (p. 24) specific surface area of spherical particles (p. 69) amplitude reduction factor (p. 162) Small Angle Neutron Scattering (p. 27) Schmidt number (p. 43) Scanning Electron Microscopy (p. 100) scale factor of phase i (p. 161) partial structure factor (p. 157)

temperature time tetragonal or cubic structure (p. 98) Transmission Electron Microscopy (p. 18) tetramethylsilane (p. 15) reactor wall temperature (p. 43) flow velocity (p. 38) Cagliotti parameters (p. 161)

volume (p. 39) volume of agglomerate particle (p. 48) volume of primary particle (p. 67) volume of monomer (p. 48) virtual crystal approximation (p. 149) unit cell volume of phase i (p. 161) volume fraction of phase i (p. 161) molar volume (p. 37) mass flow of He, or TMS (p. 57) grain boundary width (p. 50) software to analyse XAFS spectra (p. 170) X-ray Absorption Fine Structure (p. 149) X-ray Absorption Near Edge Structure (p. 164) mole fractions (p. 40) fractional coordinates (p. 160) X-ray diffraction number of atoms in the unit cell (p. 161) binary collision frequency of particles (p. 39)

zirconium-t-butoxide (p. 16) collision frequency of particles with a wall (p. 39) core radius (p. 162) inner potential (p. 162) free enthalpy of formation (p. 37) free enthalpy of cluster of critical size (p. 37) histogram bin width (p. 167) free enthalpy of reaction (p. 37) enthalpy of reaction (p. 38) change in fit residual (p. 166)

l/Ju(k, r)

r r 1: Q a f3 f3 f3 f3 X(k) Or max l/> l/>j (k)

r n Yii y;i(r, k) y; 11 11 n K ). ).

).

J.l(E) v Y;

p p Po o(i)

a, aii ag (12

J

(1Zr-O

(1Zr-Zr

C1oo

1'

1'112

'l'c

'l's

~

Abbreviations, Acronyms and Symbols XV

EXAFS phase function of atom pair ij (p. 162)

crystallinity (p. 22,81, 161)

full width at half maximum (p. 160) volume to surface ratio (p. 85) volume of diffusing species (p. 49) heat exchange coefficient (p. 43) corrected full width at half maximum ofXRD lines (p. 21) monodisperse coagulation frequency function (p. 47) normalized fit residual in RMC analysis (p. 166)

XRD full width at half maximum, corrected (p. 98) EXAFS signal (p. 162) maximum Monte Carlo step (p. 168) modified Thompson-Cox-Hastings pseudo-Voigt profile (p. 160) EXAFS phase shift (p. 162) surface (or interfacial) enthalpy (p. 37, 50) interfacial enthalpy of grain boundary (p. 95) weight coefficients for partial structure factors (p. 157) EXAFS signal corresponding to a single atom pair (p. 165) surface enthalpy (p. 95) dynamic viscosity (p. 40) profile function mixing factor (p. 160) Planck's constant(= hI 21t) thermal conductivity (p. 40) mean free path (p. 39) mean free path of photo electrons (p. 162) X-ray wavelength (p. 161) X-ray absorption coefficent (p. 162) kinematic viscosity (p. 40) stoichiometric coeffcients (p. 41) (bulk) density (p. 24) number density (p. 165) initial density (p. 95) error in RMC analysis (p. 166) collision diameter (p. 39) geometric standard deviation (p. 20) mean square vibrational amplitude (Debye-Waller factor) (p. 162) cut off radii for hard sphere potential for atom pair Zr-0 (p. 168) cut off radii for hard sphere potential for atom pair Zr-Zr (p. 168) cut off radii for hard sphere potential for atom pair 0-0 (p. 168) residence time (p. 42) half life (p. 41) characteristic time for coagulation (p. 60) characteristic time for sintering (p. 60) CVS number (p. 85)

Contents

1 Introduction .............................................................................................................. !

2 Gas Phase Synthesis ................................................................................................. 7 2.1 Background ....................................................................................................... 8

2.1.1 Chemical Vapor Synthesis ........................................................................ 8

2.1.2 Silicon Carbide ........................................................................................ 11

2.2 Experimental Methodology ............................................................................ 12

2.3 Experimental Results for Silicon Carbide ...................................................... l8

2.3.1 Particle Size and Morphology ................................................................ 18

2.3.2 Crystalline Phase, Grain Size and Crystallinity ..................................... 21

2.3.3 Surface Area, Agglomerate Size and Morphology ................................ 24

2.3.4 Influence of Decomposition Temperature ............................................. 30

2.3.5 Influence of Total Pressure ..................................................................... 30

2.3.6 Influence of Precursor Partial Pressure .................................................. 32

2.3 .7 Influence of Reactor Length ................................................................... 33

2.4 Summary ......................................................................................................... 33

3 Modeling Particle Formation and Growth ......................................................... 35 3.1 Background ..................................................................................................... 36

3.2 Basic Physical Chemistry and Chemical Engineering Relations .................. 38

3.2.1 Gas Kinetics ............................................................................................ 39

3.2.2 Chemical Reaction Kinetics ................................................................... 40

3.2.3 Hydrodynamics ....................................................................................... 42

3.2.4 Heat Transfer and Production ................................................................ .43

3.3 Reaction-Coagulation-Sintering Model ......................................................... 45

3.3.1 Particle Formation ................................................................................... 47

3.3.2 Coagulation and Sintering ..................................................................... .47

3.3.3 Heat Balance ........................................................................................... 50

3.3.4 Numerical Implementation ..................................................................... 51

3.4 Results of Numerical Simulations .................................................................. 54

3.4.1 Time-Temperature-Profile ...................................................................... 54

XVIII Contents

3 .4.2 Evolution of Particle Sizes and Specific Surface Area ......................... 56 3.4.3 Evolution of Number Densities .............................................................. 59 3 .4.4 Characteristic Times ............................................................................... 59 3.4.5 Gas Kinetic Parameters ........................................................................... 62 3.4.6 Chemical Engineering Parameters ......................................................... 64

3.5 Comparison of Numerical Simulations and Experimental Results .............. 65 3.5 .l Dependence on Reactor Temperature .................................................... 66 3.5 .2 Dependence on Reactor Length ............................................................. 70 3.5 .3 Dependence on Reactor Diameter .......................................................... 70 3.5 .4 Dependence on Precursor Partial Pressure ............................................ 72 3.5 .5 Dependence on Process Pressure ........................................................... 73 3.5 .6 Dependence on the Shape of the Temperature Profile .......................... 73 3.5.7 Predicting Power ..................................................................................... 77 3.5 .8 Considerations for Scale up .................................................................... 78 3 .5 .9 Crystallinity ............................................................................................. 81 3.5.10 From Films to Particles ......................................................................... 84

3.6 Summary ......................................................................................................... 87

4 Processing and Microstructure ............................................................................ 91 4.1 Background ..................................................................................................... 92

4.1.1 Microstructure in Nanocrystalline Materials ......................................... 92 4.1.2 Sintering .................................................................................................. 93 4.1.3 Zirconia Based Ceramics ........................................................................ 96

4.2 Experimental Methodology ............................................................................ 98 4.2.1 Powder Characterization ......................................................................... 98 4.2.2 Compaction and Sintering ...................................................................... 99 4.2.3 Pellet Characterization ............................................................................ 99 4.2.4 Zeta-potential and Particle Size Distribution in Aqueous Dispersion 100

4.3 Pure Zirconia ................................................................................................. 100 4.3.1 Powder Characterization ....................................................................... 100 4.3.2 Compaction Behavior ........................................................................... 105 4.3.3 Sintering ................................................................................................ 106

4.4 Zirconia Doped with Alumina ..................................................................... 114 4.4.1 Compaction Behavior and Microstructure of Green Bodies ............... 115 4.4.2 Sintering Behavior and Microstructural Evolution ............................. 120 4.4.3 Alumina as Grain Growth Inhibitor ..................................................... 126

4.5 Zirconia Coated with Alumina ..................................................................... 128 4.5.1 Characteristics of Powders ................................................................... 129 4.5.2 Characteristics of Aqueous Dispersions .............................................. 132

Contents XIX

4.5.3 Compaction and Sintering Behavior .................................................... 133

4.5 .4 The Function of the Al20 3 Coating on Zr02 •••••••••••••••••••••••••••••••••••••••• 135

4.6 Zirconia Doped with Yttria .......................................................................... 140

4.6 .1 Powder Characterization ....................................................................... 140

4.6.2 Sintering Behaviour .............................................................................. 141

4.7 Summary ....................................................................................................... 145

5 Local Structure and Long Range Order ........................................................... 147 5 .1 Background ................................................................................................... 148

5.1.1 Nanocrystalline Materials- Heterogeneous Disorder .......................... l48

5 .1.2 Diffraction ............................................................................................. 148

5.1.3 Spectroscopy ......................................................................................... 149

5.1.4 Zirconia .................................................................................................. 152

5.2 Methodology of Data Analysis and Experimental Procedures ................... 156

5.2.1 Scattering, Structure Factor and Partial Pair Distribution Functions .. 156

5.2 .2 Powder Diffraction and Rietveld Analysis .......................................... 159

5.2.3 EXAFS Spectroscopy and Analysis by Reverse Monte Carlo

Modeling ......................................................................................................... 161

5.2.4 Experimental Procedures ...................................................................... 169

5 .3 Crystalline and Amorphous Zirconia ........................................................... 170

5.3 .1. Monoclinic Zirconia ............................................................................. 170

5 .3 .2 Amorphous Zirconia ............................................................................. 17 5

5 .4 Scattering and EXAFS Spectroscopy of monoclinic Zirconia .................... 181

5.5 Pure Nanocrystalline Zirconia Powder ........................................................ 187

5.5 .1 Structure as a Function of Synthesis Temperature .............................. 188

5.5 .2 Local Structure in N anocrystalline Tetragonal Zirconia Powder ........ 194

5.6 Nanocrystalline Zirconia Doped with Alumina ........................................... 198

5 .6.1 Zirconia Doped with 5 mol% Alumina ................................................ 199

5.6.2 Zirconia Doped with 30 mol% Alumina .............................................. 206

5.7 Nanocrystalline Zirconia Doped with Yttria ............................................... 216

5.7.1 X-Ray Diffraction ................................................................................. 216

5.7.2 XAFS Spectroscopy .............................................................................. 218

5.8 Summary ....................................................................................................... 225

6 Conclusions and Perspectives ............................................................................. 227

7 References ............................................................................................................. 233

8 lndex ...................................................................................................................... 251

1 Introduction

M. Winterer, Nanocrystalline Ceramics© Springer-Verlag Berlin Heidelberg 2002

2 1 Introduction

The question "How many atoms create a solid ?" depends on the type of material and property studied. In metal clusters, although a crystalline structure is already formed at about 100 atoms, the macroscopic melting point is only reached for clusters with 1000 atoms (Stace 1988). In semiconductor crystals up to 10000 at-oms are necessary to show optical excitations similar to that of the bulk (Alivisa-tos 1996). Nanocrystals are located in the regime between molecular clusters and crystals of micrometer size. They are characterized by an interior component, structurally identical to a corresponding bulk solid and a distinguished, substantial fraction of atoms on the crystallite surface. The physical and chemical properties of solid matter changes as the particle size decreases well below the micrometer regime because of quantum size effects, surface and interface effects and changes in the crystallographic structure. Long range forces and cooperative phenomena are more affected by size effects than short range interactions. The crystal lattice of partially covalent oxides tends to transform into a structure of higher symmetry with decreasing crystal size (Ayyub et al. 1995). Size dependent structural metas-tability has also been found in metals (Sugano 1991) and semiconductors (Chen et al. 1997).

The scientific and technological development of nanostructures has two major approaches: superlattices and quantum well devices (Notzel and Ploog 1993) and three-dimensional ultrafine, polycrystalline microstructures (Gleiter 1992). The latter, also called nanocrystalline materials (Gleiter 1989), are defined as poly-crystalline solids with grain sizes of a few nanometers. Grains, pores, interface thicknesses and defects are of similar dimensions. This microstructure results in chemical and physical size effects which are of increasingly large scientific and technological interest. The fraction of atoms in the interface regions is large at grain diameters of the order of l 0 nm as can be seen from the volume fraction of atoms located in grain boundaries as estimated from a simple bricklayer model (Fig. 1.1).

Nanocrystalline ceramics are suited for applications in which either the atomic or crystallographic structure, microstructure (e.g. high surface area, high volume fraction of atoms in interfaces), resulting properties (e.g. superplasticity, catalytic activity), the processing route (e.g. netshape forming) or the final product (e.g. ce-ramic joints) are of advantage compared to conventional materials. Major fields of applications for such advanced ceramics are, for example, catalysis, porous (ul-trafiltration) or dense (in sensors or fuel cells) membranes and electroceramics.

In this work it will be shown that crystallography, microstructure and properties are correlated to the local structure and distribution of dopants in the nanocrystal line materials and can be controlled on the molecular level at the time of the pow-

0.8

0 0.6

> ..._ .D

> 0.4

0.2

0 0 5

b = 0.5 nm b = 1.0 nm

·- ·--

b

------

10

alnm

V I V = 1 - a3 I (a-b)3 b c

-...... _ .. ........

15

······

I Introduction 3

20

Fig. 1.1. Ratio of volume fractions of atoms in grain boundaries and grains as a function of the grain size, a, for different grain boundary widths b

der synthesis. Size effects are usually observed when the dimension of a structural

feature (e.g. grain or pore size) is close to a characteristic length of a physical or

chemical property (e .g. Debye length, mean free path). Nano-materials with new

or modified properties can be created by exploiting these effects as will be shown

in the following few examples .

The electrical resistivity increases by about 2 to 3 orders of magnitude when the

grain size of Sn02 approaches the width of the space charge layer which is about

6 nm (Xu eta!. 1991). The sensitivity of a Sn02 gas sensor increases at the same time by about 2-3 times . While electrons are the conducting species in Sn02 , oxy-

gen ions are the charge carriers in yttrium doped Zr02 where the space charge

layer is depleted of oxygen . Another example are solid oxide fuel cells with ce-

ramic membranes which are operated at high temperatures (900-1000°C) to di-

rectly convert chemical energy (fuels: e.g., H2 or CH4) into electric energy with

high efficiency. The membrane of a solid state fuel cell consists essentially of

three layers, the porous cathode (e.g. (La,Sr)Mn03) which acts as a catalyst con-

verting molecular oxygen into oxygen ions and as an electron conductor, the dense

electrolyte (e .g. (Y)Zr02) which acts as a separating membrane and as an ion con-

ductor and the porous anode (e.g. a Ni-(Y)Zr02 cermet) which is a catalyst for the

conversion of the fuel and an electron conductor (Haart 1995). Production and op-

eration of solid oxide fuel cells face difficult materials problems. The sintering of

the complex composite layer compound material requires high temperatures.

However, the composite material is unstable at these temperatures which makes it

4 1 Introduction

difficult to produce a reliable, 'gas dense' seal. For high conversion efficiency the electrolyte should have a low (absolute) resistivity, e.g. by using thin membranes, and the transport of the gaseous species and their catalytical conversion at the electrodes has to be fast so that the conversion is not limited by the gas transport of reactants. During operation at high temperatures long term stability is required and corrosion by gases has to be limited. Nanocrystalline materials provide solu-tions for these problems. Catalytic conversion, ionic conductivity, and sinterability can be improved by the small grain size and the resistance to high temperature corrosion can be increased because of a large number of grain boundaries across the membrane.

Ceramic ferroelectrica are another family of materials of high technological importance where size and interface effects play an increasingly important role, particularly in microelectronics. For example, ferroelectricity vanishes at small grain sizes. This causes difficulties with the increasing demand for miniaturisia-tion in the production of integrated circuits and devices such as DRAM's (Dy-namic Random Access Memory; Sengupta et al. 1995). In ceramic multilayer ca-pacitors of BaTi03 (Hennings et al. 1991), the electrical resistivity is too low at high electric fields, the ferroelectric losses are too large and the dielectric proper-ties of BaTi03 depend strongly on the microstructure of the thin films (Waser 1999). These properties and the short device life time are a function of the grain size, specifically the number of grains, respectively grain boundaries between the electrodes. An optimization of the microstructure can solve such materials prob-lems. For BaTi03 , an increase in the device life time has been observed if the grains are smaller than the thickness of the space charge layer. The grain bounda-ries act as barrier for oxygen vacancy migration and prevent the electrical break-down at high voltages with increasing number of grain boundaries perpendicular to the electrodes. Therefore, u1trafine grained, granular microstructures are pre-ferred to columnar microstructures (Balatu et al. 1990).

These few examples show that nanocrystalline materials are not only interest-ing for fundamental research but also for technologically important applications. However, a prerequisite for use in applications is the development of processes to economically produce large quantities of nanocrystalline materials at high and controlled qualities (Gleiter 1992). Additionally, the properties which are a result of the small grain sizes will be accessible in the sintered solid only if the nanodi-mension is retained after consolidation.

These two criteria put very stringent demands on the synthesis and processing methods of nanocrystalline powders because densification and grain growth are coupled processes. However, the decrease of the melting point with decreasing

1 Introduction 5

size is already an indication that the sinterability of nanocrystalline materials is enhanced. It has also been recognized that surface chemistry effects play an im-

portant role in the processing and properties of nanocrystalline ceramics (Mayo et al. 1999).

The goal of this work was to further develop a synthesis method for the pro-

duction of nanocrystalline ceramic powders, to study the development of the mi-

crostructure during sintering and to investigate the influence of the synthesis pa-

rameters on the structure and properties of the nanocrystalline ceramics from the

atomic to the microstructural level. Such an unified view from powder synthesis

and ceramic processing to structural characterization and determination of proper-

ties offers a detailed understanding of nanocrystalline materials and enables the

precise control of the quality of the final products.

Nanocrystalline powders were produced by a modified Chemical Vapor Depo-sition (CVD) method, called Chemical Vapor Synthesis (CVS), where the process parameters are adjusted to produce ultrafine particles instead of films. CVS pro-

vides the control of grain size and chemistry at the interfaces on the molecular

level at the time of powder synthesis.

Materials with a modified 'nano'-microstructure provide the potential for new

or improved appplications. The structure of these complex, heterogeneously dis-

ordered materials is investigated with a variety of methods. It will be shown that a combination of EXAFS (Extended Absorption Fine Structure) spectroscopy in

connection with Reverse Monte Carlo analysis (RMC) provides information not accessible by other methods.

Both, silicon carbide (Bums 2000) and zirconia based ceramics (Russo and Partis 2000) are considered advanced materials with superior performance in de-manding applications. Silicon carbide is a good model system for the CVS process because it can be synthesized from a single-source-precursor with an accurate pre-cursor delivery method. The availability of thermodynamic and kinetic data make it possible to simulate the CVS process. Results of the gas phase synthesis and the modeling of particle formation and growth are presented in Chaps. 2 and 3.

Zirconia based ceramics on the other hand are good candidates for the investi-gation of the local structures and the study of the microstructural development be-

cause EXAFS spectra of high quality can be measured and the powders can be

handled in air. Results on processing and microstructural development as well as local structure and long range order are presented in Chaps. 4 and 5.

2 Gas Phase Synthesis

~ 3 ~ · .o ~ 8~ lfl 0

IJ' .e ---4t

M. Winterer, Nanocrystalline Ceramics© Springer-Verlag Berlin Heidelberg 2002

8 2 Gas Phase Synthesis

2.1 Background

2.1.1 Chemical Vapor Synthesis

Powders consisting of nanocrystalline particles can be produced by a large variety of methods based on solid state, liquid or gas phase processes. For an overview

see Chow and Gonsalves (1996). Siegel (1991) described processes based on

physical methods such as inert gas condensation or ball milling. Brinker and

Scherer (1990) gave an overview of the synthesis of particulate sols and gels and

compared them with vapor phase methods. Segal (1989) and Klabunde et al.

(1994) gave an overview of different chemical methods.

Historically, soot produced by incomplete combustion may be the first man

made nano-material. The application of soot in inks and pigments has accompa-

nied the history of mankind since prehistoric times (Ulrich 1984). Soot (carbon black) is still produced industrially in large quantities by a vapor phase (flame)

process with production rates of up to 2.5 t/h. Most of the 4.3·106 t/a (Vohler et al.

1983) is used for reinforcing rubber. Titania, silica and alumina are other industri-

ally important, ultrafine powders which are produced by gas phase processes and

are used for a wide spectrum of applications. About half the titania ( 4.1·106 t/a in 1995) is produced by the 'chloride process' which is a flame process using TiC14 as precursor. Titania is mostly used as white pigment (Heine et al. 1992). Silica (about 105 t/a in 1991) has been produced by flame hydrolysis of SiC14 (Degussa: Aerosil or Cabot Corporation: Cab-o-Sil and other manufacturers) for more than

50 years (Florke et al. 1993). Mazdiyasni et al. showed as early as 1965 that metalorganic precursors such as alkoxides can be used for the production of ul-trafine, pure oxide powders by pyrolysis in a hot wall reactor.

Another technologically important gas phase process is the Chemical Vapor Deposition (CVD) where a solid film is synthesized from the gas phase by a chemical reaction (Hitchman and Jensen 1993). Particle contamination (called

'snowing') is a serious problem in many CVD applications, especially in microe-

lectronics (Breiland and Ho 1993) where ultrafine particles may occur as an unde-

sired, film deteriorating byproduct in CVD reactors. Usually, particles are ob-

served in CVD processes under the following conditions (e.g. Bryant 1977): o at high temperatures (in hot wall reactors)

o at high supersaturations (high partial pressure of monomers at a low vapor pressure of the bulk solid)

o at long residence times (low gas flows or long reactors) and

o for small substrates.

2.1 Background 9

This can also be seen in the characteristic CVD curves (e.g. Kodas and Hampden-Smith 1994) where the film growth rate at high temperatures and precursor partial

pressures decreases above a critical value because particle formation in the gas phase results in precursor depletion.

The cross over points between CVD film growth and CVS particle formation as

a function of process temperature and precursor partial pressure are indicated by

the dashed lines in Fig. 2.1. Therefore, a modified CVD process where the process

parameters are adjusted to generate particles instead of films can be used to pro-duce nanocrystalline powders. We call this process Chemical Vapor Synthesis ( CVS). However, other names are also used in the literature such as Chemical Va-por Reaction (CVR, H.C. Starck), Chemical Vapor Precipitation (CVP, Kruis et

al. 1993) or Chemical Vapor Condensation (CVC, Chang et al. 1994). The princi-

ple advantages of the reactions in the gas phase are very short process times and

nanoscaled powders of high purity with a narrow particle size distribution.

Small (colloidal) particles dispersed in a gas are called aerosols (Reist 1993). Aerosols are produced either by conversion of gases (vapors) to particles or by

disintegration of existing liquid or solid particles. Friedlander (1977) already

pointed out that particles produced from the gas phase are usually smaller than

those produced by the disintegration processes. The production of an aerosol by a

chemical reaction (CVS) is a special case of the gas to particle conversion. These roots of the CVS process, CVD and aerosol science supply a wealth of

expertise and methods concerning reactor design, precursor selection, particle size

detection or models for particle formation. A simple way to characterize particles

,...lo.:;..g(:._w:._) ---....,....---------.., w

diffusion I feed limited homogeneous nucleation

Pi

Fig. 2.1. Characteristic CVD curves: film growth rate w as a function of inverse process temperature and precursor partial pressure (adapted from Kodas and Hampden-Smith 1994). The dashed lines mark the cross over from film growth to particle formation

10 2 Gas Phase Synthesis

is the use of a single parameter such as the diameter or particle size (Baron and Willeke 1993). The morphology of the particles can be described by fractal theory

(Schaefer and Hurd 1990) where it is distinguished between primary particles

which are the smallest discrete objects in the aerosol and secondary particles

which are usually called agglomerates or aggregates depending on the type of

force which holds the primary particles in the secondary particle together. Ag-

glomerates consist of primary particles weakly bonded by van der Waals forces. In

aggregates the primary particles are bonded more strongly e.g. by chemical bonds

or sintering necks according to definitions in aerosol science. Usually, in materials

science and also in this work the terms primary particle and soft or hard agglom-

erate are used to distinguish between secondary particles of different strength.

Aerosol science concentrates on atmospheric aerosols because of their envi-

ronmental (air pollution, global climate; Seinfeld and Pandis 1998) as well as

safety and health implications (Hinds 1982). Also industrial flame reactors are

usually operated at or close to normal pressure (Ulrich 1971, Ulrich 1984, Ulrich

and Riehl 1982). However, work at low pressures also has been published, re-

cently, for example by Linackers et al. (1997) and Skandan et al. (1999). A new

aspect was introduced by the evolving field of nano-materials for which powders

produced by aerosol methods are used as starting materials (for reviews of aero-

sols for materials science see Kodas 1989; Kriechbaum and Kleinschmit 1989;

Siegel 1991; Pratsinis and Kodas 1993; Gurav et al., 1993; Kruis et al. 1998 and

Kodas and Hampden-Smith 1999).

Powders of small grain size, narrow size distribution, low agglomeration and

high purity are required for the production of solid nanocrystalline materials and

the exploitation of size effects in applications. Crystalline powders are usually pre-

ferred because they are well characterized. Additionally, the process should pro-

vide high production rates, large yields, the possibility for a scale up and should be

applicable for a large variety of chemical systems and complex materials such as

composites. Flagan and Lunden (1995) pointed out, that for the production of

nanocrystalline materials from nano-powders produced by gas phase methods not

only the grain size but also the particle size, i.e. the control of the particle mor-

phology or agglomeration, is important. However, a low degree of agglomeration

is usually difficult to achieve at high particle number concentrations which are

necessary for high production rates. These requirements

• small grain size

• narrow size distribution

• low agglomeration

• high purity

2.1 Background 11

• high crystallinity • high production rates

• large yields • the possibility for scale up

are well within the reach of the CVS method.

2.1.2 Silicon Carbide

Silicon carbide is an advanced material (see e.g. Harris 1995 and Haigis 1994)

with a wide spectrum of applications in high performance functional or structural ceramics (Liethschmidt 1993; Somiya and Inomata 1991). Properties include ex-

treme hardness, high mechanical toughness at high temperatures, good heat con-ductivity, chemical resistance (e.g. against corrosion by oxygen at high tempera-

tures) and high temperature semiconductivity. Silicon carbide is therefore used as

hard and wear resistant material, as high temperature material in resistance heating

elements, turbines, combustion chambers and rocket nozzles, in electronics (Ca-

pano and Trew 1997) for high temperature, high power semiconductors (Pensel and Helbig 1990), in computer technology as hard discs and support for multichip

modules and in optics as mirrors (Haigis 1994) with extreme flatness, stiffness and

surface quality. Silicon carbide is also a candidate for the first wall in future nu-

clear fusion reactors because it combines high mechanical strength at high tem-

peratures with a very low degree of neutron activation (see for example Sharafat et al. 1991).

Sintering to high density is especially difficult in the case of SiC because of low diffusion rates due to strong covalent bonds (Oreskovich and Rosolowski

1976; Krstic 1995 and VaBen et al. 1996). However, this may change at extremely small particle sizes. Ultrafine, cubic P-SiC powders can be produced in the con-

densed phase by reduction of silica with carbon (Changhong 1997) or by pyrolysis of organo-silicon gels (White et al. 1987). It has been prepared by evaporation of the elements (Ando and Ohkohchi 1982) as well as directly from the SiC bulk

(Nariki et al. 1990) but more commonly it is synthesized by gas phase reactions of

compounds containing silicon and carbon (e.g. Cannon et al. 1982; for an exten-

sive listing see Klein 1999).

In the following parts of Chap. 2 we describe the methodology to produce

nanocrystalline powders by CVS and experimental results for silicon carbide. The

particle formation and growth of SiC is modeled in Chap. 3.

12 2 Gas Phase Synthesis

2.2 Experimental Methodology

In general, a CVS reactor consists of five modules: the precursor delivery unit, the

chemical reactor, the powder collector, the pressure control system and the vac-

uum pumps. These modules are based on different principles depending on the

material to be synthesized.

The precursor delivery unit controls the flow of reactants into the reactor (for

an overview see Wahl (1993), Kodas and Hampden-Smith (1994), Schultz and

Marks 1996). The most common method is the use of thermal mass flow meters

(Hinkle and Mariano 1991; Sullivan et al. 1986) for the delivery of highly volatile

or gaseous compounds. However, the operating temperatures and delivery pres-

sure differences are limited to a certain range. On the other hand, their advantage

lies in the simplicity and the direct measurement and control of the mass flow.

This is in contrast to bubblers (Tompa 1996; Hersee and Ballingall 1990) or sub-

limators (Wahl 1993) where a controlled inert carrier gas flow is used to transport

the precursor material to the reactor from a heated precursor reservoir. These

methods can be used for volatile liquids or solids which cannot be directly deliv-

ered by thermal mass flow meters. Their main advantages are versatility and sim-

plicity. However, the precursor mass flow depends on the saturation of the carrier

gas by the reactant material which is difficult to measure or simulate accurately.

Preliminary precursor decomposition into particles is possible in both methods and

must be avoided. Fluctuations in the mass flow can be produced by droplets (Dep-

pert et al. 1994; Sacilotti et al. 1992) originating from the bubbling liquid. This

can result in a broader particle size distribution and highly agglomerated particles.

In case of atomizers (spray generators; Walzel 1988; Kodas and Hampden-Smith

1994) mists are generated through nozzles or by ultrasonic agitation of the liquid

and delivered to the reactor. However, droplets have the disadvantage already

mentioned and are not preferred for the production of powders for nanocrystalline

materials because powders consisting of larger particles with broader distribution

and hard agglomeration are usually formed. However, in spray pyrolysis processes

completely involatile materials like aqueous salt solutions can be delivered. The

disadvantages of the bubblers or sublimators are avoided in direct liquid injection

systems consisting of a liquid mass flow controller (e.g. a micropump) and a flash

evaporator (Sullivan 1994; Rees 1996). The evaporation takes place close to the

reactor and the reactant flows are accurately measured and controlled independent

of the vapor pressure of the precursor. For solid but volatile precursors the use of

solutions is possible. However, this will increase the flow of byproducts which are

the source of impurities in the product, lower the ceramic yields and provide a

2.2 Experimental Methodology 13

large contribution to the heat production in the reactor by combustion of the sol-vents. In our work we usually use metalorganic or alkoxide (Bradley et al. 1978;

Bradley 1989) and related compounds as precursor materials. Compared to chlo-rides they have the advantage that impurities (carbonaceous species) can be easily removed in the subsequent processing steps of the powders.

Many different types of reactors are used. The most important differences are the pressure regime and the supply of energy required for the decomposition of

precursors forming nanocrystalline particles. This energy can be provided ther-mally in a hot wall reactor (Mazdiyasni et al. 1965; Wu and Readey 1987; Littau

et al. 1993; Chang et al 1994), by flames (e.g. Ulrich 1971; Rulison et al. 1996;

McMillin et al. 1996; Skandan et al. 1999), laser radiation (Cannon et al. 1982; Cauchetier et al. 1991; Gonsalves et al. 1992; Besling 1998), plasmas (Rao et al.

1995; Vollath and Sickafus 1992 and 1993) or by photolysis (Elihn et al. 1999). In contrast to CVD where a solid film is formed on a substrate which is easily

removed from the reactor after deposition, the separation of the product (powder particles) from unreacted reactants and byproducts (e.g. water) is very difficult be-

cause all the components are part of the aerosol reactive flow. Powder collectors (for industrial dust filter technology, see Loffler 1988) are based on the following

principles: mechanical filter (filter membrane, cyclone, etc.), thermophoretic col-

lector, electrostatic filter or wet collector (scrubbers). The choice of a collector depends again on the type of product, production rate and conditions of operation.

For CVS processes at low pressure, the high temperature thermophoretic separator (see below) is a better choice than a thermophoretic collector using liquid nitrogen

if condensable vapors (e.g. water) are byproducts of the synthesis reaction. How-ever, in case of the former the gas stream continues to pass over the product where it can be adsorbed and the relatively high temperatures can initiate sintering of the

clusters leading to hard agglomerates. The pressure control system usually consists of an absolute pressure gauge and

a butterfly valve connected by a feedback control system. For low pressure syn-thesis a pumping system for the generation of gas-flow is placed downstream of

the reactor, consisting of a combination of pumps, e.g., a sliding vane pump and a

roots blower to provide low pressure and high pumping speed economically (Wutz

et al. 1989). This is a big disadvantage regarding the probability for scaling up the

process because large volume flows of reacting gases have to be drawn through

the pump. At normal pressures, the gas flow is produced by a high pressure, up-

stream of the reactor, where simpler compressors produce the necessary pressure

difference from clean gases.

14 2 Gas Phase Synthesis

The following process parameters can be adjusted and have to be controlled

during the CVS process. All of these parameters influence the time-temperature

profile of the process:

• reaction temperature

• reaction pressure

• mass flows of reactants and carrier gas

• precursor material (e.g. of different decomposition kinetics and reaction enthal-

pies)

• method of precursor delivery

• type of carrier gas

• reactor geometry (diameter, length and cross section)

The properties of the reactants (precursor materials) are very important because

the decomposition kinetics can change the product composition, properties and

yield and the reaction enthalpies contributes to the heat balance (compare with

Chap. 3). Therefore, a variation of the precursor material can be used to manufac-

ture different materials, one of the major differences with respect to methods

based on physical vapor deposition methods. CVS type processes are flexible due

to their modular construction and the numerous possible precursor materials. The

chemical process allows the synthesis of materials that are not accessible through

physical methods (e.g. SiC). Major disadvantages are the complexity of the proc-

ess in the hydrodynamics and chemical kinetics and the possible collection of by-

products, e.g., from precursor ligands which can lead to a high impurity level.

Figure 2.2 shows the variations of the CVS process that are possible for a bi-

nary system. Depending on the timing of the different precursor flows, doped,

coated or mixed nanocrystalline particles can be produced. These have different

properties and can be processed into nanocrystalline materials with different mi-

crostructures. This flexibility makes it possible to view the nanoparticles itself as

building blocks for novel materials not unlike monomers in polymer chemistry.

Vollath and Szabo (1994) reported on the synthesis of different coated, ceramic

nanoparticles such as zirconia coated with alumina by two sequential microwave

plasma reactors.

V ollath et al. ( 1997) proposed coatings on nanoparticles as barriers for grain

growth and surface modifications for the attachment of organic compounds and

showed an improvement in the magnetic properties of maghemite coated by zirco-

nia. Fotou and Kodas (1997) and Powell et al. (1996) produced relatively large,

pigment type titania particles and coated them sequentially in a hot wall reactor

with thin silica and alumina films with thicknesses of 6 to 20 nm. Konrad et al.

(1999) reported on CVS of europium doped yttria.

2.2 Experimental Methodology 15

• .. ~

IJ-=' = ------.~ -----+ .. ---4t Fig. 2.2. Variations of the CVS process to produce doped (top), coated (middle) and mixed (bottom) powders displayed together with envisioned microstructures of ceramic products of such powders

Two typical, modular CVS reactors which have been extensively used in this work for the production of SiC and oxide ceramics, are shown in Figs. 2.3 and 2.4. For the synthesis of nanocrystalline silicon carbide, tetramethylsilane (Si(CH3) 4 , TMS, NMR grade purity of 99.9%, Merck) is decomposed in a hot wall reactor, schematically shown in Fig. 2.3. The hot zone of the reactor consists of a silicon infiltrated SiC tube with an inner diameter of 18 mm heated by a resistance fur-nace. The TMS precursor has a high vapor pressure at low temperatures (boiling point at 105 Pa is 299 K) and is delivered to the reactor directly through a thermal mass flow controller (MKS Instruments). The precursor reservoir is held at 297 K to supply the heat of evaporation . The mass flow of helium is controlled by a sec-ond thermal mass flow controller. The total, absolute pressure was measured with a capacitance gauge (Baratron, MKS Instruments). The continuous gas flow is produced by a combination of a Roots pump (250m3/h) and a sliding vane pump (65m3/h) and stabilized at different pressures by means of a butterfly valve. The products are collected on a rotating cylinder cooled by liquid nitrogen from which the powder is scraped off and transferred continuously into a container under inert conditions. Production rates are up to 20 g per hour with yields of about 50% of the theoretical conversion of TMS. The low yield arising most probably because

of inefficient powder collection .

16 2 Gas Phase Synthesis

Fig. 2.3. CVS Reactor for the production, collection and transfer under inert conditions of nanocrystalline silicon carbide (T: thermocouple, p: Baratron pressure gauge)

The CVS reactor used for the synthesis of pure and doped nanocrystalline zir-conia powder (Srdic et al. 2000) , is shown schematically in Fig. 2.4. It consists of a bubbler as the precursor delivery system, a hot wall tubular reactor consisting of

an alumina tube with an inner diameter of 19 mm heated by a resistance furnace.

A heated quartz lamp creates a temperature gradient in a cylindrical water-cooled collector where the doped nanoparticles are separated from the gas stream by

thermophoresis. The zirconia precursor is delivered to the reaction zone by bub-

bling a controlled flow of helium gas (99 .99% purity) through the liquid zirco-

nium-tertiary-butoxide precursor (lnorgtech, England) held constant at 353 K. An

additional flow of oxygen (99.95 %purity) entering directly into the reacting zone

assures the complete oxidation of the product. The mass flows of He

(100 cm3/min) and 02 (1000 cm3/min) are measured and controlled by thermal

mass flow controllers (MKS Instruments). The continuous gas flow is established

by a combination of a roots pump (250m3/h) and a sliding vane pump (65m3/h)

and stabilized at a pressure of 1000 Pa by means of a butterfly valve. The reaction

tube with a length of 400 mm was held at 1273 K. In a typical experiment, about

3 g of pure nano-zirconia are obtained per hour, with yields in the range of 55 to

60%.

The CVS reactor used for nanocrystalline zirconia doped with 3, 5, 15 and

30 mol% of alumina (Srdic et al. 2000a) consists of two helium bubblers as the

precursor delivery units. The liquid precursors, zirconium-t-butoxide (ZTB) and

2.2 Experimental Methodology 17

bubbler

pumping system

with precursor 2 (AI-s-butoxide, Y-tetramethylheptanedionate)

Fig. 2.4. CVS Reactor for the production of pure or doped nano-Zr02 (T: thermocouple, p: Baratron pressure gauge, MFC: thermal mass flow controller)

aluminum-s-butoxide (ASB) are held at 353 K and 448 K, respectively. Precursor

vapors , helium and oxygen flows are intimately mixed by diffusion before the re-action is initiated . The alumina content in the powders is controlled by the ratio of

the helium flow rates through the bubblers containing ASB and ZTB. Nanocrystalline zirconia particles coated with alumina (3 and 30 mol%, Srdic

et a!. 1999 and Srdic et a!. 2000b) were synthesized in a modular CVS reactor consisting of two bubblers as precursor delivery units, two serial reaction zones

consisting of two alumina tubes both with inner diameters of 19 mm and sepa-

rately heated by two resistance furnaces, a powder collection zone and a pumping

unit. The liquid precursors, zirconium-t-butoxide, ZTB (lnorgtech, England) and aluminum-s-butoxide, ASB (Merck, Germany) held at 353 K and 448 K are bub-bled with helium as carrier gas (67 seem through ZTB and 33 seem through ASB

for the sample containing 3 mol% Al20 3 and 50 seem through ZTB and 125 seem through ASB for the sample containing 30 mol% Al20 3) , mixed with oxygen

(1000 seem) and delivered into the corresponding reaction zone. In the first reac-

tor zirconia precursor molecules are decomposed at a wall temperature of 1273 K forming zirconium oxide nanoparticles , which are then used as seeds (substrates

with a very large surface to volume ratio) for heterogeneous growth of alumina in

the CVD mode in the second reactor held at 1173 K. This is an example of a se-

quential CVS process.

Yttria doped zironia was produced in a manner analogous to alumina doped

zirconia, except that instead of a bubbler (ASB), a sublimator was used for the de-

livery of the solid yttrium-heptanedionate precursor (Benker 1999).

Different nanocrystalline materials have been produced by CVS in our group:

SiC and B,C, Al20 3 , Si02 , Ti02 , Fe20 3 , ZnO, Y20 3 , Zr02 , Sn02 , ln20 3 , Eu20 3 ,

18 2 Gas Phase Synthesis

SrTi03 , SrZr03, as well as doped and coated materials. An optimization of there-

actor design (reactor geometry and temperature profile) based on investigations of

the hydrodynamics, heat transfer and particle formation combined with in-situ

process analysis can provide a large improvement in product quality in view of the

simple reactor used for our investigations.

2.3 Experimental Results for Silicon Carbide

2.3.1 Particle Size and Morphology

The crystallographic structure as well as the microstructure (defects in and shape

and arrangement of the primary particles) of some of the powders were investi-

gated by transmission electron microscopy (TEM) with a Philips CM20 Ultra

Twin microscope operating at 200 kV. TEM samples were prepared either by dip-

ping the TEM grid into a dispersion of the powder in a solution of etha-

nol/collodium (volumetric ratio 100: 1) prepared by ultrasonic agitation or simply by dusting the as-synthesized powders onto the grid (Klein 1999 and Buschmann

et al. 1998). A better method would be in situ sampling directly from the aerosol flow (e.g. Dobbins and Megaridis 1987 and Lindackers 1999, p. 25).

The TEM image in Fig. 2.5 gives an overview of silicon carbide powder pro-duced at 1373 K and 2 kPa. It is evident that primary particles well below 10 nm have agglomerated to form secondary particles of much larger size. In the electron

diffraction pattern (not shown), rings corresponding to the 111, 220 and 222 re-flections of ~-SiC, the cubic phase, could be identified (Buschmann et al. 1998).

At higher synthesis temperatures the diffraction rings become sharper as the crys-tallinity increases. The high resolution TEM image in Fig. 2.6 shows nanocrystals

with defects such as twin boundaries and intrinsic stacking faults which are often

observed in silicon carbide. The number of defects decreases with increasing tem-

perature and decreasing reactor pressure.

The distribution of the diameter (size) of primary nanocrystalline silicon car-

bide particles (Fig. 2.7) obtained from TEM images can be fitted by a log-normal

distribution function with a mean diameter of 2.1 nm and a geometric standard de-

viation of 1.26. The statistics are poor because of counting problems due to the

extremely small grain size and extensive overlapping of grains and particles in the

TEM images. The primary particle size is smaller than the grain size obtained by XRD line broadening (3.3 nm) and much smaller than the particle (agglomerate)

2.3 Experimental Results for Silicon Carbide 19

Fig. 2.5. TEM image of nanocrystalline SiC powder synthesized at 1373 K and 2 kPa

Fig. 2.6. High resolution TEM image of nanocrystalline SiC particles synthesized at 1373 K and 2 kPa (enlargement of Fig. 2.5)

diameter of 8.8 nm obtained from the BET specific surface area (compare Sect. 2.3.3). This indicates a very large degree of agglomeration (about 20 primary par-ticles in one agglomerate) .

Different experimental methods measure particle sizes with different weights.

20 2 Gas Phase Synthesis

Fig. 2.7. Size distribution obtained from TEM images of SiC synthesized at 1373 K and 2 kPa; grain size from XRD line broadening and particle size from the specific surface area (BET) are indicated by dashed lines

These different diameters dP can be related by a modified Hatch-Choate equation (Reist 1993) for lognormally distributed sizes given by:

(2.1)

where pis a parameter describing the different weights (e.g. p = 2 for surface and p = 3 for volume or mass weighted distributions), dg is the geometric mean di-ameter which is identical to the median diameter for a lognormal distribution and

2.3 Experimental Results for Silicon Carbide 21

long coagulation times the size distribution of a coagulating aerosol approaches

approximately a log-normal distribution function regardless of the initial size dis-

tribution (Lee et al. 1984). The geometrical standard deviation for such a self pre-serving size distribution (Friedlander and Wang 1966 and Lai et al. 1971) is 1.35 in the continuum regime and 1.46 for free molecular coagulation where the parti-

cle diameter is much smaller than the mean free path of the suspending gas mole-

cules. This self preserving size distribution does not change its shape if plotted in

a dimensionless form regardless of time and physical properties of the dispersing medium (temperature, viscosity etc.) (Friedlander 1977). A coagulating aerosol is

said to have a narrow size distribution if the geometric standard deviation is smaller than the self preserving limit. This can be achieved if the aerosol is pro-duced by a slow chemical reaction, physical condensation or surface reaction on

an existing aerosol or by sintering of coagulating particles (Landgrebe and Pratsinis 1989 and Tsantilis and Pratsinis 2000).

2.3.2 Crystalline Phase, Grain Size and Crystallinity

The crystalline phases of the silicon carbide powders were identified by X-ray

powder diffraction with a Siemens DSOOO diffractometer using a CuKa source

(Fig. 2.8). The crystallite or grain size, dxRD• can be calculated from the XRD line broadening using Scherrer's formula:

K·A dXRD=---/3 ·cosB (2.3)

where K is a shape constant (here assumed as 1), A. is the wavelength of the X-rays

(0.154 nm for Cu-KJ, f3 is the difference of the width (full width at half maxi-mum, b) of the (220) peak at 28 = 59 .9· of the ultrafine powders and of a standard

microcrystalline ~-SiC powder (bs, Aldrich, 99.9%):

(2.4)

This formula can be applied, if the line shape is close to lorentzian i.e. when the

line broadening is dominated by small grain size (for other cases see Ananthara-

man and Christian 1956) and not by strain (Keijser et al. 1983) which is the case

for all SiC samples investigated. This size is actually the volume weighted average

unit cell column length. Krill and Birringer (1998) showed that the distribution of

grain sizes can also be obtained from a detailed analysis of XRD patterns.

The geometric mean diameter can be obtained from the grain size according to

22 2 Gas Phase Synthesis

(2.5)

(Krill and Birringer 1998). The relative statistical errors of the XRD measure-ments are estimated to about 5%. If the full width at half maximum of the diffrac-

tion peak, /3, is larger than twice the distance between the peaks, it is no longer re-solvable. Therefore, the smallest grain size that can be determined in this way is

estimated using the expression

d.= 2KA, mm l'lp · COS (J

(2.6)

to be about 1.4 nm where llp is the difference of the positions of the (111) and

(220) peak in the diffractogram of SiC.

The phase of nano-SiC was determined to be of the cubic polymorph, ~-SiC or

3C (space group F -4 3 m) by electron diffraction and Rietveld analysis of XRD

(Klein 1999) and NDP data (Chatterjee et al. 2000).

The relative degree of crystallinity, r, is estimated by the ratio of the intensities (full width at half maximum, /3, times height, h) of the (220) XRD peak of the ul-trafine powders to a standard microcrystalline powder (subscripts) of ~-SiC (Ald-

rich 99.9%) normalized by the background signal (offset o). This is similar to

methods used in polymer science (Gedde 1995, p. 157):

r= h·/3 -~ (2.7) 0 h, ·/3,

The relative errors in the degree of crystallinity were estimated to about 20%. In

some cases an additional peak in the XRD patterns was found which was due to a graphitic carbon impurity.

Chemical analysis of the powders showed a free carbon content between 9.5

and 12.7 wt% depending on process parameters and storage conditions which cor-

responds to about one monolayer of carbon on the ultrafine particles of nanocrys-

talline silicon carbide. The carbon originates from the decomposition of byprod-

ucts of the TMS pyrolysis, e.g. methane. The TMS pyrolysis reaction is formally:

(2.8)

The by-product methane can decompose into acetylene and carbon with increasing

temperature and residence time (Wu and Ready 1987) according to the following reaction:

(2.9)

2.3 Experimental Results for Silicon Carbide 23

400

7ii" 300 c :::l 0 ~ ~ "Cii 200 c Q)

~

100

0

20 30 40 50 60 70 80

28 [0 ]

Fig. 2.8. X-ray diffractogram of nano-SiC synthesized from TMS at 1573 K and Bragg re-flex positions according to the cubic ~-SiC phase

Klotz et al. (1980) determined that the first step in Eq. (2.9) is rate limiting with an

activation enthalpy of Ea = 985.7 kJ/mol and a preexponential constant of k0 = 10 13/s. Since the decomposition of methane is a subsequent reaction and slower than the formation of silicon carbide (Ea = 282.6 kJ/mol and k0 = 2·1014/s; Clifford et al. 1972), it should be, in principle, possible to avoid or lower the for-

mation of the carbon impurity. Other approaches to this problem could be the use

of a different precursor or the in situ conversion of the carbon by another reactant,

e.g. hydrogen favoring the reverse reaction in Eq. (2.9) (Klein 1999). However,

with hydrogen as co-reactant, elementary silicon is a possible by-product (Agival

and Schieber 1971). Alternatively, carbon dioxide could be used as reactant ac-

cording to the Boudouard equilibrium:

C + C02 2 CO (2.10)

with the forward reaction being favored at high temperatures. There exists an op-

timum content of free carbon in the SiC powders (Clegg 2000) for the production

of sintered SiC ceramics of maximum density. The role of the free carbon is to

remove Si02 from the SiC particle surfaces which can be formed by exposing ul-

trafine SiC powder to air.

24 2 Gas Phase Synthesis

2.3.3 Surface Area, Agglomerate Size and Morphology

The specific surface areas were measured by nitrogen adsorption according to a

multipoint Brunauer-Emmett-Teller (BET) method (Brunauer et al. 1938; Greg

and Sing 1982) using a Quantachrom Autosorb-3B instrument. Samples were de-

gassed at 423 K for 2 hours to remove any adsorbed species from the surface. The

particle size, d8ET• is calculated from the specific surface area, S, using the fol-

lowing expression

6 dBET =--

S·p (2.11)

assuming spherical particles and using the bulk density, p, of ~-SiC (3.21 g/cm3).

This assumption is reasonable because a crystallographic density of

3.2 (0.1) g/cm3 was determined from XRD and found to depend only weakly on

the process parameters.

The specific surface area and, consequently, the particle size are independent of

the degree of crystallinity. The relative statistical errors of the BET and XRD size

measurements are estimated to 1% and 5%, respectively. In the case of XRD

measurements a systematic shift to smaller grain sizes is likely due to other line

broadening effects mentioned earlier. The difference between particle and grain

size is due to the fact that BET measures the surface of a volume filled with a

solid while XRD line broadening detects the coherently diffracting domain size. The ratio of the average particle and grain volume determined by BET and XRD is

used to estimate the number of grains per particle, N, as a simple measure for the

degree of agglomeration:

(2.12)

where it is assumed that grains and particles have the same shape, that the parti-

cles are dense and/or that the porosity is closed or too fine to be detected by nitro-

gen adsorption. The relative errors in the degree of agglomeration are estimated to

about 20%.

The nitrogen adsorption isotherm (Fig. 2.9) of the powder is of type IV

(mesoporous, Brunauer 1945 and Greg and Sing 1982) and is very similar to the

isotherm of an aerogel (Brinker and Scherer 1990, 521 pp.) corresponding to a

very large pore volume with a relatively narrow hysteresis between the adsorption

and desorption branches. However, most of the pore volume is only filled at large

relative pressures. The BET (agglomerate) particle size of the SiC powder was

2.3 Experimental Results for Silicon Carbide 25

500 --+-- pellet (adsorption) -- pellet (desorption) ---o-- powder (adsorption)

400 ---o-- powder (desorption) Cl

M' E 0

300 Q)

E ::I 0 200 >

100

0 0 0.2 0.4 0.6 0.8

p/po

Fig. 2.9. Nitrogen adsorption isotherms of nano-SiC powder (synthesized at 1573 K and 2 kPa) and SiC powder pressed uniaxially into a pellet (700 MPa)

1.6

1.4

1.2

'0 Cl 0.8 0

~ 0.6 -o

0.4

0.2

0

10 100 1000 pore diameter I A

Fig. 2.10. Pore size distribution dv(log d)ldd obtained by BJH analysis of nitrogen adsorp-tion isotherms of nano-SiC powder (as synthesized) and nano-SiC powder pressed uniaxi-ally into a pellet

obtained from the linear part of the adsorption branch at low relative pressures. The desorption branches of the isotherms were analysed using the BJH theory

(Barret et al 1951; Greg and Sing 1982) to extract pore size distributions in the mesopore regime (Fig. 2.10). The isotherm of the powder pressed into a pellet, on

26 2 Gas Phase Synthesis

the other hand, has a considerably smaller pore volume and a larger hysteresis

than the powder and is similar to a (particulate) xerogel (type IV), i.e. the volumi-

nous pore structure surrounded by a highly branched network of particles is at

least partially collapsed. In case of a xerogel, the narrow part (peaks) of the pore

size distribution represent the throat of cavity pores (Brinker and Scherer 1990,

pp. 524). For the pressed pellet the pore size distribution peaks at diameters simi-

lar in size to the primary (3 nm according to XRD) and agglomerate particle size

(9.8 nm according to BET). This can be explained by a microstructure of the pellet

consisting of primary and agglomerate particles which enclose pores of sizes

similar to the particles. The pressed pellet forms a body with a hierarchical ar-

rangement of particles whereas in case of the powder most of the pore volume is

of the size of the agglomerate particles, very few pores being smaller than the

primary particles. The pore size distribution of the SiC powder is very broad as is

already expected from the inspection of the isotherm. This broad pore size distri-

bution is evidence for an object of fractal character (Pfeifer and A vnir 1983). For

surface fractals with fractal dimensions 2 < D, < 3 and for mass fractals with

Dm > 2 the slope of the pore size distribution (dvldd) in a double logarithmic plot

is negative (Ehrburger-Dolle 1998) as is observed for as synthesized SiC powder

(Fig. 2.11). From the slope a fractal dimension of D = 2.5(1) could be determined by fitting with

'0 0.001 > ""0

0.0001 •

100 1000

pore diameter I A

Fig. 2.11. Double logarithmic plot of the pore size distribution d vi dd obtained by BJH analysis of nitrogen adsorption isotherms of nano-SiC powder (as synthesized) and a fit ac-cording to Eq. (2.13)

2.3 Experimental Results for Silicon Carbide 27

(2.13)

Surface and mass fractals can not be distinguished by this method. The SANS (small angle neutron scattering) measurements were performed at

the Intense Pulsed Neutron Source at Argonne National Laboratory on the SAND beamline at room temperature. The scattering curve can be schematically divided into four regions (Fig. 2.12). For the smallest scattering vectors, q [k 1], or longest length scales, the intensity is constant and proportional to the square of the mo-lecular mass, M. In the Guinier region (0 < q·R < 1) the shape of the curve pro-vides a direct measure of the radius of the scattering object (R: radius of agglom-erates in this case). In the intermediate, so called Porod regime (0 < q·r < 1) the shape of the curve depends on the radius of the primary particle (r) and the type of fractal structure that is formed. In the Bragg regime (q·a :2: 1) distinct diffraction

peaks can be observed if the primary particles are crystalline (a being the lattice constant). In the Porod regime the scattering intensity decays by a power law:

(2.14)

where the Porod slope P is

P=-X=-2Dm+D, (2.15)

where Dm (0:::; Dm:::; 3) is the mass fractal dimension and D, is the surface fractal dimension (2:::; D,:::; 3). For non fractal objects Dm = 3 and D, = 2. For mass fractal objects Dm is identical to D, and P = -Dm. For surface fractal objects Dm = 3 and P = D,- 6 (Brinker and Scherer 1990, p. 186). However, it is possible that a polydisperse system can also obey a power law for the scattering intensity, for ex-ample, in the case of power law distribution functions (Martin 1986, Wong and Cao 1992). This may obscure the information on the fractal dimension or in other words there exist two views of the same experimental data. Both interpretations are equivalent if the scattering object consists of a hierarchical structure with larger and larger objects (i.e. a fractal structure or a size distribution). The SANS

data (Fig. 2.13) were fitted with the Schaefer-Hurd model (Schaefer-Hurd 1990):

(2.16)

where the radius of gyration (Guinier radius) or in our case, the agglomerate ra-dius is given by

(2.17)

28 2 Gas Phase Synthesis

limiting IG uinie~ Porod Bragg

logq

Fig. 2.12. Schematic scattering curve for a colloidal aggregate; after Schaefer and Hurd (1990). Symbols are explained in the text

1000

: q = 0.004 A-1 I

0.01 ~-------- q-3.74 '-- - q-2.9 ·.

0.001

0.01 0.1 q•A

Fig. 2.13. Small angle neutron scattering of SiC powder (as received). Fit with power laws and the Schaefer-Hurd-model

2.3 Experimental Results for Silicon Carbide 29