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Composite sol-gel ceramics Yang, Quanzu 1999

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C O M P O S I T E S O L - G E L C E R A M I C S B y Quanzu Yang M . S., Changchun Institute of Optics and Fine Mechanics Chinese Academy of Sciences, 1989 B . S., Dalian Institute of Light Industry, P. R. China, 1985 A THESIS S U B M I T T E D I N P A R T I A L F U L F I L L M E N T O F T H E R E Q U I R E M E N T F O R T H E D E G R E E O F D O C T O R O F P H I L O S O P H Y in T H E F A C U L T Y O F G R A D U A T E S T U D I E S (The Department of Metals and Materials Engineering) We accept this thesis as conforming to the required standard T H E U N I V E R S I T Y O F B R I T I S H C O L U M B I A May, 1999 © Q u a n z u Yang, 1999 In presenting this thesis in partial fulfilment of the requirements for an advanced degree at the University of British Columbia, I agree that the Library shall make it freely available for reference and study. I further agree that permission for extensive copying of this thesis for scholarly purposes may be granted by the head of my department or by his or her representatives. It is understood that copying or publication of this thesis for financial gain shall not be allowed without my written permission. Department of The University of British Columbia Vancouver, Canada DE-6 (2/88) Abstract ABSTRACT ii The fundamental goal of the present study was to develop an understanding of the mechanisms of dispersion, gelation, drying, interfacial bonding and densification of composite sol-gel (CSG) ceramics. The general applied objective was to fabricate high performance C S G ceramics and to produce novel, non-permeable, adherent C S G coatings on stainless steel substrates for high temperature corrosion and wear protection. The properties of the alumina sols and C S G were studied by measuring viscosity, conductivity, ionic strength, and p H of the sol. The dispersion and stability of ceramic particles in alumina sols were investigated by measuring particle size distributions, measuring zeta potentials, and calculating the interaction energy according to D L V O theory. The C S G technology has been developed to fabricate high performance engineering composite ceramics and coatings through dispersing ceramic fillers (alumina, zirconia, S iC) into alumina sols, gelcasting, drying, and pressureless sintering. A sintering model for C S G was developed and validated by experiment results. The model was then used successfully to predict sinterability and to optimize the processing technologies of C S G . The research results indicate that hydrated alumina sols can be used as a sintering and dispersion additive for alumina-based ceramics. The sol-gel matrix provides fast diffusion paths for mass transport during sintering C S G . Dispersion of alumina and S i C particles is substantially improved in alumina sols, as compared to pure water of similar acidity, e.g. the average agglomerate size is decreased by at least 50%. For alumina/alumina C S G ceramics sintered at 1400 °C, the microhardness is 20 G P a and porosity is less than 1 vo l%. The C S G composite with composition of 50vol%SiC-50vol%Al2O3 has been sintered successfully to full densification and microhardness of 22.9 GPa. A novel process for ceramic coatings on the metallic substrates has Abstract iii been developed successfully by combining chemical bonding and C S G technologies. Non-permeable, crack-free, thick ceramic coatings (2-600 pm) on the substrates were fabricated by spraying and dipping, followed by low temperature (500 - 600 °C) sintering. The correlations between the processing methods, microstructure, and mechanical properties of C S G coatings were investigated by varying the preparation methods, studying morphology, and measuring mechanical properties of the ceramics. The chemically bonded C S G coatings have the best performance. The bonding strength between the substrates and coatings is about 42 M P a , and the surface microhardness of the coatings is about 6.5 GPa. Table of Contents iv TABLE OF CONTENTS Abstract ii List of Figures vii List of Tables x v i Nomenclature x v n Acknowledgments x x i CHAPTER 1 INTRODUCTION 1 1.1 Sol-Gel Ceramics 1 1.1.1 Background 1 1.1.2 Advantages and Disadvantages of Sol-Gel Processing 3 1.2 Composite Sol-Gel. 4 1.2.1 Composite Sol-Gel Processing 5 1.2.2 Limitations of Composite Sol-Gel Ceramics 6 1.3 Focus of the Present Study 7 CHAPTER 2 LITERATURE REVIEW 8 2.1 Sol-Gel Processing 8 2.1.1 Sol-Gel Chemistry 8 2.1.2 Alumina Sol 11 2.2 Drying of Gels 14 2.2.1 Driving Forces for Drying 14 2.2.2 Drying Stages 16 2.2.3 Mass Transport Processes 18 2.3 Sintering of Gels 20 2.3.1 Driving Forces for Sintering 20 2.3.2 Viscous Sintering Model 20 2.3.3 Sintering Models for Crystalline Materials 22 2.3.4 Sintering of Alumina Gels 23 2.3.5 Sintering Sol-Gel with Hard Inclusions 25 2.4 Sol-Gel Ceramic Coatings 28 2.4.1 Fabrication of Sol-Gel Coatings 28 2.4.2 Deposition Methods 29 2.4.3 Curing Processes 30 2.4.4 Electrophoretic Deposition of Coatings 31 2.5 Dispersion of Ceramic Powders 33 2.5.1 Zeta Potential of Ceramic Particles 33 2.5.2 Dispersion of Ceramic Powders 35 Table of Contents v 2.6 C o m p o s i t e S o l - G e l ( C S G ) C e r a m i c s 40 2.6.1 Composite Sol-Gel Processing 40 2.6.2 Gelcasting 41 2.6.3 Microstructure of Al 20 3/SiC Composites 43 2.7 C o m p o s i t e S o l - G e l ( C S G ) C o a t i n g s 47 2.7.1 Preparation of CSG Slurry 47 2.7.2 Properties of the Coatings 48 2.7.3 Interfacial Bonding between Substrate and Coatings 49 CHAPTER 3 SCOPE AND OBJECTIVES 53 3.1 S c o p e o f the I n v e s t i g a t i o n 53 3.2 O b j e c t i v e s 54 CHAPTER 4 EXPERIMENTAL METHODOLOGY 56 4.1 S o l - G e l P r o c e s s i n g 57 4.1.1 Alumina Sol 57 4.1.2 Properties of Alumina Sol 60 4.1.3 Dispersion of Ceramic Particles in Alumina Sol 61 4.2 P r o c e s s i n g o f C o m p o s i t e S o l - G e l C e r a m i c s 62 4.2.1 AI2O3/AI2O3 and Zr02/A1203/Al203 Composites 62 4.2.2 SiCAl 2 0 3 / AI2O3 Sol-Gel Composites 62 4.3 P r o c e s s i n g o f C o m p o s i t e S o l - G e l C e r a m i c C o a t i n g s 63 4.3.1 Coating Processing Technology 63 4.3.2 Electrophoretic Deposition 64 4.4 C h a r a c t e r i z a t i o n o f C S G C e r a m i c s 64 4.4.1 Microstructure 64 4.4.2 Properties 65 CHAPTER 5 EXPERIMENTAL RESULTS AND DISCUSSION 68 5.1 C o m p o s i t e S o l - G e l P r o c e s s e s 68 5.1.1 Alumina Sol 68 5.1.1.1 Physical Properties of Alumina Sol 68 5.1.1.2 Electrical Conductivity of Alumina Sol 71 5.1.1.3 Zeta Potential of Ceramic Particles in Alumina Sol... 75 5.1.2 Dispersion of Ceramic Particles in Alumina Sol 79 5.1.2.1 Electrostatic Stability of Ceramic Particles in Alumina Sol 79 5.1.2.2 Steric Stability of Ceramic Particles in Alumina Sol.. 81 5.1.2.3 Interaction Energy of Ceramic Particles in Alumina Sol 82 5.1.2.4 Dispersion of Ceramic Particles in Alumina Sol: Experimental Verification 84 5.1.3 Summary 87 Table of Contents vi 5.2 A l u m i n a / A l u m i n a a n d A l u m i n a / Z i r c o n i a C S G C e r a m i c s 89 5.2.1 Thermogravimetric Analysis of C S G Ceramics 89 5.2.2 Interaction between Sol-Gel Matrix and Ceramic Fillers 90 5.2.3 Properties of C S G Composites 92 5.2.4 Effect of M g O on Sinterability of C S G Ceramics 95 5.2.5 Microstructures of C S G Ceramics 97 5.2.6 Summary 101 5.3 S i C / A I u m i n a C o m p o s i t e S o l - G e l C e r a m i c s 102 5.3.1 Mechanical Properties of S i C / A l 2 0 3 C S G 102 5.3.2 Microstructure of S i C / A l 2 0 3 C S G 106 5.3.3 Summary I l l 5.4 C o m p o s i t e S o l - G e l C e r a m i c C o a t i n g s 113 5.4.1 Processing of the C S G Coatings 113 5.4.2 Post Deposition Treatment of C S G Coatings 117 5.4.3 Microstructure and Properties of C S G Coatings 120 5.4.4 Summary 125 CHAPTER 6 SINTERING MODEL FOR COMPOSITE SOL-GEL CERAMICS 127 6.1 S i n t e r i n g o f A l u m i n a G e l 127 6.2 S i n t e r i n g M o d e l f o r C o m p o s i t e S o l - G e l ( S M C S G ) 130 6.2.1 Ideal Composite Sol-Gel Systems 130 6.2.2 Viscous Sintering of C S G 134 6.2.3 Formulation of Sintering Model for C S G 135 6.3 E x p e r i m e n t a l V e r i f i c a t i o n o f the S i n t e r i n g M o d e l f o r C S G 137 6.3.1 Model Coefficients Calculated from Experimental Data 137 6.3.2 Validation of Sintering Model for C S G 144 6.4 D i s c u s s i o n s o f S i n t e r i n g M o d e l f o r C S G 147 6.4.1 Analysis of the Sintering Model 147 6.4.2 Microstructure of C S G at the Intermediate Sintering Stage 154 6.5 F i n a l S t a g e o f S i n t e r i n g C S G 157 6.6 S u m m a r y 161 CHAPTER 7 SUMMARY AND CONCLUSIONS 163 7.1 S u m m a r y 163 7.2 C o n c l u s i o n s 165 CHAPTER 8 RECOMMENDATIONS FOR FUTURE WORK 169 REFERENCES 171 APPENDIX I PARAMETERS OF SINTERING MODEL FOR CSG 179 APPENDIX II SINTERING STRESS IN CSG 181 List of Figures vii Figure 1.1-1 Figure 1.2-1 Figure 2.1-1 Figure 2.1-2 Figure 2.2-1 Figure 2.2-2 Figure 2.3-1 Figure 2.3-2 Figure 2.3-3 Figure 2.3-4 Figure 2.4-1 Figure 2.4-2 Figure 2.4-3 LIST OF FIGURES Schematic diagram of the sol-gel processes. 2 Schematic diagram of composite sol-gel processing: (a) dissolving of 5 organic-metallic compound into solution; (b) dispersion of ceramic fillers into the solution; (c) formation of composite sol; (d) formation of networked composite gel by gelation; (e) formation of composite aerogel by drying; (f) formation of composite sol-gel ceramics by sintering. Schematic diagram of bond formation between silica particles. 10 (a) the A l i 3 unit, (b) horizontal slices of AI13 unit in (a) depicted from the 12 top looking downward (B = bottom layer, M = middle layer, T = top layer), (c) horizontal slices of the AI13 unit in (a) depicted from the bottom looking upward. To prevent exposure of the solid phase (A); the liquid must adopt a curved 15 liquid/vapor interface (B); Resulting compressive forces on the solid phase cause shrinkage. Schematic illustration of drying process: black network represents solid 17 phase and shaded area is liquid fill ing pores. Kinetics for transformation of a -AhCVseeded boehmite to a - A l 2 0 3 at 24 1025 °C . Density of seeded gels as a function of temperature for 100 min heating 25 time. Plot of normalized stress as a function of time for heterogeneity with 27 lower initial density than matrix. Plot of normalized stress as a function of time for heterogeneity with (a) 27 higher initial density and (b) larger initial particle size than the matrix. Flow chart for thin f i lm formation indicating processing steps. 28 (a) F i l m thickness as a function of withdrawal rate for dip-process silica, 29 and (b) F i l m thickness as a function of oxide concentration for dip-process silica. Schematic of the formation of an electrophoretic coating in sol/particle 32 suspension. Note that the zeta potential of deposited particles is reduced, which has been schematically shown as a low charge density on the deposited particles List of Figures viii Figure 2.5-1 Electrostatic potential in the neighborhood of a charged sphere of radius 1 33 cm. The broken line indicates how the measured potential near a conductor would be affected by the image force i f an electron were used as the test charge. Figure 2.5-2 Zeta potential of alumina and silicon carbide in aqueous suspension as 35 function of p H . Figure 2.5-3 Forces acting on and between particles in laminar flow. 36 Figure 2.5-4 Methods of stabilizing colloidal ceramic particles in liquids. 36 Figure 2.5-5 Potential energy diagrams for interparticle energy of van der Waals 38 attraction (VA), steric repulsion (Vs) and electrostatic repulsion (VR). Figure 2.5-6 Viscosity of OC-AI2O3 suspensions from 20 to 58 v o l % solid concentrations 39 as a function of p H . Figure 2.6-1 Processing routes for sol-gel-derived composites. 40 Figure 2.6-2 Gelcasting processing flowchart. 42 Figure 2.6-3 Strength and toughness of A l 2 0 3 - S i C nanocomposites as a function of S i C 44 content (•) by three point bonding test and Vickers indentations; ( • ) by four-point bonding test; (A) by four-point bend test and indentation-strength method; (V) by three-point bend test and notched beams. Figure 2.6-4 S i C particles in intergranular positions: (a) lattice fringe image of 0001 45 base planes (6H, oc-SiC) in medium-sized S i C particle, (b) large S i C particles agglomerated in pores area show rim and residual glass phase. Figure 2.6-5 T E M micrograph of an A l 2 0 3 / 5 v o l % S i C nanocomposite showing an 46 interface. Figure 2.6-6 Composition profiles across interface (automated step scans): (a) S i - K a 46 signal grain boundary; (b) S i - K a signal across Al 2 03-mul l i te -SiC phase boundary. Figure 2.7-1 Homophase boundaries and heterophase boundaries. Homophase 50 boundaries wave on both sides materials of the same composition and structure, where at heterophase boundaries the structure and/or composition of both components is different. Figure 2.7-2 (a) Schematic diagram of mechanical anchor points: dendrite theory of composite to metal adhesion, and (b) schematic diagram of mechanical anchor points: electrolytic theory of composite to metal adhesion. 51 List of Figures ix Figure 2.7-3 Simplified 2-D schematic representation Si02-to-metal bonding: (a) 52 ceramic saturated with substrate metal oxide in the interfacial region to give strong chemical bonding via a "mono-layer"; (b) as above, with chemical bonding via a "bulk" oxide layer; the strength of the resulting system is dependent on properties of this bulk oxide layer; (c) interface is not saturated with metal oxide; only weak bonding via van der Waals forces is achieved. Figure 4.1-1 Schematic diagram of the processing of alumina sol: (A) aluminum 58 isopropoxide particles dispersed into water; (B) aluminum isopropoxide rapidly hydrolyzed on surface to form the shell of aluminum hydroxide which slows down the rate of peptization; (C) aluminum hydroxide on surface dissolved into the solution; (D) clear alumina sol. Figure 4.1-2 The Distribution of hydrolysis products of A l 3 + in 1 M alumina sol as a 59 function of p H . Figure 4.4-1 Cross-section of the permeability j ig . 67 Figure 5.1-4 The viscosity of alumina sol at p H = 4 as a function of the concentration. 68 Figure 5.1-5 The interparticle separation distance of alumina sol cluster as a function of 70 concentration of alumina sol at p H = 4. Figure 5.1-6 The van der Waals attraction energy of alumina sol clusters as a function 70 of interaction distance in alumina sol at p H = 4. Figure 5.1-7 The viscosity of 1 M alumina sol as a function of p H . 71 Figure 5.1-8 The conductivity of 1 M alumina sol as a function of p H . 72 Figure 5.1-9 The conductivity of alumina sol at p H = 4 as a function of concentration. 72 Figure 5.1-10 The normalized radius of hydro-aluminum-ions (or sol cluster) by charge 74 of 1 M alumina sol as a function of p H . Figure 5.1-11 Zeta potential of alumina particles dispersed in 1 M alumina sol and in 76 water solution, as a function of p H . Figure 5.1-12 Zeta potentials of S i C particles dispersed in 1 M alumina sol and in water 76 solution, as a function of p H . Figure 5.1-13 Schematic diagram of ceramic particles interacting with the alumina sol 78 clusters; (a) hydrolyzed alumina particles react with alumina sol clusters; (b) hydrolyzed S i C particles with adsorbed alumina sol clusters. Figure 5.1-14 The van der Waals attraction energy of S i C - S i C and AI2O3-AI2O3 particles in aqueous solution as a function of interparticle separated distance. 80 List of Figures Figure 5.1-15 Schematic diagram of the two approaching ceramic particles with a f i lm 81 thickness of adsorbed sol clusters of H0p/2. Figure 5.1-16 Electrostatic repulsive energy of alumina particles dispersed in 1 M 82 alumina sol and water solution as a function of the interparticle separation distance. Figure 5.1-17 Electrostatic repulsive energy of S i C particles dispersed in 1 M alumina 83 sol and water solution as a function of the interparticle separation distance. Figure 5.1-18 Total interaction energy of alumina particles dispersed in 1 M alumina sol 84 and water solution as a function of the interparticle separation distance. Figure 5.1-19 Total interaction energy of S i C particles dispersed in 1 M alumina sol and 84 water solution as a function of the interparticle separation distance. Figure 5.1-20 The particle size distribution of alumina A16 dispersed in 1 M alumina sol 85 and water solution at p H = 4. Figure 5.1-21 The S E M morphology of A16 dispersed in 1 M alumina sol (a) and water 86 solution (b) at p H = 4. Figure 5.1-22 The particle size distribution of nano-alumina (NA) dispersed in 1 M 86 alumina sol and water solution at p H = 4. Figure 5.1-23 The particle size distribution of S i C dispersed in 1 M alumina sol and 87 water at p H = 4. Figure 5.2-1 Weight loss vs temperature of pure alumina sol-gel and composite sol-gel 90 with 86.2 vol % calcined alumina after drying at 100 °C for 20 hours. Figure 5.2-2 Schematic of the anticipated interaction of the alumina sol with 91 hydrolyzed alumina particles. Figure 5.2-3 Relative density of alumina/alumina composite sol-gel ceramics with 92 different sol-gel matrix content as a function of sintering temperature. Figure 5.2-4 Microhardriess ( H V i K g ) vs sintering temperature for alumina-alumina 93 C S G . Figure 5.2-5 Porosity as a function of sintering temperature for alumina-alumina C S G 94 Figure 5.2-6 Microhardness (HViKg) vs sintering temperature for the composite of 95 Zr02-3wt%Y203 and alumina sol-gel matrix. Figure 5.2-7 The porosity of Zr02/Alumina composites as a function of sintering 95 temperature. List of Figures Figure 5.2-8 Figure 5.2-9 Figure 5.2-10 Figure 5.2-11 Figure 5.2-12 Figure 5.2-13 Figure 5.2-14 Figure 5.3-1 Figure 5.3-2 Figure 5.3-3 Figure 5.3-4 Figure 5.3-5 Figure 5.3-6 Figure 5.3-7 Figure 5.3-8 xi Microhardness ( H V i K g ) vs M g O content in alumina sol-gel matrix phase, 96 for C S G sintered at 1300°C and 1400°C. Porosity vs M g O content in alumina sol-gel matrix phase,for C S G sintered 96 at 1400°C. S E M micrograph of C S G with 86.2 v o l % calcined alumina sintered at (a) 97 1300°C and (b) 1400°C for 3 hours. Bar=2 pun. S E M micrograph of C S G with 86.2vol% calcined alumina sintered at 98 1400°C for 3 hours, with different M g O contents in the sol-gel matrix: (a) 0.5 mol% M g O and (b) 1.0 mol% M g O , 2.0 mol% M g O . Bar=2 pm. The X-ray map of the M g O distribution in the alumina/alumina C S G . 99 S E M morphology of alumina/zirconia composites: (a) 6 4 v o l % Z r 0 2 - 100 3 6 v o l % A l 2 0 3 and (b) 5 0 v o l % Z r O 2 - 5 0 v o l % A l 2 O 3 composites. Back Scatter picture of S E M of Zr0 2 /a lumina sol-gel matrix. The lighter 101 phase is alumina grains and sol-gel matrix phase and the darker phase is Z r 0 2 . Microhardness of both types of A l 2 0 3 - S i C composites: dispersed in 103 alumina sol ( C S G , closed points) and dispersed in water ( C W D , open points), as a function of S i C content, sintered at 1850 °C for 1 hr. Microhardness of S i C - A l 2 0 3 sol-gel composites only (CSG), as a function 103 of S i C content, sintered at 1700°C, 1800°C, 1850°C and 1900°C for 1 hr. Relative density of both types of A l 2 0 3 - S i C composites: dispersed in 104 alumina sol ( C S G , closed points) and dispersed in water ( C W D , open points), as a function of S i C content, sintered at 1850 °C for 1 hr. Schematic interactions of hydrated silica f i lm on the surface of S i C 104 particles with the alumina gel. The phase diagram of system A l 2 0 3 - S i 0 2 . 106 The S E M morphology of 2 0 S i C - 8 0 A l 2 O 3 sol-gel composites sintered at 107 1700°C for 1 hour. The S E M morphology of 5 0 S i C - 5 0 A l 2 O 3 sol-gel composites sintered at 107 1700°C for 1 hour. Triple junction morphology ( S i C / S i C / A l 2 0 3 ) for 50vol%SiO 2 - 108 5 0 v o l % A l 2 O 3 composite dispersed in alumina sol, and sintered at 1850 °C for 1 hour. S G is alumina sol-originating integranular phase. List of Figures Figure 5.3-9 Alumina sol-originating coating on surface of S i C particles (sample from 108 Fig . 5.3-8). Figure 5.3-10 Triple junction and micropore morphology for 50vol%SiO2-50vol%Al2C»3 109 composite dispersed in alumina sol, and sintered at 1800 °C for 1 hour. S G is alumina sol-originating integranular phase. Figure 5.3-11 S i C particles in intragranular positions in alumina matrix. 110 Figure 5.3-12 The grain distribution of S i C and A 1 2 0 3 in the composites. 110 Figure 5.3-13 Agglomerates of S i C particles joining through necks rather than grain 111 boundaries, in 50vol%SiO 2 -50vol%Al 2 O3 composite, water-dispersed ( C W D ) , sintered at 1850 °C for 1 hr. Figure 5.4-1 Thickness of C S G coatings as a function of withdrawal speed of the 114 substrate at a constant viscosity of the liquid. The solid line presents the result calculated using Eq.(5.4-1). The points present the experimental data. The dotted line is the calculation result of Eq.(5.4-1) after constant c was changed from 0.8 to 0.6. Figure 5.4-2 Thickness of C S G coatings as a function of the viscosity of l iquid at the 115 constant withdrawal speed of the substrate. The solid line presents the result calculated using Eq.(5.4-1). The points present the experimental data. The dished line is the calculation result of Eq.(5.4-1) after constant c was changed 0.8 to 0.6. Figure 5.4-3 Schematic diagram of curing process of composite sol-gel coatings: (a) the 116 coating before drying; (b) the coating dried at 100°C for 1 hour; (c) the coating sintered at 550°C for 20 min. Figure 5.4-4 Infiltrated porosity of C S G coatings as a function of infiltrating time with 117 0.5 M alumina sol and at 200 mTorr vacuum. Figure 5.4-5 The weight gained and infiltrated porosity as a function of the 118 concentration of alumina sol at the constant infiltration time (20 min) and vacuum (200 mTorr). Figure 5.4-6 The gas permeability of C S G coating as a function of concentration of 119 alumina sol with multiple impregnations. Figure 5.4-7 The morphology of composite sol-gel coatings fabricated by the methods 120 from A to D . Figure 5.4-8 The morphology of the composite sol-gel coatings fabricated by the method from the E to H. 121 List of Figures xiii Figure 5.4-9 The surface microhardness of composite sol-gel coatings as a function of 122 different processing methods Figure 5.4-10 The bonding strength between composite sol-gel coatings and substrates 122 as a function of different processing methods Figure 5.4-11 The weight gain for T i (coated and uncoated) as a function of test time at 125 600°C and 800°C. Figure 6.1-1 Sintering as a process of microstructural change involving contributions 127 from: densification and coarsening. Figure 6.1-2 Linear shrinkage of alumina gel doped with 2vol% OC-AI2O3 as a function 129 of sintering temperature (30 min hold at each temperature). Figure 6.1-3 Relative density of alumina gel doped with 2vol% OC-AI2O3 as a function 129 of sintering temperature (30 min hold at each temperature). Figure 6.2-1 Two-sphere sintering model of composite sol-gel particles with the 131 development of the interparticle bond during sintering, a) starting with a point contact at sol-gel coating layers of inclusions, b) neck growth creates a new grain boundary at particle contact of sol-gel matrix layers. Figure 6.2-2 Relative density of composites with different initial relative density as a 133 function of linear shrinkage. Figure 6.3-1 Relative densities of AI2O3/AI2O3 sol-gel composites sintered isothermally 137 at 1350°C as a function of sintering time. The lines were calculated using S M C S G and the points are experimental data. Figure 6.3-2 A plot of Ln(p - po) as a function of Ln(t); experimental data are same as 138 inFig.6.3-1. Figure 6.3-3 A plot of Ln{p -po) as a function of 1/T. 139 Figure 6.3-4 Coefficient Ln(b) of the S M C S G as a function of volume fraction of sol- 140 gel matrix of composites. Figure 6.3-5 Coefficient B of the S M C S G as a function of volume fraction of sol-gel 141 matrix of composites Figure 6.3-6 The relative densities of AI2O3/AI2O3 sol-gel composites with different 141 initial relative densities sintered isothermally at 1350°C as a function of sintering time. The lines were calculated using S M C S G and the points are experimental data. List of Figures Figure 6.3-7 Figure 6.3-8 Figure 6.3-9 Figure 6.3-10 Figure 6.3-11 Figure 6.3-12 Figure 6.3-13 Figure 6.4-1 Figure 6.4-2 Figure 6.4-3 Figure 6.4-4 Figure 6.4-5 Figure 6.4-6 xiv A plot of Ln(p - po) as a function of Ln(t); experimental data are same as 142 in Fig.6.3-7. Coefficient (3 as a function of the initial relative density of C S G 142 The relative densities of AI2O3/AI2O3 sol-gel composites sintered 144 isothermally at 1250°C as a function of sintering time. The lines were calculated using S M C S G and the points are experimental data. The relative densities of A1203/A1203 sol-gel composites sintered at 144 1300°C as a function of sintering time. The lines were calculated using S M C S G and the points are experimental data. The relative densities of AI2O3/AI2O3 sol-gel composites sintered at 145 1450°C as a function of sintering time. The line was calculated using S M C S G and the points are experimental data. The relative densities of alumina gel doped 2wt% (X-AI2O3 sintered 145 isothermally at 1200°C as a function of sintering time. The line was calculated using S M C S G and the points are experimental data. The relative densities of composites with different sol-gel matrix contents 146 sintered isothermally at 1200°C as a function of sintering time. The densification rate of C S G as a function of sintering time for different 148 heating rates at constant initial density and volume fraction of sol-gel matrix. The densification rate of C S G as a function of heating rate for different 149 sol-gel matrix content at a constant initial density after the samples were sintered from 1000°C for 2 min. The densification rate of C S G as a function of sintering temperature for 149 different sol-gel matrix content at constant initial density and heating rate The densification rate of C S G as a function of sintering time for different 150 initial densities at constant heating rate and volume fraction of sol-gel matrix. The densification rate of C S G as a function of sintering time for different 151 sol-gel matrix contents at constant initial density and constant sintering temperature. The relative density of C S G as a function of sintering temperature for 152 different heating rates at constant initial density and volume fraction of sol-gel matrix. List of Figures Figure 6.4-7 Figure 6.4-8 Figure 6.4-9 Figure 6.4-10 Figure 6.4-11 Figure 6.4-12 Figure 6.5-1 Figure 6.5-2 Figure 6.5-3 Figure 6.5-4 Figure 6.5-5 XV The relative density of C S G as a function of sintering temperature for 152 different volume fractions of sol-gel matrix at constant initial density and heating rate. The relative density of C S G as a function of sintering time for different 153 initial densities at constant volume fraction of sol-gel matrix and constant heating rate. The S E M morphology of AI2O3/AI2O3 sol-gel composites with 14vol% 154 sol-gel matrix sintered at 1350°C for 10 min. The S E M morphology of AI2O3/AI2O3 sol-gel composites with 14vol% 154 sol-gel matrix sintered at 1350°C for 30 min. The S E M morphology of AI2O3/AI2O3 sol-gel composites with 14vol% 155 sol-gel matrix sintered at 1350°C for 60 min. The S E M morphology of AI2O3/AI2O3 sol-gel composites with 14vol% 155 sol-gel matrix sintered at 1350°C for 120 min. Morphology of A 1 2 0 3 / A 1 2 0 3 sol-gel composites with 14vol% sol-gel 157 matrix sintered at 1450°C for 30 min. The smaller grains of sol-gel matrix are located at the triple junctions and/or grain boundaries of large ceramic particles. G 2 - Gg2 as a function of sintering time, where G is average grain size and ^ 9 G0 is average initial grain size. Morphology of AI2O3/AI2O3 sol-gel composites with 14vol% sol-gel 159 matrix sintered at 1450°C for 50 min. It is shown that the content of the smaller grains of the sol-gel matrix has decreased. Morphology of AI2O3/AI2O3 sol-gel composites with 14vol% sol-gel 160 matrix sintered at 1450°C for 100 min. The content of the smaller grains of sol-gel matrix further decreases. Morphology of AI2O3/AI2O3 sol-gel composites with 14vol% sol-gel 160 matrix sintered at 1450°C for 420 min. The smaller grains of sol-gel matrix disappear and the larger grains grow larger. List of Tables LIST OF TABLES XVI Table 4.1-1 The equilibrium constants for chemical reactions (5.1-4) to (5.1-8) at 25 59 °C. Table 6.3-1 The experimental data fit into S M C S G model at 1350°C. 138 Table 6.3-2 Ln(p - po) as a function of 1/T for C S G samples sintered for 1 hour. 139 Table 6.3-3 The coefficients b and B as a function of volume fraction of sol-gel 140 matrix. Table 6.4-1 Influences of parameter of Eq.(6.4-1) on densification rate of C S G . 151 Table 6.4-2. Influences of parameter of Eq.(6.2-21) on relative density of C S G . 153 Table 6.4-3 Relative density as a function of initial relative density of C S G sintered 156 isothermally at 1350°C for 3 hrs. Nomenclature NOMENCLATURE xvii Latin Symbols A Hamaker constant Ac constant in Eq.(6.2-4) Ar area of the sample A constant a radius sphere B coefficient in Eq.(6.2-21) b coefficient in Eq.(6.2-21) C constant vc concentration gradient c molar concentration of the solution Ci constant D particle diameter in Eq.(6.2-4) DP gas permeability Ds distance between two ions or particles d distance e charge of electron F Faraday constant Ff form factor of two atoms Fo open porosity Pj total porosity G grain diameter Gc apparent shear modulus of porous C S G Gs shear modulus XVI1 Nomenclature xviii H separation distance H0 interparticle separation distance H0P minimum separation distance h constant in Eq . (6.2-19) ht thickness of f i lm I ionic strength J flux of fluid of liquid JD diffusion flux K bulk modulus gas permeability k Boltzmann's constant lcc constant L linear length M modulus of a porous system Mo modulus at full density m, constant NA Avogadro constant n relative to reciprocal of pore size P surface tension PA ambient vapor pressure Pc capillary pressure Po initial pressure Pv evaporation pressure VPL gradient in pressure in liquid P coefficient in Eq.(6.2-18) Q activation energy for viscous flow in E q (6.2-16) Qk flow rate of gas Qn equilibrium constant R radius of coated particle in Eq.(6.2-1) xviii Nomenclature xix Rg gas constant r distance r from sphare surface S specific surface area of porous body T temperature t time V withdraw speed in Eq.(5.4-1) VA interparticle energy of van der Waals attraction Vm volume fraction of sol-gel matrix Vp volume ratio VR interparticle energy of electrostatic repulsion Vs interparticle energy of steric repulsion in Figure 2.5-5 Vs volume fraction of solid in Eq . (6.2-17) Vr Total interaction energy VE rate of evaporation v m volume of sol-gel matrix U grain boundary energy u dummy variable. ui concentration of ion i Wt weight fraction Wad work of adhesion X neck diameter Y constant Z, charge of ion Greek Symbols a coefficient of thermal expansion ac coefficient of thermal expansion of composites YLV liquid/vapor interfacial energy Ysv solid/vapor interfacial energy YSL solid/liquid interfacial energy £ 0 dielectric permittivity of vacuum er relative permittivity. x i x Nomenclature xx £ linear strain rate uniaxial strain rate differential viscosity of liquid reference viscosity of full densification material, r\oh reference viscosity of liquid Ke conductivity of ion K reciprocal width of electrical double layer X coefficient in Eq.(6.2-18) chemical potential e contact angle p relative density Ps density of the solid skeleton pb bulk density <7 viscoelastic sintering stress Cm sintering stress on matrix characteristic time zeta potential potential I sintering potential Abbreviation C S G composite sol-gel C W D composite without sol-gel matrix D L V O Derjaguin-Landau-Verwey-Overbeek E P D electrophoretic deposition S E M scanning electron microscope S M C S G sintering model for composite sol-gel T E M transmission electron microscope X R D x-ray diffraction XX Acknowledgments ACKNOWLEDGMENTS xxi I would like to express my sincere gratitude to my advisor, Associate Professor Dr. Tom Troczynski of the Metals and Materials Department of University of British Columbia, for his collaboration and invaluable guidance and encouragement through out the course of this work I would like to thank Dr. Sue Bradley, Dr. Geoff Kelsal l , and Dr. George Oprea for their help and suggestions on this work. I would like to thank graduate student Carolyn Moorlag for her help and proof reading of the manuscript. M y thanks are extended to all the members of the Ceramic Group at M M A T for their friendship and pleasant discussions, along with special thanks to Mary Mager for her cheerful help in the running S E M and T E M , as well as M M A T Secretary Joan Kitchen for her help. I also would like to thank the Natural Science and Engineering Research Council of Canada ( N S E R C ) and Ceramic Industry Ltd . for their financial supports. Especially, I would like to express my sincere gratitude to my wife Donghui and my daughter Michelle (Yuqiao) for their unfailing supports and understanding throughout this long process! Chapter I Introduction 1 CHAPTER 1 INTRODUCTION Sol-gel processing is a chemical route for ceramic synthesis involving (aqueous) precipitation of sub-micrometer aggregates of nanometer size primary particles. It has the highest profile among these routes as measured by the number of publications in the scientific and technical literature. Compared with conventional ceramic powder processing using solid-state reactions between powder reactants, sol-gel has the potential to yield ceramics with tailored properties and with numerous advantages over conventional materials. This is because using sol-gel allows chemical control in the preparation of precursors and their conversion to oxide products. Precursors are very important in sol-gel processing as they affect both the way a sol-gel process is carried out as well as the economic viability of scale-up operations. 1.1 Sol-Gel Ceramics 1.1.1 Background Contemporary sol-gel processing emerged during the 1960s, primarily as a result of specialized requirements for ceramic nuclear fuels [1,2]. The most severe constraint on industrial sol-gel processing was, and still is, cost. However, during the past 20 years, the demand for advanced ceramics with high purity, homogeneity and well-controlled, tailored properties has led to a renewed interest in sol-gel technology [3-7]. Traditionally, the term sol-gel (solution-gelation) was applied to a solution process that initially formed a colloid followed by a gel phase. Recently the term has been used to include any solution processes involving hydrolysis and formation of gel, irrespective of whether an intermediate colloid is formed. There is considerable Chapter I Introduction 2 interest today in the application of sol-gel technology for the production of powders, spheres, coatings, and fibers for a variety of applications [8]. A sol is a dispersion of solid particles in a liquid phase where the particles are small enough to remain suspended indefinitely by Brownian motion [8]. For aqueous sols, this means a particle size less than approximately 1 p:m. Sols are classified as lyophobic if there is a relatively weak solvent/particle interaction and lyophilic if this interaction is relatively strong. A gel is a solid containing a liquid component and an internal network structure so that both the liquid and solid are in highly dispersed state [9]. Not all sols can be converted to gels. An important criterion for gel formation is that there be strong particle/solvent interaction so that at least part of the solvent is bound [10]. Their Products Figure 1.1-1 Schematic diagram of the sol-gel processes [11]. Chapter I Introduction 3 Typical colloidal sol-gel processes are represented schematically in Figure 1.1-1 [11]. The precursor material (either an inorganic salt or metal alkoxide solution) is chemically processed to form hydrous metal oxides or hydroxides. Colloidal dispersions (sols) of the hydrolysate are prepared by peptization, then a gel is formed by dehydration or p H control, and the resulting body is calcined to form the stable product. The technique is generally confined to those hydrolysable metal-ion species that produce aqueous sols, e.g. S i 0 2 , A 1 2 0 3 , Z r 0 2 , T i 0 2 , C e 0 2 etc. 1.1.2 Advantages and Disadvantages of Sol-Gel Processing The Advantages which sol-gel processing has over conventional melting can be listed as follows [5,12-16]: 1. Colloidal particle compacts have high surface energies, allowing sintering well below the melting temperature. Lower sintering temperatures not only mean lower energy costs, but also translate into higher purity glass because less metal oxide contaminants are likely to be released from refractory-lined furnaces. This is particularly true for fused silica, which requires a melting/forming temperature of about 2000 °C (3630 °F). 2. High-purity raw materials, such as tetraethylorthosilicate (TEOS), containing less than 100 ppb total metals do exist. If care is exercised, these levels can be retained or reduced in the final product. 3. Improved homogeneity of multicomponent species can be obtained by blending a variety of metal alkoxides, colloidal dispersions, or easily diffused soluble salts. The primary concerns are whether the reactivities of the various species can be controlled to produce the desired level of homogeneity, and whether that distribution can be retained throughout the remaining processing steps. Chapter 1 Introduction \ 4. Low-temperature sintering of near-molecularly dispersed components could lead to new non-crystalline compositions that might otherwise phase separate or crystallize i f produced by the conventional melting approach. 5. Gelation permits the molding of near-net shapes in applications where machining of those shapes is very costly. Although shrinkage factors must be considered, the shape and surface configuration wi l l be retained, despite the large dimensional change. 6. Special materials or configurations have been developed, such as films that are easily applied in thicknesses under 1 pim, controlled-size spherical powders, fibers, and others. Disadvantages of sol-gel processing include [12-17]: 1. The colloidal gel monoliths have very small pore structures and relatively low densities. Removal of the solvents from these open networks and the overall shrinkage in processing require special techniques to avoid cracking. In addition, thermal processing must take into account water and carbonaceous residues that can lead to bloating, residual bubbles or crystal formation i f not properly removed. 2. The high-purity alkoxides are relatively expensive raw materials. 3. Multi-step processing adds time and expense. 1.2 Composite Sol-Gel In order to overcome the disadvantages mentioned above, ceramic powder, fibers and/or whiskers may be dispersed into sols as reinforcement phase (ceramic fillers) to fabricate high performance composite sol-gel ceramics. Examples are composite sol-gel coatings and sol-gel composites. The shrinkage of these bodies is expected to decrease because of the presence of the Chapter 1 Introduction 5 significant amount of inert ceramic powder. Homogenization, stabilization, and dispersion of ceramic particles in liquids are very important steps in producing high quality ceramics. a Solu t ion O o ! o !"o - o o , 0~_ • - o o _ _ • - . -d e f C o m p o s i t e Ge l O C o m p o s i t e A e r o g e l O C o m p o s i t e Sol id F i g u r e 1.2-1. Schematic diagram of composite sol-gel processing: (a) dissolving of organometallic compound into solution; (b) dispersion of ceramic fillers into the solution; (c) formation of composite sol; (d) formation of networked composite gel by gelation; (e) formation of composite aerogel by drying; (f) formation of composite sol-gel ceramics by sintering. b Dispers ion o -- 0 - 0 0 « 0 ~ • o o - -C o m p o s i t e So l 1.2.1 Composite Sol-Gel Processing Sol-gel matrix composites reinforced with ceramic particles, fibers or whiskers have received much attention as high performance materials, for a wide range of engineering applications and ceramic coatings [18-20]. The potential advantages of sol-gel processing for Chapter I Introduction 6 ceramic composites are fine scale mixing and low densification temperature, leading ultimately to improved properties. For example, composites of CC-AI2O3 seeded sol-gel-derived alumina-zirconia were fabricated having sub-micrometer alumina grains and small intergranular zirconia particles of 0.4 p,m average grain size [21]. The different components of sols may be tailored so that they do not react with each other to form new components. A variety of solid phases, such as fine powders or fibers, can be dispersed into a sol before gelation, leading to a composite with good homogeneity and intimate contact between the components. The composite slurry is typically dried at 25-250 °C, and sintered at temperatures several hundred degrees lower than the counterpart calcined ceramics. Typical composite sol-gel processing is described in Figure 1.2-1 [22]. 1.2.2 L i m i t a t i o n s o f C o m p o s i t e S o l - G e l C e r a m i c s A n issue of agglomeration of ceramic particles in sol precursor becomes critical to further improving technology of composite sol-gels. It is also observed that different thermal expansion coefficients of inclusions and sol-gel ceramic matrices result in a stress field around the inclusion [23]. This stress builds up during cooling after hot fabrication. Hence, cracks may be formed from small pre-existing defects at or near the inclusion/matrix interfaces. The composite sol-gel technology can be used to fabricate non-permeable and crack-free thick ceramic coatings (1-100 pirn) on metallic substrates. However, composite sol-gel coatings cannot be densified enough to gain high strength and hardness at curing temperatures below 1000 °C. Some metallic substrates, such as aluminum, magnesium, and alloys, require curing temperatures to be below 600 °C or even lower. Different thermal expansion coefficients of ceramic coatings and metallic substrates result in a stress field build up at interfaces. This Chapter I Introduction 7 thermal stress may cause the ceramic coatings to flake and delaminate at the interface between the coating and the substrate. 1.3 Focus of the Present Study The present investigation focuses on the dispersion of ceramic particles in alumina sol and an explanation of the sintering mechanism of composite sol-gel ceramics. The composite sol-gel technology is further developed to fabricate high performance composite sol-gel ceramics and non-permeable composite sol-gel ceramic coatings. Chapter 2 Literature Review 8 CHAPTER 2 LITER A TURE REVIEW 2.1 Sol-Gel Processing 2.1.1 Sol-Gel Chemistry Sol-gel processes can be divided into two categories: aqueous-based processes that start from a solution of metal salt and alcohol-based processes that start from a metal alkoxide [24]. In the aqueous-based process, the first step is sol formation, which is accomplished by hydrolysis of the metal ions, according to general chemical reaction: M n + + n H 2 0 - » M(OH)„ + n H + (2.1-1) In most cases, the reaction (2.1-1) is driven to the right by the addition of a base. The sol can be prepared by a condensation or a dispersion method. For a condensation method, the sol particles are prepared by slow, controlled nucleation and growth of crystals at an elevated temperature. In a dispersion process, the metal salt is hydrolyzed rapidly at room temperature with an excess of base to form a gelatinous precipitate. The second step in the aqueous-based process is gelation. Gelation of the sol is accomplished by either the removal of water (dehydration gelation) or an increase in the p H (alkaline gelation). A s water is removed during dehydration gelation, the energy barrier to gelation is reduced by the increase in electrolyte concentration in the diffuse layer round individual particles. In alkaline gelation, an increase in the p H reduces the magnitude of positive surface charge on the sol particles which, in turn, reduces the repulsive force between particles and lowers the height of the energy barrier [10]. Chapter 2 Literature Review 9 The alcohol-based process involves reactions with metal alkoxides. Much of the work in this area has focused on the preparation of glasses starting from tetraethyl or tetramethyl orthosilicate. The reactions involve hydrolysis, S i (OR) 4 + n H 2 0 -> S i (OR) ( 4 . n ) (OH) n + n H O R (2.1-2) and condensation, 2Si(OR) (4. f l)(OH)„ -> [Si(OR W O H ) ( „ . ; ) ] 2 0 + H 2 0 (2.1-3) with the overall reaction given by S i (OR) 4 + 2 H 2 0 -> S i 0 2 + 4 H O R (2.1-4) Here there is no distinct sol-formation step, but rather simultaneous hydrolysis and condensation reactions that proceed, ultimately, to the formation of a gel [10]. The silicon alkoxide tetraethylorthosilicate (TEOS) is used most often in sol-gel processes, because it reacts slowly with water, comes to equilibrium as a complex silanol, and in a one-quarter hydrolyzed state has a shelf life of about six months. T E O S , or S i ( O C 2 H s ) 4 is the product of the reaction of S i C l 4 with ethanol. It is a clear, colorless liquid, has a density of about 0.9 g/cm 3 , is easy to handle safely, can be obtained extremely pure when distilled, and is available from several producers [25]. The two other ingredients of TEOS-based sols are alcohol and water. A s soon as T E O S is dissolved in ethanol to make it soluble in water, hydrolysis and polymerization reactions begin [26]. These chemical reactions are [26]: Hydrolysis: S i ( O C 2 H 5 ) + H 2 0 -> S i (OH) 4 + C 2 H 5 O H (2.1-5) Polycondensation: S i ( 0 C 2 H 5 ) + Si(OH) - » = S i - 0 - S i = + C 2 H 5 0 H . (2.1-6) A n inorganic acid is usually used to control the rates of these reactions which are higher at low p H . The solution viscosity increases and the complete hydrolysis of S i ( 0 C 2 H 5 ) 4 would give silicic acid (Si(OH) 4 ) . A condensation reaction can take place between a silanol and an Chapter 2 Literature Review 10 ethoxy group to give a bridging oxygen or siloxane group, =Si-0-Si=. To ensure there is no organic residue in the gel, the amount of water should be in excess of the amount calculated for complete reaction. The choice of solvent is frequently determined by cost and safety, but the more complex compositions are easier to mix in longer-chain alcohols [27]. The bond formation between silica particles is illustrated in Figure 2.1-1 [17]. Figure 2.1-1. Schematic diagram of bond formation between silica particles [17] The processing of zirconia and titania sols is similar to that of T E O S [28]. For example TiC»2-Si02 sol can be prepared using Ti(OC2H5)4 and Si(OC2H5)4 as sources of titania and silica. Zirconium isopropoxide (Zr(OC3H7)4) diluted in isopropanol (C3H7OH) is usually used as of source of zirconia [28]. Following the dissolution of zirconium alkoxide in isopropanol, acetic acid is added to a solution to adjust p H . After 20 min, a homogenous mixture is obtained. Excess water is then added under ultrasonic mixing to complete the reaction. Chapter 2 Literature Review 11 The sol-gel transition is reached when the one-phase liquid becomes a two-phase alcogel of solid plus liquid. The alcogel is an oxide polymer which condenses in the presence of solvent. The term "alcogel" is used to differentiate gels prepared with alkoxides from those prepared from ion-exchanged solutions or colloidal sols (those gels are called hydrogels). The sol-gel transition in alcogels is irreversible and occurs without change in volume. The time of the transition depends on the reaction rate of hydration and polycondensation of the solution. Once through the sol-gel transition, the solvent phase is removed to create xerogels (dry gels) by ordinary evaporation [29]. It is not well understood what exactly takes place during this process at the molecular level. 2.1.2 Alumina sol A n alumina sol is a colloid which consists of a homogeneous medium (H2O + isopropanol) and colloidal particles dispersed therein. The alumina sol clusters are particles that may be formed by polycondensations of A104Ali2(OH)24(H20)i2+? ions. The central tetrahedrally coordinated [AIO4] units are surrounded by twelve edge-linked octahedrally coordinated [AlOg] units [30-36], as illustrated in Figure 2.1-2 [35]. The sets of three octahedra are bound together by bridges at the other end of the shared edges, and these A I 3 O 1 3 units are interlinked by double O H bridges. The average particle size of the sol clusters is approximately 1-2 nm. Aluminum hydroxide sols can be prepared from various aluminum compounds. As early as in 1854, sols were prepared from aluminum acetate and later from AICI3 [28]. Commercial hydroxides can be dispersed in aqueous media with various electrolytes. The process of making transparent monolithic alumina [17], involves four basic steps: (1) Hydrolysis of aluminum alkoxides, (2) Peptization of the hydroxide to a clear sol, (3) Ge l formation, (4) Hydrolysis to alumina. Chapter 2 Literature Review 12 b B T Figure 2.1-2. (a) The A l i 3 unit, (b) Horizontal slices of A l o unit in (a) depicted from the top looking downward (B = bottom layer, M = middle layer, T = top layer), (c) Horizontal slices of the A l o unit in (a) depicted from the bottom looking upward [35]. There are many ways to prepare alumina sols. In the alcohol-based route, aluminum isopropoxide (Al(OC3H 7)3) diluted in isopropanol (C3H7OH) at a temperature of 85 °C for ten hours is used as a source of alumina, and then processed like zirconia sols [37]. B y another method, aluminum isopropoxide or aluminum secondary butoxide ( A K O C ^ H ^ ) is dissolved into excess water under stirring [38]. The molar ratio of water to alkoxide is kept at about 100:1. Deionized water with an initial temperature of 75 °C is used. A c i d additions are generally specified in terms of H + / A 1 ratio. The molar ratio of nitric acid to alkoxide ranges from 0.0154 to Chapter 2 Literature Review \ 3 0. 246. The solution is stirred continuously for 24 hours at temperature of 90 °C, resulting in a clear alumina sol [38]. Yoldas [17] indicated that at least 0.03 moles of acid per mole of alkoxide (or hydroxide) must be added to the slurry to peptize and disperse the system to a clear sol. Complete peptization cannot be attained below this concentration. There appear to be two general requirements for the type of acids listed for alumina sol processing: 1. The anion of the acids must be non-complexing (or very weakly complexing) with aluminum ions at these low concentrations. Therefore, failure of sulfuric and hydrofluoric acids to peptize can be explained in terms of F" and SO4" 2 ions forming complexes with aluminum [39]. 2. The acid must also be sufficiently strong to produce the necessary charge effect in relatively small quantities with respect to aluminum concentration. In other words, the amount of acid in relation to aluminum must not be large enough to prevent the formation of a continuous aluminum bonding through oxygen or through hydroxide [39]. Chapter 2 Literature Review \ 4 2.2 Drying of Gels 2.2.1 Driving Forces for Drying Removal of liquid is particularly troublesome in sol-gel processing because gels tend to warp and crack during drying. Very slow drying rates are required to avoid fracture. The driving forces and mass transport mechanisms are very important during the drying of gels. The driving forces for shrinkage include chemical effects, such as condensation reactions, and physical effects such as capillary pressure [40,41]. Fluid transport can occur by flow down a pressure gradient or diffusion down chemical potential gradients. Deformations of a network involve elastic, plastic, and viscoelastic strains. A s liquid is removed by evaporation, tension that develops in the pores produces contraction of the network. When the pressure is not uniform, warping and cracking of the gel can result [42-48]. (1) Chemical reactions The shrinkage of gels during aging is attributable to ongoing condensation reactions between M - O H groups. The driving force for shrinkage (less than 10% linear) provided by chemical reaction is small compared to other factors operating during evaporation [49]. On the other hand, mechanical properties of gels are profoundly affected by the formation of new bridging bonds. The elastic modulus and viscosity of gels increases during aging, and rises even during drying-induced shrinkage. It seems likely that the contraction brings reactive M - O H groups into proximity so that further condensation is possible, and shrinkage is irreversible [50]. Chapter 2 Literature Review 15 (2) Capillary Pressure If evaporation of liquid from the pores were to expose the solid phase, a solid/liquid interface would be replaced by a solid/vapor interface of higher surface energy. To prevent such an increase in the energy of system, liquid tends to spread from interior of the body to cover that interface. Since the volume of liquid has been reduced by evaporation, the meniscus must become curved as indicated in Figure 2.2-1. r (A) (B) Figure 2.2-1. To prevent exposure of the solid phase (A); the liquid must adopt a curved liquid/vapor interface (B); Resulting compressive forces on the solid phase cause shrinkage [41]. The excess pressure (P) in this liquid is related to the radius of curvature (r) of the meniscus by P = -2%v/r (2.2-1) where yLV is the liquid/vapor interfacial energy (or surface tension). When the center of curvature is in the vapor phase, the radius of curvature is negative and the liquid is in tension (P>0). The maximum capillary tension (PR) in the liquid occurs when the radius of meniscus is small enough to fit into the pore; for liquid in a cylindrical pore of radius a, the minimum radius of meniscus is r = -aJcos(Q) (2.2-2) where 9 is the contact angle. If contact angle 6 is 90°, then the liquid does not wet the solid and the solid/liquid interface is flat (as r —> °o, p —> 0). If the 6 is 0°, then the solid surface is covered with a liquid film. The maximum tension is related to surface-to-volume ratio of pore space, SJ/VP [51,521: PR = (YSV-YSL)SP/VP = yLVcos(d)SI/VP (2.2-3) Chapter 2 Literature Review 16 where jsv and YSL are the solid/vapor and solid/liquid interfacial energies, respectively. The specific surface area of porous body, S, is related to surface-to-volume ratio by [52]: where p is relative density, p = p\Jps, Pb is the bulk density of the solid network (not counting the mass of the liquid), and ps is the density of the solid skeleton (the skeletal density). (3) Osmotic pressure Osmotic pressure is produced by a concentration gradient. If the pores are large, the diffusive flux is matched by a counterflow of liquid toward the interior, and no stress develops. However, i f the pores are small enough to inhibit flow, diffusion away from the interior can produce tension in the liquid in that region. The balancing compression in the solid phase can produce shrinkage [50]. (4) Disjoining pressure Disjoining forces are short-range forces resulting from the presence of a solid/liquid interface. The most important examples are double-layer repulsion between charged surfaces and interactions caused by the structure created in liquid by dispersion forces [53,54]. 2.2.2 Drying Stages (1) Constant Rate Period The stages of gel drying are illustrated in Figure 2.2-2 [41]. The first stage of drying is called the constant rate period (CRP) (Figure 2.2-2B), because the rate of evaporation per unit area of the drying surface is independent of time [55,56]. The evaporation rate is close to that from an open dish of liquid [57]. The rate of evaporation, VE, is proportional to the difference between Py and the ambient vapor pressure, PA. SpWp = SppAl-p) (2.2-4) VE=kc(Pv-PA) (2.2-5) Chapter 2 Literature Review 17 where kc is a factor that depends on the temperature, mass transport rate coefficient, and the geometry of the system. The vapor pressure of the liquid is related to the capillary tension (P) by Pv = P0 exp < PV^ RJ (2.2-6) The fact that the evaporation rate is similar to that of bulk liquid indicates that the vapor pressure reduction is insignificant during the C R P . (A) Initial condition Stages of Drying Liquid/vapor meniscus flat Pore liquid Solid phase (B) Constant rate period Pressure in liquid at exterior: Evaporation Shrinkage (C) Falling rate period Maximum capillary pressure: L P J ySV ySL )SP v p m* Empty pores Minimum radius ol curvature E R Figure 2.2-2. Schematic illustration of drying process of gels: black network represents solid phase and shaded area is liquid filling pores [41]. Chapter 2 Literature Review \ g (2) First Falling Rate Period When shrinkage stops, further evaporation drives the meniscus into the body, as illustrated in Figure 2.2-2C. A s air enters the pores, the surface may begin to lose its translucency [58]. In the first falling rate period (FRP1), the rate of evaporation decreases and the temperature of the surface rises above the wet bulb temperature. Most of the evaporation is still occurring at the exterior surface, so the surface remains below the ambient temperature, and the rate of evaporation is sensitive to the ambient temperature and vapor temperature [56]. The liquid in the pores near the surface remains in funicular condition, so there exist contiguous pathways along which flow can occur. A t the same time, some liquid evaporates within the unsaturated pores and the vapor is transported by diffusion. (3) Second Falling Rate Period A s the meniscus recedes into the body, the exterior does not become completely dry right away, because liquid continues to flow to the outside. A s long as the flux of liquid is comparable to the evaporation rate, the funicular condition is preserved. However, as the distance from the exterior to the drying front increases, the capillary pressure gradient decreases and therefore so does the flux. Eventually it becomes so slow that the liquid near the outside of the body is isolated in pockets, so flow to the surface stops and liquid is removed from the body only by diffusion of its vapor. A t this stage, drying is said to enter the second falling rate period, where evaporation occurs inside the body [41]. 2.2.2 M a s s T r a n s p o r t Proces se s (1) Darcy's law Fluid flow through porous media obeys Darcy's law [59] which states that the flux of fluid of liquid, J, is proportional to gradient in pressure in the liquid, VPL: Chapter 2 Literature Review 19 J = - (D/7]L)VPL (2.2-7) The flux of liquid is in units of volume per area of the porous body (not the area occupied by the liquid) per time, PL is the force per unit area of liquid, T\L is the viscosity of the liquid, and Dp is called the permeability and has unit area. Positive flux moves in the direction of increasingly negative pressure (i.e. the flow is toward regions of greater tension in the liquid). In gels, the pores are so small that a large portion of the liquid may be in structured layers within about 1 nm of a solid surface, so that the effective viscosity may be larger than in the bulk. (2) Diffusion transport According to Fick ' s law, diffusive flux (JD) is proportional to the concentration gradient (PC) [60]: JD = -DVC = -(DcC/RgT) Vp (2.2-8) where Dc is the chemical diffusion coefficient, C is the concentration, and p is the chemical potential. Diffusion can contribute to the shrinkage of gels in special cases (e.g. when the gel is immersed in salt solution) and may be important during evaporative drying, i f a concentration gradient develops in the pores by preferential evaporation of one component of the pore liquid. 20 Chapter 2 Literature Review 2.3 Sintering of Gels 2.3.1 D r i v i n g F o r c e s f o r S i n t e r i n g Sintering is a process of densification driven by decreasing area and hence interfacial energy. Material moves by viscous flow or diffusion in such a way as to eliminate porosity and thereby reduce the solid/vapor interfacial area. In gels, the pore diameters are typically in the order of 1-10 nm, corresponding to large area (500 m /g), so the driving force is great enough to produce sintering at exceptionally low temperatures. Amorphous materials sinter by viscous flow rather than diffusion, so the path along which material moves, and the relationship between the rate of transport and the driving forces, are quite different. Viscous sintering is driven by the energy gained by reduction in surface area of the porous body. The energy gained when this strain occurs is the product of specific surface energy and the change in surface area. When a viscous body flows, energy is expended, and the rate of this dissipation of energy is proportional to the square of the strain rate [61,62]. 2.3.2 V i s c o u s S i n t e r i n g M o d e l Frankel [62] suggested that the rate of strain of a viscously sintering body could be found by equating the rate of change in the surface energy to the rate of energy dissipation. This insight is the basis for all analyses of viscous sintering, which differ only in the models they adopt to represent the geometry of the body. In every case, the same dimensionless group appears: % = ysv t/dr], where d is a characteristic dimension of the structure. Densification is completed when x ~ 1, or t ~ drf/jsv- Therefore, the sintering rate is faster for bodies with smaller particles or pores. In gels, the pore diameters are typically in the order of 1-10 nm, so that they densify Chapter 2 Literature Review 21 orders of magnitude faster than bodies made using conventionally crushed powders (particle size 1=10 pm). Mackenzie and Shuttleworth (MS) analyzed [63] the shrinkage rate of a spherical shell according to Frankel's method [62]. The shell can be used to represent the densification of a body containing spherical pores. The dimensions of the shell are chosen so that the central void occupies the same volume fraction as the pores in the sintering body. This is a much more elegant treatment than the analysis of the coalescence of spheres, because the shell remains spherical as it shrinks. Exact expressions can be written for the change in surface area and the energy dissipated in viscous flow as the shell contracts. The result is [62,64]: .1/3 7] 3 W) P J ( I - P ) 2 / 3 P 1 / 3 dp (2.3-1) The solution of Eq.(2.3-1) is given by: where: = fus(P)-fMs(Po) (2.3-2) fMS(P) = -/3 2 (I±P!_) 1(1 + P)3J - 3I / 2 t a n - ' 2p-l >l/2 (2.3-3) where p = pi/Ps is relative density, Pb is bulk density of porous body, and ps is density of the solid phase. The time scale uses the dimensionless variable (reduced time) ysvn1/3t/ri, where n is the number of pores per unit volume of solid. For a given relative density, the larger n, the smaller the pore size; therefore, n is relative to reciprocal of pore size. This model is ideally applied to the final stage of the sintering, when the pores have become isolated spherical voids. It is considered the cases of insoluble gas and of a gas that dissolves into the solid phase at such a rate that the pressure in the bubble remains constant. If the gas is insoluble, the bubble shrinks 22 Chapter 2 Literature Review until the pressure of the gas equals Pc=2ysv/r, where Pc is the capillary pressure driving sintering and r is the radius of the bubble. If, at the moment when the pore becomes a closed sphere, the radius of the bubble is r0 and the pressure of the gas P0, then the equilibrium size of the bubble wi l l be requ2=P0r03/2ysv [63]. This phenomenon can prevent sintered bodies from reaching full density, so it is advisable to sinter in a vacuum or in an atmosphere of soluble gas. 2.3.3 Sintering Model for Crystalline Materials As for viscous sintering, the driving force for densification of crystalline materials is surface energy. Material tends to move from regions of positive (convex) curvature to the regions of negative (concave) curvature, and this leads to the filling of necks between particles. Models for the initial stage are based on the two-sphere geometry analyzed by Frankel [62]. Surface energy produces an excess of vacancies in the crystal lattice in the regions of negative curvature, and this results in a flux of vacancies away from the perimeter of necks (or equivalently, a flux of atoms toward the perimeter). In the intermediate stage, the structure is assumed to consist of polyhedra with cylindrical pores along the edges. The density is predicted to increase linearly with time when mass transport occurs by lattice diffusion with the grain boundary as a source of atoms [66]: (dp/dtk = 336/T L (2.3-9) where %i is the characteristic time for sintering by lattice diffusion. The densification rate varies as l/o3 when the mass transport path is by lattice diffusion from the grain boundary. The same result is found for the final sintering stage. When diffusion occurs along the grain boundary, rather than through the lattice, the densification rate varies as l/o4, so this path is favored for very small particles. The rate of coarsening by the surface diffusion also varies as l/o4, so the latter two processes are expected to be in competition in gels, 23 Chapter 2 Literature Review because of their small particle size. Which process wi l l dominate depends on the relative magnitudes of grain boundary and surface diffusion coefficients. Since those properties have different activation energies, either one may dominate depending on the temperature of sintering [62]. Simultaneously with sintering, the grains tend to grow to reduce the total area of grain boundaries. The rate of increase of the grain diameter, G is found to obey [66]: G3(t)-G3(t)oct, so dG/dtocl/G2 (2.3-10) This means that grain growth is relatively rapid for small particles. Thus, gels are subject to rapid growth and rapid surface diffusion, both of which tend to inhibit densification. One common method for facilitating sintering of crystalline materials is to introduce a liquid phase. The liquid phase tends to accumulate near the neck between particles, where it helps to dissolve the solid phase, then provides an easy path for transport of atoms away from the points of contact [67,68]. The liquid may also help to "lubricate" the particles, allowing them to rearrange into a denser packing under a capillary pressure provided by the liquid. 2.3.4 S i n t e r i n g o f A l u m i n a G e l s Small particles provide short diffusion paths from the pore to the point of contact between particles. Most importantly, the particles must be densely and uniformly packed to minimize the pore volume while maximizing the number of particle contacts. This creates a situation in which there are many grain boundaries feeding atoms into a small pore space along a short diffusion path. In general, monolithic gels have relatively low densities, so crystalline gels do not sinter well to full density in spite of their very small pore size. Yoldas [69] showed this clearly for monolithic alumina gels with 2.5 nm pores and 7 nm crystallites. Typically, the relative density Chapter 2 Literature Review of the dried gels is 0.37, and when it is sintered at 1200 °C it shrinks only to p = 0.62. The sintering is accompanied by a growth in pore size to 10 nm, and then transformation of amorphous gel to oc-alumina because of the large grain size of the crystalline phase. Kumagai and Messing [70,71] showed that the sintering kinetics of alumina gels are dominated by the microstructure produced during crystallization. The transformation kinetics to OC-AI2O3 for samples seeded with 0.05, 0.15, and 1.5 wt% a - A l 2 0 3 and heated at 1025 °C for various times are compared to unseeded samples in Figure 2.3-1 [71]. The unseeded AI2O3 is = 50% transformed after 3 h and by extrapolation approximately 4.7 h would be required for complete transformation. In contrast the 1.5 % sample is 50 % transformed after 3 min and completely transformed in < 15 min. The 0.15 % sample follows kinetics that are similar to the 1.5 % sample. Clearly, the 0.15 % and 1.5 % seeded samples transform faster because the a-Al 2 C»3 particles act as nuclei for the transformation process. < 1 e i£ f"^~  a ~ A I I ° » : 0 1 5 * ' a - A i t O s : Owt% I025°C I 2 I 3 Time (h) Figure 2.3-1. Kinetics for transformation of a -A^CVseeded boehmite to (X-AI2O3 at 1025 °C [71] Sintered density as a function of temperature for the seeded and unseeded samples of boehmite is presented in Figure 2.3-2 [71]. The density of all samples is at least 65% of Chapter 2 Literature Review theoretical at 1050 °C and for the 1.5 % sample is 72 % of theoretical. This density difference is attributed to the fact that the seeded samples have already commenced sintering whereas the unseeded sample is still transforming. The unseeded sample does not reach > 95 % relative density until 1600°C. However, >98 % relative density is achieved at 1300 °C for the 0.15 % sample and at 1200 °C for the 1.5 % seeded sample. The 1.5 % sample is almost 95 % dense at 1150 °C. Boehmite ( A l O O H ) goes through a complicated series of phase transformation on heating, converting through the series y—» 5 —> 9 and finally to a-alumina [72]. Temperature ("CJ F i g u r e 2.3-2. Density of seeded gels as a function of temperature for 100 min heating time [71] 2.3.5 S i n t e r i n g S o l - G e l w i t h H a r d I n c l u s i o n s Heterogeneities in a green compact typically sinter at different rates than the host powder and thereby affect the sintering body. In particular, stresses develop in association with the differential shrinkage of the heterogeneity. The stresses may cause sintering damage, such as cracks, arrays of voids, or isolated pores [73-78]. The heterogeneities also influence the net sintering rate of the compact. Some of the important trends in stress development and sintering rates have been studied by several authors [79-84]. One of the well-known studies led to Scherer's theory [176], based on a composite sphere model. In this model, the sintering matrix containing multiple inclusions is idealized through a single inclusion surrounded by a matrix 26 Chapter 2 Literature Review whose size is chosen to obtain the same volume fraction of inclusion as in real materials. Hsueh et al. have analyzed the viscoelastic stresses and sintering damage in heterogeneous powder compacts [85]. The development of stress during sintering is illustrated by considering a spherical heterogeneity in an infinite AI2O3 compact. The constitutive laws have been used to analyze densification rate and hydrostatic sintering stress of ceramic matrix. Viscoelastic modeling of sintering is conveniently conducted by establishing an elastic solution to the basic problem and transforming it into the associated time dependent regime using a Laplace transform. Such solutions exhibit the general from Equation [85]: where a is the viscoelastic sintering stress, t is time, Gs and K are the shear and bulk moduli, respectively, 77 is the viscosity, Aes is the uniaxial strain rate differential that establishes the stress, F is a stress relaxation function, and u is a dummy variable. When the heterogeneity has either a smaller density or a smaller particle size than the host, it tends to shrink more rapidly and thus is subject to a tensile stress [86]. This stress is limited by the "sintering potential" of the heterogeneity since sintering stops when stress am —> Z, where am is actual stress determined by differential sintering rates and viscoelastic response of porous materials, and X is sintering potential, i.e., driving force for sintering determined by chemical potential, as illustrated in Figure 2.3-3 [85]. Two stress plateaus are found to occur. The initial rapid increase in stress reflects the initial difference in sintering rate. Then as <7m increases, the sintering of the heterogeneity is impeded and further stress development in the presence of the reduced driving force (crm - X) tends to be relaxed by creep of the host. Heterogeneities with a larger density or particle size than the host are subject to a smaller densification rate and thus experience compression. In this instance, the compression is not (2.3-11) 0 27 Chapter 2 Literature Review limited by the "sintering potential" Z, because the mean heterogeneity stress <rm in the matrix is zero. Hence, net matrix shrinkage is unaffected. Consequently, very large heterogeneity stresses develop as illustrated in Figure 2.3-4 [85]. The period of rapidly increasing stress appears to be associated with an effective relaxation time for the viscosity of the host. T I M E , t <») Figure 2.3-3. Plot of normalized stress as a function of time for a heterogeneity with lower initial density than the matrix [85]. - 3 0 Q -TIME , I (») WITIAL HETEROGENEITY T I M E , t (*) Figure 2.3-4. Plot of normalized stress as a function of time for heterogeneity with (a) higher initial density and (b) larger initial particle size than the matrix [85]. Chapter 2 Literature Review 2.4 Sol-Gel Ceramic Coatings 28 2.4.1 Fabr ica t ion of Sol -Gel Coatings The flowchart of the sol-gel process for monolithic thin films is shown in Figure 2.4-1 [29]. Coatings are formed from a solution using dipping, draining, and spinning methods. Prepare Coating Solution Age overnight Prepare Substrate Wash in alcohol Dip Substrate into Solution I Withdraw Substrate from Solution I Dry at low temperature I Heat Treat I Room Temp«ra tur« -25 cm/min •too°c >300°C Evaluate Figure 2.4-1. Flow chart for thin fi lm formation indicating processing steps [29] The physical properties of the solutions that are monitored include viscosity, surface tension, and time to gelation. The surface tension of the solution is usually measured by rise in a capillary [120]. The surface tension can be adjusted by changing the solvent. Generally, a low surface tension is desirable for coating a substrate with low surface energy [87]. In most cases, the effect of viscosity is far greater than that of the surface tension, particularly in dip coating. Practically, the time to gel is determined by inspection. The time when a solution shows no visible flow is one such criterion. Another practical test that applies to thin coating films is the Chapter 2 Literature Review 29 time when the gel surface shows a fingerprint when touched lightly. Alternatively, the viscosity of the solution is determined and a reasonable value, perhaps 40 mPa-s in an acid-catalyzed system or 1000 mPa-s in a base-catalyzed system, is chosen as indicative of gelation [120]. 2.4.2 D e p o s i t i o n M e t h o d s The majority of sol-gel coatings are applied by a dip coating. This is a simple approach that allows the properties of the solution to control the deposition. A substrate is lowered into a vessel containing the solution. A meniscus develops at the contact of the liquid and the substrate. As the substrate is withdrawn, the meniscus, controlled by the viscosity, surface tension, and time-to-gelation, generates a continuous fi lm on the substrate. The fluid mechanics of film formation in this process has recently been summarized [88]. The expressions relating to fi lm thickness as a function of withdrawal rate and oxide content were derived. Figure 2.4-2a [89] shows that the coating fi lm thickness increases with increasing withdrawal rate. F i g u r e 2.4-2. (a) F i l m thickness as a function of withdrawal rate for dip-process silica, and (b) F i l m thickness as a function of oxide concentration for dip-process silica [89] In Figure 2.4-2b [89], it is shown that the coating f i lm thickness, for a given withdrawal speed, increases with an increase in oxide content [89]. A n increase in oxide content generally means a higher viscosity and shorter time-to-gel, within some reasonable range. Based on these two semi-Chapter 2 Literature Review 30 empirical relationships, it is possible to select the proper fi lm thickness for a dip coating process [89]. For one-side coating, a technique that is often used is spin coating. In this case, the substrate is placed on a spinner, rotated at approximately 200 rpm while the solution is dripped on the center of the substrate. In most cases, the film thickness is between 50 and 500 nm and is controlled by adjusting the sol viscosity. Typical solution viscosities are 3-10 mPa-s and typical surface tensions are 30-50 x 10"3 N / m . Knowing the viscosity and oxide concentration of the solution, a desired fi lm thickness can be achieved by controlling the spinning rate [89]. 2.4.3 C u r i n g P r o c e s s e s After the deposition, the fi lm must be dried at approximately 250 °C. From the time the solution is applied to the time it gels, the coating loses approximately 50% of its weight due to drying. Then, as the gel dries further, there is another reduction in weight again by 50%. The drying is accompanied by a volume reduction of more than 70 vo l%. The goal is to produce a coating which remains adherent and continuous and maintains complete surface coverage. It has been shown repeatedly that all shrinkage is taken up in the thin dimension (perpendicularly to the substrate surface) and not in the plane of the substrate, as long as the thickness is less than about one micrometer [90]. The heat treatment to obtain a hard gel typically takes approximately 30 minutes. Coatings can be dried quickly because water and solvent escape easily through interconnected pores within 1 | i m thin fi lm. The pores remain open to the surface until the coatings are fired at temperatures above 600 °C. The microporosity in silica sol-gel coatings is not removed entirely up to 1000 °C, but it may already behave as an oxidation barrier or passivation coating at 600 °C [91]. This is because the pores, which are only 1-5 nm in diameter, limit diffusion of gas to the surface. This ability of Chapter 2 Literature Review 31 the microporous f i lm to behave in many ways like the bulk oxide is an attractive property of the sol-gel technology. In the cases where some protection is desired from the porous oxide fi lm, a heat treatment at 500 °C is sufficient. Especially in cases where higher temperature would degrade the substrate, or may exceed its softening point, a porous sol-gel coating can be applied. 2.4.4 E l e c t r o p h o r e t i c D e p o s i t i o n o f C o a t i n g s Electrophoretic deposition (EPD) is a colloidal process wherein ceramic bodies are shaped directly from a stable colloid suspension by a dc electric field. The depositing electrode is the shape of the ware required, and it is designed such that deposit release is facilitated. The electrophoretic deposition process is schematically shown in Figure 2.4-3 [92]. A dc field causes the charged particles to move towards, and deposit on, the oppositely charged electrode. E P D is a combination of two processes: electrophoresis and deposition. Electrophoresis is the motion of charged particles in a suspension under the influence of an electric field. Deposition is the coagulation of particles to a relatively dense mass [93-100]. Chapter 2 Literature Review 32 F i g u r e 2.4-3. Schematic of the formation of an electrophoretic coating in sol/particle suspension. Note that the zeta potential of deposited particles is reduced, which has been schematically shown as a low charge density on the deposited particles [92]. Surface charge is required for particle stability in suspension. Charged particles attract oppositely charged ions around the particle. This atmosphere is defined as "lyosphere". When the particles move in the liquid, some ions in the lyosphere fluid mechanically "shear off ". The potential at plane which defines the boundary between the moving and stationary phases is defined as the zeta potential. The higher the zeta potential, and more stable the suspension against coagulation. Interactions that occur between charges fixed at a particle surface and those free in the solution are important to the stability of colloidal systems [101,102]. Chapter 2 Literature Review 2.5 Dispersion of Ceramic powders 33 2.5.1 Z e t a P o t e n t i a l o f C e r a m i c P a r t i c l e s A l l atoms and ions on a material's surface have unbalanced chemical bonding so that excess charges (positive or negative) are produced. Zeta potential (t) is the average potential of the surface of shear which forms a sheath enveloping the particles. The charge on the object may be assumed to be located on its surface and the potential falls off with measuring distance away from the surface in accordance with Coulomb's Law. For example, i f the object is a sphere of radius a, the potential ( f ) at a distance r from its surface is given by the following equation A plot of "Pas a function of r for a = 1cm is given in Figure 2.5-1 [103]. It is apparent that near [103]: 4> = Q (2.5-1) 4K£o(a + r) to the surface (r < 10"2 cm) the potential is essentially constant and equal to: 4 / = Q (2.5-2) 1 0 , - 8 - o •4 -2 l o q ) 0 (r/cm) F i g u r e 2 .5-1. Electrostatic potential in the neighborhood of a charged sphere of radius 1 cm. The broken line indicates how the measured potential near a conductor would be affected by the image force i f an electron were used as the test charge [103]. Chapter 2 Literature Review 34 The stability of lyophobic sols is determined by the balance between the repulsive (electrostatic) and attractive (van der Waals) forces. The zeta potential is a very good index of the magnitude of repulsive interaction between particles. When two colloidal particles approach one another so that their double layers overlap, they exert a force on one another. The force/(a) of repulsion per unit area between two approaching particles is given by the following equation [103]: f(a) = -^-B(d%2 (2.5-3) where A is the Hamaker constant, C is a constant that depends on the geometry of the system, and B(d) is a function of separation between the particles and depends on the assumptions made about the conditions under which the approach of the particles occurs. It has been shown that strongly repulsive interparticle potentials can be achieved in aqueous alumina slurries by adjusting the p H (pH=4) and adding an indifferent electrolyte [104]. The short-range repulsive potential, due to added salt, has recently been determined for a sapphire plate using the surface force apparatus by Ducker et al [105]. The short-range interparticle potential is currently believed to be due to a layer of hydrated water similar to that found for mica and clay minerals [106,107]. Short-range repulsion causes the particles within the slurry to be weakly attractive, but non-touching. Such slurries have been termed "coagulated" as opposed to "flocculated" slurries, where the particle network is strongly attractive and touching, due to the presence of only the van der Waals potential [108]. Classic Derjaguin-Landau-Verwey-Overbeek ( D L V O ) theory sums the attractive van der Waals potential and electrostatic potential produced by charged surface sites partially shielded by counterions. It concludes that a strongly attractive particle network can be formed by either changing the p H to the isoelectric point (diminishing the net surface charge density to zero) or by Chapter 2 Literature Review 35 adding sufficient counterions to diminish the electrostatic potential [108]. The zeta potentials of (X-AI2O3 and a -S iC aqueous suspensions as a function of p H are shown in Figure 2.5-2 [109]. 60 1 1 1 1 | I 1 1 1 | 1 1 1 1 • Alumina, AKP53 ; 40 - f j C L • SiC. UF45 -20 — .11 c 1 \ & u 0 \ \ Q_ -20 a . \ 3 \ Q \ \ N -40 • \ \ • V ^ ^ L J -60 1 1 1 1 1 I , 1 1 1 , 1 : 0 5 10 15 pH F i g u r e 2.5-2. Zeta potential of oc-alumina and silicon carbide in aqueous suspension as function of p H [109]. 2.5.2 D i s p e r s i o n o f C e r a m i c P o w d e r s The dispersion processing of ceramic particles in liquid is controlled by interaction forces between particles, including van der Waals attraction, electrostatic repulsion, Brownian motion, kinetic energy of sedimentation, kinetic energy of stirring, and steric forces, as illustrated in Figure 2.5-3 [110]. For large particles (1-10 pm), inertial and gravitational forces become important, and fluid flow can confer large energies to the particles. For colloidal particles, the suspension stability can be increased by the use of a dispersing agent. The chemical dispersants can produce a strong electrostatic repulsive force or reduce the van der Waals attractive potential between interacting particles. Chapter 2 Literature Review 36 Elactroslatle repulsion Buoyancy Figure 2.5-3. Forces acting on and between particles in laminar flow [110] Electrostatic Steric stabilization stabilization -:QT- - : Q T --:a- --GL-Stabilization V * Electrosteric by hydration \ / stabilization lorces , " x X Stabilization •&) ^ by masking Depletion van der Waals forces stabilization Figure 2.5-4. Methods of stabilizing colloidal ceramic particles in liquids [110]. In addition, the particles may be also be stabilized by an adsorbed polymer (steric stabilization) or by adsorption of a strongly hydrophilic fi lm, causing structural hydration forces. Also, the masking of the van der Waals forces which can be achieved by selecting a suitable dispersing medium to match the ceramic properties, and "depletion stabilization" using high Chapter 2 Literature Review 37 concentrations of nonadsorbing polymer, are well established methods of stabilizing ceramic suspensions. These mechanisms of stabilization of ceramic particles are illustrated in Figure 2.5-4 [110]. In conventional ceramic processing, the ceramic powders are dispersed into liquids using wet or dry mil l ing and stirring to break the hard agglomerates [111]. The suspension stability can be increased by using a dispersing agent. In an aqueous environment, this may simply involve changing the solution p H because the zeta potential of ceramic particles changes with p H value. For example, A 1 2 0 3 particles can be dispersed in an alumina alkoxide sol by adjusting the p H to a value of p H = 4 [112]. No degradation of sol/particle suspension was observed within 14 days after the preparation of sol/particle suspension [104]. In other cases, adsorption of specific ions on the particle may occur, causing an increase in the repulsive charge. In both aqueous and non-aqueous solution, chemical dispersants can produce a strong electrostatic repulsive force or reduce the van der Waals attractive potential between interacting particles. In addition, particles may also be stabilized by an adsorbed polymer or by adsorption of strongly hydrophilic ions, causing structural hydration forces. For example, Primin [113] used citric acid as a dispersant for aqueous alumina suspension. The citric acid adsorption on alumina particles resulted in the formation of a negative charge on the alumina particles. The electrostatic repulsive force was increased and the van der Waals attractive potential was reduced, so, the system of particle/liquid suspension was stabilized. Figure 2.5-5a shows that there are only van der Waals attraction and steric repulsion forces, so that the potential energy of interaction is negative (attraction). Figure 2.5-5b shows that the additional strong electrostatic repulsive force overcome the van der Waals attractive force, so the potential energy of interaction is positive (repulsion) [110]. Chapter 2 Literature Review 38 o u a | + o >* O) .2 c • _ o a. Figure 2.5-5. Potential energy diagrams for interparticle energy of van der Waals attraction (VA), steric repulsion (Vs) and electrostatic repulsion (VR) [110]. Cesarano III et al. [114] investigated the processing of highly concentrated aqueous a-AI2O3 suspensions stabilized with polyelectrolytes. The relationship between p H and viscosity for different solids loading was discussed. Figure 2.5-6 shows the effect of p H on viscosity for stabilized suspensions at various solids loading [114]. In 20 v o l % AI2O3 suspensions, all of the stabilized systems have very low viscosity between p H 4.5 and p H 10. The minimum viscosity occurs at p H = 8.8, with viscosity increasing with both increases and decreases in pH. This p H value coincides with zero zeta potential of alumina, which occurs because there is no net adsorption of charge on the surface of ceramic particles [109]. For p H >9.0, the viscosity increase is due to the presence of excess polymer in solution which is not adsorbed on the AI2O3 particles. However, for p H < 8.0, the viscosity increases because the amount of polymethylmethauylate ( P M A A ) absorbed on the surface of alumina particles increases greatly from 0.16 mg/m 2 at pH=10 to 1.4 mg/m 2 at pH=4, while at the same time the polyelectrolyte chains become less dissociated and behave more like neutral polymers [115]. 1 1 •— W v , X \ / / / / ' V / / / / i Chapter 2 Literature Review 39 1.0 0.8 d QL 0.6 g 0.4 o in 0.2 1 1 -5 8 V / 0 \ ^ / -\ 5 0 20 \ ^ i i 4 5 6 7 8 9 10 P H Figure 2.5-6. Viscosity of a - A l 2 0 3 suspensions from 20 to 58 v o l % solid concentrations as a function of p H [51] It has been shown that zirconium hydrogel can be used as an effective processing aid for various ceramic powders [116]. The hydrogel enhances the stability of alumina, titania, and zirconia dispersions, and modifies the zeta potentials of these ceramics. It is believed that the zirconia hydrogel interacts with the dispersed ceramics to form a surface layer on each particle that provides stability. The resulting better dispersion of the suspension improves sinterability and the resulting physical properties of ceramics [117]. Different competitive interaction forces, such as electrostatic repulsion forces, van der Waals attractive forces and steric repulsion forces [118] control dispersion of ceramic particles in liquids. Chapter 2 Literature Review 2.6 Composite Sol-Gel Ceramics 40 2.6.1 Composite Sol -Gel Processing With increasing demand in the quality of advanced ceramics, sol-gel matrix composite materials with ceramic particles or whiskers as reinforcement have received much attention as high performance materials for a wide range of engineering applications [18-20]. The potential advantages of sol-gel processing for ceramic composites are fine scale mixing, low densification temperature and ultimately improved properties. l.s- MIXING OF TWO SOLS [sou] [solll [sou] IcoMPOsrrEl l.b- MIXING OF A SOL AND A SOLIDPHASE • + MH>g. GcbtiOB Diyinf •••• | SOU | ISECONDPHASg Figure 2.6-1. Processing routes for sol-gel-derived composites [119] Chapter 2 Literature Review 41 In general, there are five steps for sol-gel composite processing. (1) Organometallic compounds (such as tetraethylorthosilicate [120], aluminum isopropoxide [121], and zirconia iospropoxide [122] are used to fabricate the sol of composite matrix phase. The sol could be one component or two or more components mixed together to form a homogeneous solution, as illustrated in Figure 2.6-l(a) [119]. The different components may be tailored so that they do not react with each other to form new components. This method allows for good uniformity of the composites. (2) The different solid phases, such as fine powders or fibers, are dispersed into sol before gelation. This leads to a composite with good homogeneity and intimate contact between particles and sol-gel matrix, as illustrated in Figure 2.6-1(b). (3) Impregnation of the fine interconnecting pores by organic or inorganic phases, as illustrated in Figure 2.6-1(c). Surface coating or full impregnation may be achieved. (4) The infiltration or coating of fibers, laminates, or three-dimensional fiber fabric by a low-viscosity sol, as illustrated in Figure 2.6-1(d). (5) The combination of #2 and #3 above to give a "triphasic" composite [119]. 2.6.2 Gelcasting The fabrication of complex-shaped ceramic parts in large quantities has been limited to injection molding, slip casting, or processes requiring extensive machining [123,124]. However, the current injection molding technology for ceramics, has several shortcomings that limit its use as a complex shape forming method. Among these limitations are the following: (1) molding defects; (2) long vehicle-removal times; (3) low green strength after binder removal; (4) warpage during binder removal; (5) differential binder removal; and (6) thick-section problems [123]. The new generic process of gelcasting deals explicitly with the underlying problems of injection molding by the development of a molding process that introduces a new setting mechanism which is generically different from the current practice [125]. This process separates Chapter 2 Literature Review 42 the mold-filling operation from the setting operation, and uses a solution-based vehicle instead of a 100% wax- or polymer-based vehicle. The gelcasting technique is based on a synthesis of ideas from traditional ceramic processing and from polymer chemistry. The kernel of the process is the use of a monomer solution, which can be polymerized to form a strong, cross-linked polymer-solvent gel. The monomer solution provides a low-viscosity vehicle for carrying the ceramic powder. Cross-linking, to form a polymer-solvent gel, provides a mechanism for permanently immobilizing the ceramic slurry to the desired shape after it has been poured into a mold. Because the cross-linked polymer-solvent gelled vehicle contains only 10 to 20 wt% polymers, the solvent can readily be removed from the gelled part by a drying step. Also, because the polymer that is present in the gel is cross-linked, it cannot migrate with the solvent during drying. Figure 2.6-2. Gelcasting processing flowchart [125]. Furthermore, because the setting mechanism is the formation of a gel by polymerization of the monomer in solution, and because this is accomplished after the mold cavity is completely filled, there is a complete separation of the fluid flow operations necessary to f i l l a mold from the Chapter 2 Literature Review 43 setting operation necessary to convert the ceramic slurry to a formed green part [125-128]. A gelcasting process flowchart is shown in Figure 2.6-2. 2.6.3 AUCVSiC Composite Composites of SiC-Ai2C»3 hold great promise for application as structural components and as wear-resistant elements, e.g. cutting tools and forming dies [129-132]. The incorporation of a second phase (SiC) into alumina ceramic matrix can lead to significant increases in the fracture toughness; the level of toughening being strongly dependent on morphology of the second phase (i.e. particles, whiskers, or fibers) [133]. An increase of resistance to fracture initiation is believed to be the mechanism behind the strengthening of ceramics through dispersion of a nanometer-sized (10 - 100 nm) secondary phase. For example, the strength of the nanocomposite with 5 vol% silicon carbide increases from 320 MPa to 1050 MPa, and its fracture toughness increases from 3.2 M P a m 1 / 2 to 4.7 MParn 1 7 2 [134]. Unfortunately, silicon carbide/alumina composites are difficult to sinter to full density because of the covalent nature of the SiC bond. The common densification method for these composites is therefore hot pressing. Numerous studies [135-137] have shown that the addition of sintering additives (mainly Y2O3 , AI2O3, and MgO) to silicon carbide systems improves the densification properties of the materials. Lange [138] has investigated the hot pressing behavior of SiC/A^C^ composites and achieved more than 99% densification at 1950 °C. Boehmite gels have been used as a source for AI2O3 that is either coated on crystalline nanosized SiC precursor, to produce A^CVSiC nanocomposites [139-141]. A 14 wt% boehmite powder and deagglomerated SiC was mixed with distilled water at pH =3.5 for 6 h to form a transparent sol. After drying and calcination, the ultrafine powder was hot-pressed at 1600°C Chapter 2 Literature Review 44 [20]. Compared to nanocomposites fabricated by conventional ball milling, sol-gel processing leads to smaller AI2O3 matrix grain sizes due to the better dispersion of SiC particles. Figure 2.6-3 [46] shows a comparison of strength and fracture toughness of alumina-based nanocomposites at room temperature as a function of SiC content. The strength of the monolithic alumina used as a reference varies from 350 to 560 MPa depending on the study. Following Niihara's original work [142-146], the addition of only 5 vol% nanosized SiC increased the strength to 1050 MPa. A further increase of SiC content lowers the strength to a constant value of approximately 800 MPa. However, Niihara [146,147] explains the decrease of strength value for higher SiC contents as being due to agglomeration problems. 0 10 20 30 SiC content [vol%] Figure 2.6-3. Strength and toughness of Al203-SiC nanocomposites as a function of SiC content (•) by three point bonding test and Vickers indentations; (•) by four-point bonding test; (A) by four-point bend test and indentation-strength method; (V) by three-point bend test and notched beams [46]. Schmid et al. [148] has studied microstructures of Al203-SiC nanocomposites using transmission electron microscopy (TEM). Small SiC particles appear to be either perfect single Chapter 2 Literature Review 45 crystals or show micro-twinning characteristics [149]. The particles with diameters >200 nm usually are more irregularly shaped and are situated exclusively in intergranular positions, as shown in Figure 2.6-4a [149]. The large intergranular SiC particles tend to form clusters in association with open macropores, pockets of an amorphous secondary phase, as was frequently observed in these materials, and hard agglomeration, as illustrated in Figure 2.6-4b [148]. Investigations of SiC/Al203 interfaces, as presented in Figure 2.6-5 [46], revealed that the boundary is free of a secondary phase [150]. Though all S i atom in these grain boundaries may be located within the thin intergranular film (1 nm), measured profiles show a width ~ 5nm, which is to the order of the probe size dominating the convolution product. Profiles of Si concentration across Al 203-mullite phase boundaries, on the other hand, showed a considerably broader maximum = 12nm wide on the mullite side of the interface region, as illustrated in Figure 2.6-6 [148]. The principal features of microstructure which have developed in Al 2 03-SiC composite ceramics are a consequence of the dispersion behavior of S i C particles, and the formation of secondary intergranular phases by interactions between constituent phases and Si0 2 impurities. F i g u r e 2.6-4. S i C particles in intergranular positions: (a) lattice fringe image of 0001 base planes (6H, a-SiC) in medium-sized S i C particle, (b) large S i C particles agglomerated in pores area show rim and residual glass phase [148] Chapter 2 Literature Review Figure 2.6-5. TEM micrograph of an A l 2 C V 5 v o l % S i C nanocomposite showing an interface [46] (b) scan distance (nm) 3 80 40 J AMPS I - • - S i " K a / S 10 15 20 scan distance (nm) Figure 2.6-6. Composition profiles across interface (automated step scans): (a) S i - K a signal grain boundary; (b) S i - K a signal across A ^ C V m u l l i t e - S i C phase boundary [148] Chapter 2 Literature Review 2.7 Composite Sol-Gel Coatings 47 2.7.1 P r e p a r a t i o n o f C S G S l u r r y It is difficult to make crack-free monolayer sol-gel coating films thicker than 10 iim due to the large shrinkage during fi lm densification. Therefore, this restricts the potential applications of these coatings, for wear and corrosion, which usually requires thicknesses greater than 10 urn. There are many piezoelectric applications which also require much thicker films [151,152]. A novel sol-gel based coating process has recently been proposed for fabricating ceramic films with the thickness required to address some of these applications [18]. In this approach, the coating films are ceramic composites formed by dispersing ceramic powders into a sol. The technology has all the benefits of the sol-gel, so it is easy to fabricate the composite coatings on components with complex geometries. A wide range of ceramic composite coatings have been fabricated with thickness of 5 - 200 u\m on a variety of substrate materials and shapes. The coatings were applied to the substrates by dipping, spraying, etc. [18]. The thick coatings may be composed of many thin films by multiple dip coatings and drying [153]. A n analogy describing this process might be to say that the sol-gel makes up the mortar and the ceramic particles make up the bricks in a wall . The fact that coatings do not crack during processing is attributed to two factors. First, the coatings form a network through a sol-gel phase firmly bonded between ceramic particles. Second, because the presence of a significant the amount of ceramic filler decreases the percentage of sol-gel in coatings and less shrinkage occurs [19]. Chapter 2 Literature Review 48 2.7.2 Properties of the Coatings The sol-gel matrix phase of C S G coatings bonds the ceramic particles together to form a strong 3-D network. The interfacial bond between ceramic particles and the sol-gel matrix phase determines the material properties: for example, microhardness and Young's modulus of C S G . The coefficient of thermal expansion of the composite coatings must match that of the substrate as closely as possible in order to prevent the formation of undesirable stresses in the coatings after cooling from fabrication or the heat treatment temperature. The properties of ceramic composite are a complex function of the microstructure and composition. Thermal expansion and thermal conductivity determine the thermal shock resistance of materials [154]. It is observed that different thermal expansion coefficients of inclusions and sol-gel ceramic matrix result in a stress field around the inclusion. This stress builds up during cooling after hot fabrication [23]. The presence of such a stress field affects the fracture behavior of homogeneous brittle materials. Cracks may be formed from small pre-existing defects at or near the inclusion/matrix interface. It is found that the fracture incidence from inclusions increases with increasing number and size of the inclusions and thermal expansion mismatch strain between inclusions and the matrix. In ceramic composite systems, complex non-linear thermal expansion characteristics can be achieved. This allows very close thermal expansion matching between the ceramic coatings and a wide variety of metal substrates, including those that exhibit non-linear behavior due to phase transformations. The coefficient of thermal expansion, a, of ceramic composite coatings is an additive function of the thermal expansions of the various phases present, and can be represented by the following relationship [155]: a = {axKxWx Ipx +a2K2W21p2 + (KiWl I p, + K2W21 p 2 + • • • ) (2.7-1) Chapter 2 Literature Review 49 where ctv, ce2 etc., are the thermal expansion coefficients of the various phases present in the ceramic composite, Kj, K2 etc., are the bulk moduli of these phases, Wj, W2 etc., are the weight fractions, and pu p2 are the densities of the phases. The ability to tailor the thermal expansion characteristics of composite ceramics is a direct consequence of the ability to design the microstructure of the materials, such as composition, reinforced materials, weight fractions, and fabrication technique, etc. 2.7.3 I n t e r f a c i a l B o n d i n g b e t w e e n S u b s t r a t e a n d C o m p o s i t e C o a t i n g s In order to form mechanically strong, adherent and hermetic composite coatings, the formation of chemical bonds at the interface between the composite ceramic coatings and the substrate is very important. There are two kinds of interfacial boundaries [156]: homo-phase boundary and hetero-phase boundary. The homo-phase boundary separates two crystals of the same materials with the same composition. In contrast, the hetero-phase boundary separates two crystals which possess different structures and / or compositions as illustrated in Figure 2.7-1 [157]. Almost all interfacial boundaries between composite coatings and substrates are hetero-phase boundaries. A t the interface between two dissimilar materials, segregation of impurities and formation of reaction products may occur. These processes lead to a modification of the properties of the system. The interface between a metal and a ceramic composite sol-gel coating belongs to the group of hetero-phase boundaries. The free energy per unit area, y, of the interface represents a fundamental thermodynamic property of interface. The free energy is correlated to measurable work of adhesion Wad by the Dupre equation [157]: Wad = ym + yc-r (2-7-2) Chapter 2 Literature Review 50 Where ym is surface energy of the metal and yc is surface energy of the ceramics. Attractive interaction between the two constituents results in Wad > 0. Unfortunately, the Wad is very difficult to measure. Jilavi determined Wad by measuring contact angles at pores present at the metal/ceramic interface. A theoretical description of Wad on the atomic level requires knowledge of interatomic potential at the interface [158]. homophase boundaries grain boundaries twins domain boundaries stacking faults heterophase boundaries M(I) / M(II) ZrO-2 (t) / Zr02 (c) Metal / Ceramic Metal / Semiconductor Metal / Polymer Figure 2.7-1. Homophase boundaries and heterophase boundaries. Homophase boundaries wave on both sides of materials of the same composition and structure, whereas in heterophase boundaries, the structure and/or composition of both components is different [157]. Another aspect of mechanical bonding occurs due to a mechanical keying effect between the roughened substrate and composite coatings. There are several theories describing the mechanism of mechanical bonding. For example, dendrite theory considers the precipitation within ceramic composite coatings of a metallic phase which provides anchor points. The dendrites hold the coating in place by a mechanical keying using anchor points between the metal substrates and the coatings, Figure 2.7-2a [159]. The dendrites are believed to result from the reaction of a metal oxide present in the coatings with a metallic element of the substrate. For Chapter 2 Literature Review 51 example, composite ceramic coatings containing C o O bond to iron substrates by what is believed to be the following reaction [159]: C o O ( ceramic coatings) + (substrate) ^ Co (dendrite) + FeO (ceramic coatings) The bonding of CoO-containing ceramic composite coatings is believed to form through precipitation contact with an iron substrate [48]. The net effect is the dissolution of iron into the ceramic coatings and formation of a pitted surface, thus forming a number of mechanical keys which hold the composite ceramic coatings in the substrate, as shown in Figure 2.7-2b [159]. The reaction is as follows: 2Co (precipitates) + 0 2 (from atmosphere) ~~* 2Co + 2 0 (2.7-3) C o 2 + + 2e" -> Co Fe - 2e" -> F e 2 + (2.7-4) Figure 2.7-2. (a) Schematic diagram of mechanical anchor points: by dendrite theory of composite to metal adhesion, and (b) schematic diagram of mechanical anchor points: by electrolytic theory of composite to metal adhesion [159]. Formation of a chemical bond must occur through a transition zone in which the metallic bonding of the metal is gradually substituted for the ionic-covalent bonding of ceramics [160]. Strong chemical bonding forms at the interface between ceramic and metal substrates i f the ceramic becomes saturated with an oxide of the substrate metal. When the metal oxide is dissolved in the ceramic up to its saturation point, the metal ions wi l l remain at the surface and promote metal-metal bonding across the interface, as illustrated in Figure 2.7-3 [159]. When Si02 is saturated with FeO, strong bonding occurs. A t elevated temperatures, metal ions from Chapter 2 Literature Review 5 2 the ceramic coating wi l l diffuse into the metal where they wi l l gain electrons and become zero valent atoms, similarly, metal atoms wi l l diffuse into the ceramic composite coating and become ionized. A state of dynamic equilibrium wi l l therefore exist at the interface between the composite ceramic coating and the metal substrate [159]. Figure 2.7-3. Simplified 2-D schematic representation SiCvto-metal bonding: (a) ceramic saturated with substrate metal oxide in the interfacial region to give strong chemical bonding via a "mono-layer"; (b) as above, with chemical bonding via a "bulk" oxide layer; the strength of the resulting system is dependent on properties of this bulk oxide layer; (c) interface is not saturated with metal oxide; only weak bonding via van der Waals forces is achieved [159]. Chapter 3 Scope and Objectives 53 C H A P T E R 3 SCOPE AND OBJECTIVES 3.1 Scope of the Investigation Recently, composite sol-gel technology has received much attention for processing high performance materials, such as sol-gel matrix composites and non-permeable composite sol-gel ceramic coatings. However, there are some problems; for example, large densification shrinkage leading to cracks, agglomeration of ceramic particles, and sintering interstresses, that restrict the applications of the composite sol-gel technology. The mechanisms of sintering and dispersion of composite sol-gel systems are not clear. There is no model available to simulate the sintering process of composite sol-gel ceramics and to explain the behavior of these composites. The principle direction of this investigation is to construct a new sintering model of composite sol-gel ceramics and, using the model, to analyze the various experiments to better understand the sintering mechanisms, to predict the sintering behavior, and to optimize the processing of composite sol-gel ceramics. The goal of applied research is to eliminate the above-mentioned problems through development of the novel composite sol-gel technology and to fabricate the nanocomposite sol-gel ceramics and coatings. The models constructed and experimental techniques developed in this investigation are generally applicable to other composite sol-gel ceramics. The dispersion mechanisms of ceramic particles in sols are studied to develop the homogenous, stable suspension system for C S G . Chapter 3 Scope and Objectives 3.2 Objectives 54 T h e g e n e r a l r e s e a r c h o b j e c t i v e is to study and develop an understanding of the mechanisms of dispersion, gelation, drying, interfacial bonding and densification of composite sol-gel ceramics (CSG). This objective is was addressed by the following research tasks: 1. The mechanism of deagglomeration of ceramic particles in alumina sol was investigated by measuring the zeta potential of ceramic particles and particle size distribution in alumina sol. The experimental data were analyzed using D V L O theory. 2. The microstructure of nanocomposite sol-gel ceramics was studied by S E M and T E M . A sintering model for composite sol-gel ( S M C S G ) was proposed and developed based on mass transport through viscous sintering. The parameters of the S M C S G model were obtained using experimental data from AI2O3-AI2O3 C S G system. The S M C S G was used to analyze the sintering stresses and densification rate of the composite and to optimize the experimental procedures. T h e g e n e r a l t e c h n o l o g i c a l o b j e c t i v e is to fabricate high performance C S G ceramics and to produce novel non-permeable, adherent C S G coatings on metallic substrates for high temperature corrosion protection. This objective was approached through execution of the following tasks: 1. Dispersion of ceramic particles (SiC, A 1 2 0 3 and Zr02) in alumina sol precursors. 2. Investigation of the optimum ratio of the sol-gel matrix phase to filler ceramic particles for producing the best performance composite sol-gel ceramics and composite sol-gel coatings. Chapter 3 Scope and Objectives 55 3. Fabrication of S i C / A l 2 0 3 nanocomposites with high content of S i C (up to 50vol%). 4. Fabrication of Z r 0 2 / A l 2 0 3 and AI2O3/AI2O3 nanocomposites with uniform grain size. 5. Development of chemical bonding technology to form non-permeable and high hardness composite sol-gel coatings for wear and corrosion protection. 6. Characterization of the microstructure and properties of nanocomposite sol-gel ceramics and C S G coatings, including microhardness, porosity, gas permeability, and interfacial strength. 7. Post-deposition treatment of ceramic composite sol-gel coatings to close open pores and to improve interfacial bonding between the coating and the substrate. Chapter 4 Experimental Methodology 56 C H A P T E R 4 EXPERIMENTAL METHODOLOGY The experiments were divided into four categories: 1) sol-gel processes, 2) processing of composite sol-gel ceramics (CSG), 3) deposition of composite sol-gel coatings, and 4) characterizations of the ceramics and coatings. The first series of experiments involved preparing alumina sols, measuring their properties, and dispersing ceramic particles in the sols. The properties of alumina sols were varied to control the processing and fabrications of composites and coatings. The second series of experiments were conducted to provide experimental data for the sintering model and to develop the novel, high performance CSG ceramics. The experimental methods included gelcasting, drying and sintering of CSG in the system Al203/Zr02/SiC. The parameters investigated in these experiments included the concentration of alumina sol, the loading content of ceramic particles, gelation time, drying time and sintering time and temperature. The third series of experiments were aimed at fabrication of the non-permeable crack-free composite ceramic coating on metallic substrates. New ceramic sealants were developed for post-deposition treatment of ceramic coatings, such as composite sol-gel ceramic coatings and thermal barrier coatings (TBC). The final series of experiments were performed to assess properties of the composite sol-gel ceramics and coatings. The results of these experiments were used to evaluate the sintering model and fabrication processes and to optimize the experimental parameters. Chapter 4 Experimental Methodology 4.1 Sol-Gel Processes 57 4.1.1 Alumina Sol Aqueous alumina sols were selected as they are stable and have long shelf lives. Aluminum isopropoxide (Al(OC3H7)3, (99.0 wt% purity, Aldr ich Chemical Company, Inc. Milwaukee, U S A ) was used as the starting material for alumina sols, hydrolyzed in excess water. Nitric acid ( I M HNO3) was used to peptize the reaction of hydrolysis of alkoxide [28]. During processing, the solution was stirred constantly and temperature was kept between 90 and 95 °C for 48 hours and cooled to room temperature. The concentration of alumina sol solution studied was from 0.5 to 2 M . Alumina sol is prepared through a hydrolysis reaction of solid aluminum isopropoxide with water. Aluminum isopropoxide particles rapidly react with water to form an aluminum hydroxide surface layer, according to the following reactions, (refer to Figure 4.1-1A and 4.1-1B). A l ( O C 3 H 7 ) 3 + H 2 0 => A l ( O C 3 H 7 ) 2 ( O H ) + HOC3H7 (4.1-1) A l ( O C 3 H 7 ) 2 ( O H ) + H 2 0 => A l ( O C 3 H 7 ) ( O H ) 2 + H O C 3 H 7 (4.1 -2) A l ( O C 3 H 7 ) ( O H ) 2 + H 2 0 => A l ( O H ) 3 + H O C 3 H 7 (4.1-3) The shell of aluminum hydroxide growing on the surface of particles slows down reaction of aluminum isopropoxide with water. In the aqueous aluminum hydroxide system, the following species may exist: A l ( O H ) 2 + , A l ( O H ) 2 1 + , Al (OH) 3 ° , A l ( O H ) 4 ' , A l 1 3 0 4 ( O H ) 2 4 ( H 2 0 ) 1 2 7 + , dimer, and/or trimer. A l l of theses species form rapidly and reversibly [161]. The following reactions were identified in the aqueous aluminum hydroxide system [161]: Chapter 4 Experimental Methodology 58 A l 3 + + H 2 0 <=> A 1 ( 0 H ) 2 + + H + , Q n ^ ^ (4.1-4) [Ar] A l 3 + + 2 H 2 0 <=> A 1 ( 0 H ) 2 + + 2 H + , Gl2 = [A/(0/7) 2 +][// + ] 2 [A/ 3 + ] (4.1-5) 13A1 3 + + 2 8 H 2 0 ^ A l 1 3 0 4 ( O H ) 2 4 7 + + 3 2 H + , g 1 3 3 2 = lAh,0<(OH)^][H+ ]32 ^ J g ) [A/ ] A l 3 + + 3 H 2 0 <=> A1(0H) 3 + 3 H + , + n3 _[A/(0/7) 3][rY + ] ^ ~ [A/ 3 + ] (4.1-7) A l 3 + + 4 H 2 0 <^> A1(0H) 4 " + 4 H + , _[Al(OH)4-][H + ]4 [Ali+] (4.1-8) where Qn is equilibrium constant for a hydrolysis product resulting from reaction with water. H 2 0 Adjust pH=4 Constant Stirring Constant Stirring SAH : shell of aluminum hydroxide Figure 4.1-1. Schematic diagram of the processing of alumina sol: (A) aluminum isopropoxide particles dispersed into water; (B) aluminum isopropoxide rapidly hydrolyzed on surface to form the shell of aluminum hydroxide which slows down the rate of peptization; (C) aluminum hydroxide on surface dissolved into the solution; (D) clear alumina sol. Chapter 4 Experimental Methodology 59 The equilibrium constants for chemical reactions (4.1-4) to (4.1-8) are listed in Table 4.1-1 Table 4.1-1. The equilibrium constants for chemical reactions (4.1-4) to (4.1-8) at 25 °C [161] Equil ibrium Constant Medium Qn 0.0IM N a C 1 0 3 -4.97 Ql2 0.1 M NaC10 3 -9.9 Ql3,32 0.01 M NaC10 3 -100.4 Ql3 0.1 M NaC10 3 -15.6 Qj4 0.0 M N a C l Q 3 -23.0 The concentration of different ions as a function of p H was calculated in the present work and plotted in Figure 4.1-2 according to the data from Table 4.1-1. Note that dimer and trimer are ignored in the previous calculations because the equilibrium constants for dimer and trimer are unknown. 1 M Alumina Sol \ A | 3 + ^ \ / \ [AI 1 3 0 4 (OH) 2 4 (H 2 0) 1 2 ] 7 + / •"r \ \ AIOH 2 + 7 7 * " " - ^ AI(OH)2+ 2 2.5 3 3.5 4 4.5 5 PH Figure 4.1-2. The distribution of hydrolysis products of A l 3 + in 1 M alumina sol as a function of p H . According to Figure 4.1-2, in saturated solution below p H 3.0, only a small fraction (about 1%) of aluminum is hydrolyzed. A n increase in p H leads to removal of H + from the Chapter 4 Experimental Methodology 60 coordinated water molecules, followed by condensation of the O H groups. The species (Ali 304(OH)24(H20)i2 7 +) is the only significant hydrolysis product in saturated solutions above p H 3.3. The p H value in saturated solution determines the hydrolysis products and the reactive direction of alumina sol processing. Therefore, theacid is introduced into a slurry to enhance the rate of peptization. A s hydration of isopropoxide proceeds, aluminum hydroxide continues to react with H + to form the species Al i30 4 (OH )24 (H 2 0) i 2 7 + , and thus the shell of aluminum hydroxide on the surface of aluminum isopropoxide particles becomes thinner. Thus, the rate of water transport through the shell and subsequent reaction with aluminum isopropoxide increases. The acid addition alone does not cause peptization of the slurry. The slurry must be heated above 80°C and held at constant temperature until a clear sol is formed. The rate of peptization drastically drops below 80°C since it is the removal of residual (OR) groups in hydroxide above 80°C which significantly enhances the peptization. If cold water is used in the initial hydrolysis of alkoxide, the slurry must be heated to 85 °C before a substantial conversion to bayerite ((3-Al(OH) 3) takes place [162]. It takes about 24 hours to obtain a clear alumina sol in the final step. The largest particle size of alkoxide, p H value, and temperature dominate the time of sol preparation, the particle size distribution and concentration of alkoxide still affect the rate of peptization. However, N M R study results show that other species, such as dimer, trimer, and/or octamer, may exist in alumina sol [30]. 4.1.2 Properties of A l u m i n a Sol Viscosity is one of the primary factors in determining the sol-gel transition time, stabilization of the slurry of composite sol, gelcasting for processing composite ceramics, and dip and spraying coatings. The viscosity of alumina sol was measured, as a function of the p H for 1 M alumina sol, and as a function of concentration of alumina sol at p H = 4, with a Chapter 4 Experimental Methodology 61 Bookfield Model R V T D V II viscometer (Bookfield Engineering Laboratories, M A , U S A ) , using the A S T M standard C936-81. The electrical conductivity of alumina sol was measured, as a function of the p H of 1 M alumina sols and the concentration of alumina sol at p H = 4, with a M E T R O H M 660 conductometer (Brinkmann Instruments, Switzerland). Ionic strength / of alumina sol affects its zeta potential, the viscosity, the stability, and the double layer thickness. It is difficult to directly calculate the ionic strength of alumina sol from the formula / = — ^ w , Z , . , where M , is concentration of ion i and Z, is charge of ion /, because the exact charge number and concentration of sol clusters are unknown. Normally, it is assumed that solutions with the same electrical conductivity have the same ionic strength. Therefore, the HNO3 was chosen as a reference of the ionic strength of alumina sol because p H of alumina sol was adjusted using HNO3. 4.1.3 D i s p e r s i o n o f C e r a m i c P a r t i c l e s i n A l u m i n a S o l In order to decrease shrinkage and to improve the properties of the composite sol-gel ceramics and coatings, different ceramic powders (SiC, A 1 2 0 3 , Z r 0 2 ) were dispersed into alumina sol solution. A n ultrasonic dispersion was used to investigate the deagglomeration of ceramic particles in alumina sol, using Ultrasonic Disruptor Probe (Horiba). The particle size distributions in alumina sol at p H = 4 and water solution (pH = 4) were measured automatically by Particle Size Distribution Analyzer (Horiba C A P A 700), using the principle of liquid phase photo-sedimentation. Zeta potentials of alumina and S i C were measured using a Zeta-Meter System 3.0+ in 1 M alumina sol (pH=4, and ionic strength / = 0.005 M ) and in water solution of the same ionic strength and p H as that of the sol. Chapter 4 Experimental Methodology 62 4.2 Composite Sol-Gel Ceramics 4.2.1 AI2O3/AI2O3 a n d Z r C V A U C V A h O a C o m p o s i t e s Calcined alumina A-16 (0C-AI2O3, 0.31 (xm, Alcoa Industrial Chemicals, Pittsburgh, P A , U S A ) and yttria stabilized zirconia (Toyo Soda M f g . Co . Ltd. , Tokyo, Japan) were dispersed in I M and 1.5 M alumina sol, respectively. The concentration of the ceramic powders was from 4 to 24 vo l% of the sol solution. The slurry was mixed using an ultrasonic mixer for 10 min or ball mill ing for 24 hours. The viscosity of the composite sol-gel slurry was adjusted by changing pH. A water-soluble precursor of magnesia (Mg(N03)2) was added to the alumina sol to control sintering of the composite sol-gel ceramics. The concentration of magnesia in the alumina sol was varied from 0 to 2 mole%. The slurry was cast into plastic mold for gelation for 24 hours. The hardened bodies were then removed from the mold and dried at room temperature and humidity 60-80% for one week. The green bodies were calcined at 550° (aluminum hydroxide transforms to y -A l20 3 ) , and then sintered at 100-1500°C. 4.2.2 S i C / A l 2 0 3 / A l 2 0 3 S o l - G e l C o m p o s i t e s In order to produce the composite sol-gel specimens (CSG), calcined alumina A-16 (0.3-0.5 ixm, Alcoa Industrial Chemicals, Pittsburgh, P A , U S A ) and S i C powders (0.5-3 u\m, Norton Company, Worcester, Massachusetts) were dispersed in alumina sol followed by ultrasonic mixing for 10 min. The content of S i C in the composites was varied from 5 to 60 vol%. The slurry viscosity was controlled by adjusting p H between 3 and 4.5. The stable slurry was cast into plastic mold. After 24 hour gelation, the samples were removed from the mold and Chapter 4 Experimental Methodology 63 dried at room temperature and humidity 60-80% for one week. The alumina phase contained about 86 v o l % calcined alumina, and 14% alumina from decomposition of the sol-gel phase. This resulted, for example, in C S G composition of 50vol% S i C , 43 v o l % alumina from calcined phase and 7 v o l % alumina from the sol phase, for the specimen containing 50 v o l % S i C . For samples containing less S i C , the content of alumina from sol phase increased proportionally. In parallel, samples resulting in identical ratio (vol%) of SiC/Al203 were dispersed in de-ionized water with p H = 3 - 4.5, adjusted using nitric acid, followed by similar shaping and drying procedure. These samples are coded C W D , for composite water-dispersed (as opposed to the composite sol-gel, C S G ) . This procedure resulted, for example, in 50vol% S i C and 50 v o l % alumina from calcined phase. A l l the samples were sintered using tungsten heating element furnace (Centorr Associates Inc., Sancook, N .H. ) in argon atmosphere, at 1600 to 1950 °C for 1 hour. 4.3 Composite Sol-Gel Ceramic Coatings 4.3.1 Coating Processing Technology The preparation of the composite sol-gel was the same as described in Section 4.2. However, different methods were investigated to fabricate the thick ceramic coating films. In dip coating method, the substrate was lowered into a slurry, and then withdrawn from the vessel at constant speed. The coating thickness was a function of withdrawal speed (1-20 cm/min) and slurry viscosity (2-30 mPa-s). The thickness of coating was measured using a P O S I T E C T O R 6000 coating thickness gauge (DeFelsko Corporation, N Y , U S A ) . Alternatively, the coatings were applied by a spray gun. The composite sol-gel coatings were dried at room temperature for 30 min, at 100°C for 10 min, and then were sintered at 450-600°C for 10 min. In order to reduce Chapter 4 Experimental Methodology 64 thermal stress build up during cooling, the thermal expansion coefficient of the C S G coatings was graded to match that of the substrates using alumina ( a = 6.0 x 10~ 6K _ 1) and zircona ( a = 10.23 x lO^K" 1 ) powders. The surfaces of the substrates were roughened by sand blasting, to provide mechanical bonding at the interface. The surfaces of the substrates were also oxidized for 30 min at 450-550 °C to form a metal oxide fi lm which aided the formation of a chemical bonding at the interface. In some experiments, the surfaces were also treated by phosphoric acid to form a iron phosphate fi lm at interface. 4.3.2 E l e c t r o p h o r e t i c D e p o s i t i o n ( E P D ) E P D coatings were applied by imposing a constant D C voltage (VT1187 Coulometric Power Supply) across the electrode surface, resulting in the migration of positively charged colloids and ceramic particles toward the cathode. A n electric field ranged from 1 V / c m to 2 V/cm, for coating time 2 to 30 min. Once a coating was applied, the sample was withdrawn from the suspension at a rate of 10 cm/s. Coating thickness was controlled by varying the voltage and concentration of the slurry. 4.4 Characterization of CSG Ceramics 4.4.1 M i c r o s t r u c t u r e The Scanning Electron Microscopy (SEM) with E D X analysis probe and backscattered electron (BSE) mode (Hitachi-2300) was used to examine the particle size distributions and to evaluate the microstructure of composite sol-gel ceramics and coatings. The powder samples for S E M were prepared by dispersing the powder into water or alcohol, and then the suspensions Chapter 4 Experimental Methodology 65 were dropped on the glass pellets, dried at room temperature followed by depositing a carbon fi lm on them. To observe the cross sections of coatings, the samples were infiltrated and mounted using epoxy resin, and then cut by slow speed diamond saw followed by polishing to 1 pm finish. To observe the microstructure of the composite sol-gel ceramics, the samples were polished to 1 pm finish, followed by thermal etching for 20 min at 1400°C. The E D X and X-ray were used to analyze the phase composition and the distribution of elements, respectively. Transmission Electron Microscopy ( T E M ) with E D X analysis probe (Hitachi 800, Tokyo, Japan) was used to investigate the microstructure and composition of the SiC-Al 20"3 nanocomposites. The samples were first cut into small pallets with approximately 200 pm thick, and then thinned up to 100 pm, followed by dimple grinding/polishing on the center of the samples to 10 pm thick. Finally, the samples were thinned by ion-beam mill ing (5 keV A r ions, 20° angle of incidence for 15 hours) until there was a small hole the center, and then coated with thin carbon fi lm. 4.4.2 Properties The open porosity of C S G was measured by infiltrating the samples with deionized water at 25 °C under 10"3 Torr vacuum for 6 hours. Before infiltration the samples were weighed at air Wj, after infiltration samples were weighed at air W2, and then infiltrated samples were weighed at 25 °C deionized water W3. The density p c of composites can be expressed as a function of density of water pw\ Pc = (4.4-1) w The open porosity FQ of C S G is given by: Chapter 4 Experimental Methodology 66 w 2 - w ; W2-W3 (4.4-2) The total porosity FT of C S G is given by: P f - P c Pf (4.4-3) where p / i s the density of fully dense composite. The microhardness tests were performed using B U E H L E R MICROMET®3 Hardness Tester at loads of 0.5 and 1 kg for 10 seconds. The samples were polished to one micrometer finish and indentation cracks were examined by S E M . The open porosity of C S G was measured by impregnating the samples in deionized water at 25°C and 1 0 3 Torr vacuum for 6 hours. The bonding strength of the ceramic coatings to substrate was tested according to the procedure of A S T M standard ( A S T M C-633-79). The coatings were deposited on 1 inch diameter rod substrate, and a thermoset epoxy ( 3 M type 1386) was used to glue the another same rod substrate on the surface of coating. The samples were cured at 180 °C for 8 hrs. A n F N S T R O N universal testing machine with 10,000 lb load cell was used in sample tension to measure adhesion of coatings. The A S T M standard ( A S T M C-577-78) test method for gas permeability of refractories was extended to study porous coatings. The 2-3 mm thick films of C S G were prepared for the tests as shown in Figure 4.4-1 [38]. The specimen was placed between two rubber gaskets, enclosed between the two halves of an airtight holder. Dry nitrogen gas was passed through the coatings under a fixed pressure difference across the specimen. The coating permeability was calculated according to the following equation [38]: xlO (4.4-4) A rAP Chapter 4 Experimental Methodology 67 where Kp is the permeability, expressed in centidarcys, 77 (mPa-s) is the viscosity of the gas, Qk (cm3/s) is the flow rate of the gas, L (cm) is the sample thickness, Ar (cm2) is the area of the sample, and AP (atm) is absolute pressure drop across the sample. One centidarcy is defined as a flow of 0.01 cm3/s of a fluid of 1 mPa-s viscosity through a 1 cm cube under a pressure difference of 1 atmosphere. Figure 4.4-1 Cross-section of the permeability jig. Chapter 5 Experimental Results and Discussion 68 C H A P T E R 5 EXPERIMENTAL RESULTS AND DISCUSSION 5.1 Composite Sol-Gel Processes 5.1.1.1 Physical Properties of Alumina Sol Viscosity of alumina sols as a function of concentration is plotted in Figure 5.1-1 for p H = 4. The viscosity increases slowly with increasing concentration of alumina sol for concentrations below 2.8 M . However, the viscosity increases rapidly when the concentration is over 3 M . 60 50 A pH=4 10 A 0 -9r 0 2 3 4 Concentration of Alumina Sol, (M) Figure 5.1-1. The viscosity of alumina sol at p H = 4 as a function of the concentration The van der Waals intermolecular attractive energy is given by [110]: Chapter 5 Experimental Results and Discussion 69 V =-F - ^ L (5 1-1) f6H0 P ' where H0 is the separation distance of interparticles, Ff is the form factor Ff = — — — , a\ and a 2 a, + a2 are radius of particles, and A123 is the net Hamaker London function for materials 1 and 2 in medium 3. In the case of monodispersed particles, a\ = a2, and A123 becomes A121 given by [163]: A21 = A12 ~ A i + A2 — 2 A ] 2 ~ (V A1 — VA2 ^ (5-1-2) where A n and A 2 2 are the Hamaker constants of materials in vacuum. If it is assumed that the sol particles are uniform and homogeneously dispersed in the medium, the center distance Ds between two particles can be given by: 10"8 D , = - _ ( n m ) (5.1-3) where c is molar concentration of alumina sols and NA is Avogadro's constant. Thus the interparticle separation distance Ha can be obtained: H0=Ds-ax-a2 (5.1-4) The interparticle separation distance H0 of sol clusters as a function of concentration c of alumina sol is plotted in Figure 5.1-2. The separation distance between particles decreases with increasing concentration of alumina sol at p H = 4. The van der Waals attraction energy, VA as a function of the separation distance of interparticles, Ha was calculated using Eq.(5.1-1) and plotted in Figure 5.1-3, where A121 = 4.0x10 2 0 J [110]. The energy of attraction between sol clusters increases from 0.64x10" 2 0 J to 8x l0* 2 0 J for an increase in concentration of alumina sol from 1 M to 3 M . A s an increase of concentration of alumina sol above 2 M causes a reduction of the separation distance of interparticle to below 1 nm, the van der Waals attraction force wi l l dominate the system. Chapter 5 Experimental Results and Discussion 70 7 0 0.5 1 1.5 2 2.5 3 3.5 Concentration of Alumina Sol, (M) Figure 5.1-2. The interparticle separation distance of alumina sol cluster as a function of concentration of alumina sol at p H = 4. 0 0 0.5 1 1.5 2 2.5 3 H0, (nm) Figure 5.1-3. The van der Waals attraction energy of alumina sol clusters as a function of interaction distance in alumina sol at p H = 4. Viscosity of the alumina sols as a function of p H is shown in Figure 5.1-4. The viscosity is approximately equal to that of water for p H below 5.3. When the p H exceeds 5.6, the viscosity rapidly increases from 1.2 cP to more than 30 cP. This is because a rearrangement of the sol Chapter 5 Experimental Results and Discussion 71 cluster structure takes place when gelation begins. This causes an increase in the polycondensation reaction rate, and gelation time becomes very short ( l-5min). 4 PH Figure 5.1-4. The viscosity of 1 M alumina sol as a function of p H . 5.1.1.2 Electrical Conductivity of Alumina Sol The electrical conductivity of 1 M alumina sol as a function of p H is shown in Figure 5.1-5. The conductivity of alumina sols decreases sharply above pH=2 because the simple species in solution, such as A l ( O H ) 2 + , A l ( O H ) 2 + Ali304(OH)24(H20)2 7 + , dimer and/or trimer, condense to large polyvalent clusters. The electrical conductivity of alumina sol at p H = 4 as a function of concentration is plotted in Figure 5.1-6, the viscosity of which indicates that the molar conductivity of alumina sol is independent of concentration. Chapter 5 Experimental Results and Discussion 72 Figure 5.1-6. The conductivity of alumina sol at p H = 4 as a function of concentration Under the influence of an applied electric field, the random motion of clusters wi l l be sufficiently perturbed to produce a small component of acceleration in the direction of the field. The electrical conductivity of an ionic solution is given by the equation [164]: Chapter 5 Experimental Results and Discussion 7 3 Ke=J^i[Fciui (5.1-5) i where w, is the electrical mobility of ion i, c, is concentration of ion i, Zt is the valency of ion i, and F is the Faraday constant. The mobility of ions can be estimated using Stokes' formula [165]. u=-^- (5.1-6) 6707a where the e is the charge of electron, 17 is viscosity of the solution, and a is the hydrodynamic radius. In comparison with concentrations of NO3" and alumina sol clusters, the H + and OH" concentrations are negligible in a solution with a p H between 3.3 and 5. Therefore, it can be assumed that there are only two types of species in the solution, NO3" and alumina sol clusters, structure and charge of which are unknown. The condition of electroneutrality is ^ c , Z , = 0 , i and thus, the Eq . (5.1-5) reduces to: Ke =Z+Fc+(u++u_) or Ke =Z_Fc_(u+ +u_) (5.1-7) The mobility of N 0 3 " (a = 1.89A) was calculated using Eq . (5.1-6), and the radius a of alumina sol cluster normalized by the charge Z was estimated by using Equations (5.1-7) and (5.1-6). The — of alumina sol cluster as a function of the p H of 1 M alumina sol is plotted in Figure 5.1-7. When the p H of alumina sol is below 2, the ratio — is less than 1.5A, indicating that there may exist only simple species, such as A l 3 + , A l (OH) 2 " , and/or Al(OH)2~ in the solution. Increase of — with increasing p H indicates that sample hydrolysis products start to polymerize to form larger polyvalent ions which are called here sol clusters. The size of alumina sol clusters continues to increase with increasing p H . If there are 7 positive charges per particle, as for Chapter 5 Experimental Results and Discussion Ali304(OH) 24(H 20)i2 7 + , the particle size is about 3.5 nm at p H = 4. The size of sol clusters wi l l increase rapidly for p H > 5 due to polycondensation. 74 Figure 5.1-7. The normalized radius of alumina sol cluster in 1 M alumina sol, as a function of p H According to conventional Debye-Huckel theory, the thickness of the electrostatic diffuse double layer thickness (K ) can be expressed as [161]: 2e2I 1/2 (5.1-8) where / = — ^ n , Z , 2 is ionic strength of solution, e is the electronic charge, n, the number of ions per cm 3 , Z, is valency of the ionic species, e0 is the dielectric permittivity of a vacuum, and er is the relative permittivity. Therefore, the thickness of the double layer increases with decreasing ionic strength. This causes a reduction of the repulsion force in alumina sols and an increase of van der Waals attraction force, and thus ion species easily condense to large clusters. Chapter 5 Experimental Results and Discussion 75 5.1.1.3 Zeta Potential of Ceramic Particles in Alumina Sol Atoms and ions at surfaces have unbalanced chemical bonding, producing excess charges (positive or negative) at surface of materials. When a charged particle (including a closely bound solvent layer) moves relative to the surrounding liquid, the zeta potential is defined as the potential difference between the slip plane and the bulk. The stability of a lyophobic sol is determined by the balance between the repulsive interaction and attractive forces. Zeta potential is very good index of the magnitude of repulsive interaction between particles. When two colloidal particles approach one another so that their double layers overlap, they exert a force on one another. The zeta potential of oc-alumina (A 16) (Alcoa Industrial Chemicals, Pittsburgh, P A , U S A ) and S i C (Norton Company, Worcester, Massachusetts) particle suspensions in water and 1 M alumina sol are shown in Figure 5.1-8 and 5.1-9, respectively, as a function of pH. The zeta potential of A16 alumina in 1 M alumina sol increased marginally as compared to that of alumina in water. N o other significant changes were recorded between the two solutions. The ionic strengths of both water solution and alumina sols were 0.005 M . The measurement of zeta potential of alumina in sols was limited to p H < 5. Increasing the p H of I M alumina sol f rom 5.5 to 5.9 causes a rapid increase in the sol viscosity from 1.2 cP to more than 10 3 cP due to polycondensation reactions of the type = A l - 0 - H + H - 0 - A l = - » = A l - 0 - A l = + H 2 0 . This effectively prevents the zeta potential measurements in the sol for p H > 5.5. In contrast to alumina, a dramatic change of the zeta potential function was observed for water and sol dispersed silicon carbide, as shown Figure 5.1-9. The sign of zeta potential of S i C in alumina sol is reversed from negative to positive, for 3.5 < p H <5.5. Alumina sol consists of a homogeneous medium and colloidal particles dispersed therein. The size of alumina sol particles is dominated by the p H value of the sol. The sol particles have Chapter 5 Experimental Results and Discussion 76 more than 7 positive charges on the surface and the size of sol particles is in a range of 3.5 -50 nm at p H = 4. Figure 5.1-8. Zeta potential of alumina particles dispersed in 1 M alumina sol and in water solution, as a function of pH. > E c o Q. 0> N 50 40 30 20 10 0 -10 -20 -30 -40 -50 SiC Dispersed in 1 M Alumina Sol SiC Dispersed 1 in Water P • • — • 10 PH Figure 5.1-9. Zeta potentials of S i C particles dispersed in 1 M alumina sol and in water solution, as a function of p H . If 8 v o l % alumina powder (assuming a 0.3 pm diameter uniform sphere) is homogeneously dispersed into 1 M alumina sol at pH=4, the average center distance between the particles is approximately 0.5 pm and the average separation distance between the particles is Chapter 5 Experimental Results and Discussion 77 approximately 0.2 u\m. The atoms with unbalanced charges on the surface of alumina particles hydrolyze with OH". It is believed that the sol clusters interact with hydrolyzed AI2O3 particles by a polycondensation reaction, as shown in 5.1-10a. Effectively, the sol clusters coat the surface of ceramic particle. However, as the surface structure of sol clusters is similar to that of the surface of aluminum hydroxide, the zeta potential of this hybrid particle remains unchanged. The interaction of S i C with the sol clusters is very different from alumina, as shown in Figure 5.1-10. The surface of S i C is normally oxidized to form an ~1 nm S i02 f i lm, which, in a humid environment, is further hydrolyzed to form a negatively charged =Si(OH)~. When such S i C particles are dispersed in 1 M alumina sols, they adsorb the positively charged sol clusters to produce a coating layer, as schematically illustrated in 5.1-10b. The coordinate number of A l 3 + in sol clusters is 6. The following chemical reaction may occur at the interface between the particle and sol cluster, leading to a reversal of surface charge: s S i ( O H ) - + A104Al 1 2 (OH )24 (H 2 0) ,2 7 + -> = S i - 0 - A l - A 1 0 4 A l i i ( O H ) 2 3 ( H 2 0 ) , 2 6 + + H 2 0 (5.1-9) As a result of these interactions, a thick double layer forms on the surface of S i C particles with alumina sol clusters that are similar in structure to aluminum hydroxide. Therefore, the zeta potential of S i C in alumina sol is similar to that of alumina. Additionally, the adsorbed sol clusters aid dispersion of S i C through steric interactions. Effectively, the electrostatic repulsion and steric interaction forces dominate the dispersion process, leading to deagglomeration of S i C particles, as detailed in the following section. Chapter 5 Experimental Results and Discussion Double Layer Oxidized Layer Si0 2 Thickness < 1 nm Al 2 0 3 Particle <D = 0.1-0.5nm •* Alumina Sol Clusters <D = 4-35 nm + if \ ^Al-O-H HO J 2O a 0-AI= •AJ-O-H T H-O / / Al- O-H +  H-Q A/- o -oV | % SiC Particle <& = 3-5u.m U Double Layer Alumina Sol Clusters O = 4-35 nm Al-O-H a Figure 5.1-10. Schematic diagram of ceramic particles interacting with the alumina sol clusters; (a) hydrolyzed alumina particles react with alumina sol clusters; (b) hydrolyzed S i C particles with adsorbed alumina sol clusters. Chapter 5 Experimental Results and Discussion 79 5.1.2 D i s p e r s i o n o f C e r a m i c P a r t i c l e s i n A l u m i n a S o l The homogenization, dispersion, and stability of ceramic particles in liquids are important issues in the processing of high-performance ceramics produced by slip casting, gelcasting, and novel composite sol-gel method. Recent studies [110,165] show that the distribution and packing of ceramic particles throughout the green body controls the porosity and microstructure and plays an important role in determining the reliability of final products. The aim of the dispersion process is to achieve a homogeneous suspension with high solid content, free from agglomerates, and with well-defined rheological behavior. 5.1.2.1 Electrostatic Stability of Ceramic Particles in Alumina Sol Dispersion of ceramic particles in liquid is controlled by interaction forces between particles, including van der Waals attraction, electrostatic repulsion, Brownian motion, kinetic energy of sedimentation, kinetic energy of stirring, and steric forces. Normally, van der Waals attraction, electrostatic repulsion, and steric forces dominate the dispersion processes. The D L V O theory [110] provides a quantitative explanation of the coagulation by equating the interaction potental equation W term as the summation of the dispersion attraction and the electrostatic repulsion VR: To calculate VR we need to consider the interaction between approaching double layers as the particles come together. The complete mathematical expression for VR is given by [110]: (5.1-10) l + expC-KfVj^ l - e x p ( - K » J > + ( V A 12 +! / 22 ) ln( l -exp ( -2KfO) (5.1-11) Chapter 5 Experimental Results and Discussion 80 „ a\a7 where F, =——— ax + a2 is the form factor or size for two particles with radius ai and a2> £ is dielectric constant; y/i and y/2 are the surface potentials of the particles; K is the reciprocal width of the electrical double layer, which is related to the ionic strength of the medium; and Ha is the interparticle separation distance. As I/A values are difficult to measure, it is common to substitute them with zeta potential values. This assumption is a good approximation when the absolute zeta potential is less than 50 m V . In the case of VA, the van der Waals interaction energy is given by [110]: VA = - F x -*123 6H„ (5.1-12) where A is the Hamaker constant, as discussed in section 5.1.1.1. The van der Waals attraction energy between S i C - S i C (a = 0.8 pm, A123 = 0.6xl0" 2 0 J) and alumina-alumina (a = 0.155 pm, A123 = 4xlO" 2 0 J) versus the interparticle surface distances is plotted in Figure 5.1-11. ~ -2 x < > -4 -6 pH = 4 — Alumina SiC-SiC 20 40 60 80 100 H 0 , (nm) Figure 5.1-11. The van der Waals attraction energy of S i C - S i C and AI2O3-AI2O3 particles in aqueous solution as a function of interparticle separation distance (Eq5.1-12). Chapter 5 Experimental Results and Discussion 81 5.1.2.2 Steric Stability of Ceramic Particles in Alumina Sol Steric stability allows for polymer chains, large ions, or clusters (in some cases) to attach to the surface of ceramic particles by chemical bonding, electric adsorption, or hydroxide bonds. Such a coating layer on the ceramic surface reduces the van der Waals attraction energy and increases stability of the slurry, contributing to steric stabilization of the suspension. Alumina Sol Clusters + - V " + + +:.+,4. + + + + + SiC : + > ~ — + + + + + + + • + + / • ' + + + V + • + " + ' • -- + + -V+ SiC :!+ + K + + + Figure 5.1-12. Schematic diagram of the two approaching ceramic particles with a f i lm thickness of adsorbed sol clusters of Ha 12 In C S G , the ceramic particles (SiC and alumina) absorb the alumina sol clusters on their surface to form a steric layer, as illustrated in Figure 5.1-12. A s a result, the short-range steric repulsive term Vs rises rapidly at short interaction particle distance (Ha< HP0 ) and the dispersion is stabilized. According to the previous analysis, the size of alumina sol clusters at pH = 4 is in the range of 3.5-35 nm. Thus, the minimum separation distance H0P of ceramic particles dispersed in alumina sol is larger than 7 nm because the surfaces of ceramic particles are coated by at least a single layer sol cluster with a minimum size of 3.5 nm. Chapter 5 Experimental Results and Discussion 82 5.1.2.3 Interaction Energy of Ceramic Particles in Alumina Sol The electrostatic repulsive interaction energy VR of a-alumina and S i C in water and alumina sol as a function of separation distance is estimated from Eq.(5.1-11) and plotted in Figures 5.1-13 and 5.1-14, respectively. In order to avoid the influence of ionic strength on thickness of double layer of ceramic particles in solution, the ionic strength of alumina sol and water solution is 0.005 M . The particles of S i C (1.6 pm) and AI2O3 (0.31 pm) are assumed to be equal-sized spheres. In comparison with alumina dispersed in water solution, the repulsive energy in alumina sol was increased slightly. The repulsive energy of S i C dispersed in alumina sol significantly increases because of the large increase of the zeta potential of S i C in alumina sol. 5 Alumina/Water — Alumina/Sol 0 0 5 10 15 20 25 30 H 0 , (nm) Figure 5.1-13. Electrostatic repulsive energy of alumina particles dispersed in 1 M alumina sol and water solution as a function of the interparticle separation distance. The average diameter of alumina particles is 0.31 pm and ionic strength of alumina sol is 0.005 M . Chapter 5 Experimental Results and Discussion 83 H 0 , (nm) Figure 5.1-14. Electrostatic repulsive energy of S i C particles dispersed in 1 M alumina sol and water solution as a function of the interparticle separation distance. The average diameter of alumina particles is 1.6 pm and ionic strength of alumina sol is 0.005 M . The interaction energy of AI2O3-AI2O3 and S i C - S i C in alumina sol and in water as a function of interparticle separation is calculated and plotted in Figures 5.1-15 and 5.1-16, respectively. The interaction energy of alumina-alumina dispersed in sol is only slightly higher than that in water because electrostatic repulsive forces increase only slightly in alumina sol. The interaction energy of S i C - S i C in alumina sol is significantly increased (as compared to water) because of large electrostatic repulsive energy. The Brownian motion energy (~kT) of ceramic particles in alumina sol was not included during calculations of the interaction energy in present study. The results of calculations indicate that electrostatic repulsive energy dominates the dispersion processing of S i C particles in alumina sol. Chapter 5 Experimental Results and Discussion 84 H 0, (nm) Figure 5.1-15. Interaction energy of alumina particles dispersed in 1 M alumina sol and water solution as a function of the interparticle separation distance 4 H 0 , (nm) Figure 5.1-16. Interaction energy of S i C particles dispersed in 1 M alumina sol and water solution as a function of the interparticle separation distance 5.1.2.4 Dispersion of Ceramic Particles in Alumina Sol: Experimental Verification Figure 5.1-17 shows the particle size distribution for A16 alumina dispersed in water, and alumina sol. For the lowest particle size intervals (0.2-0.3 | i m and 0.3-0.4 u\m), the sol-dispersed fraction is larger compared to the water-dispersed fraction. The mode particle size in alumina sol (-0.27 \sm) is smaller than that in water (-0.36 u\m). The 0.6-1 u\m agglomerates are largely Chapter 5 Experimental Results and Discussion 85 removed upon dispersion of the alumina powder in the alumina sol. This result is confirmed by taking high magnification S E M micrographs of alumina A 1 6 powder dispersed in alumina sol, Figure 5.1-18a, and in water, Figure 5.1-18b. Figure 5.1-19 presents the respective results for the nano-alumina (NA) (V-AI2O3 Nanophase Tech. Corp., Burr Ridge, IL) powder, dispersed in the two media. The sol-dispersion effect is even more dramatic. Although there is only a negligible amount of N A particles smaller than 0.1 um when dispersed in water, majority of N A particles are below 0.1 pjm when dispersed in the sol. The powder is effectively deagglomerated in the sol, and contains less than 3% particles larger than 1 p:m. When dispersed in water (pH=4), more than 35 % of N A alumina agglomerates into particles larger than 1 pirn. Figure 5.1-20 shows that all S i C agglomerates larger than 3 um are removed from alumina sol. The water suspension contains particles larger than 2 u\m. The experimental results confirm that dispersion processes of ceramic particles in sol are more effective than in water. These experimental data are in good agreement with that of theoretical predictions from D V L O theory. 40 35 H g 30 £ 25 c 0) 3 20 ty 4) £ 15 10 • Alumina (A16) Dispersed in Water at pH=4 Alumina (A16) Dispersed in Alumina Sol at pH=4 <0.1 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1< Particle Size, (nm) Figure 5.1-17. The particle size distribution of alumina A16 dispersed in 1 M alumina sol and water solution at p H = 4 Chapter 5 Experimental Results and Discussion 86 Figure 5.1-18. The S E M morphology of A16 dispersed in 1 M alumina sol (a) and water solution (b) at p H = 4. 40 35 30 ^ 2 5 > o | 20 I 15 10 • Alumina (NA) Dispersed in Water at pH=4 ^Alumina (NA) Dispersed in Alumina Sol at pH=4 <0.1 0.1 0.3 0.4 0.5 0.6 0.7 Particle Size, (pm) 0.8 0.9 1< Figure 5.1-19. The particle size distribution of nano-alumina (NA) dispersed in 1 M alumina sol and water solution at p H = 4. Chapter 5 Experimental Results and Discussion 87 80 -i 70 -60 -50 -SK o 40 -c 0) 3 cr Q> 30 -ul 20 -10 -0 -HSiC Dispersed in Alumina Sol at pH=4 • SiC Dispersed in Water at pH=4 <1.0 1 2 3 Particle Size, (urn) Figure 5.1-20. The particle size distribution of S i C dispersed in 1 M alumina sol and water solution at p H = 4 5.1.3 Summary 1. It is believed that the p H value and temperature of solution dominate the time of alumina sol preparation. 2. Below 2.8 M and at p H =4, the viscosity increases slowly with increasing concentration of alumina sol. However, the viscosity increases rapidly when the concentration is over 3 M . This is because the reduction of separation distance of interparticles results in an increase in attraction energy between sol clusters from 0.64x10" 2 0 to 8x lO" 2 0 J . 3. The viscosity of alumina sol dispersed by ultrasonic mixer and then immediately measured is approximately equal to that of water for p H below 5.3. When the p H exceeds 5.6, the viscosity rapidly increases from 1.2 cP to more than 30 cP. This is because a major rearrangement of the sol cluster structure occurred due to polycondensation when gelation begins. Chapter 5 Experimental Results and Discussion 88 4. The conductivity of 1 M alumina sol decreases sharply above p H = 2 because the simple species, such as Al (OH) 2 " , Al (OH) 2 " , A l i30 4 (OH )24 (H 2 0)i2 7 + , dimer, and/or trimer in solution may condense to large polyvalent ions. 5. The zeta potential of AI2O3 in 1 M alumina sol increases slightly, as compared to water; however, the zeta potential of S i C reverses its sign over a wide range of p H values from negative to positive because the structure of S i C particles has been changed by coating of a layer of alumina sol clusters. 6. The stability of ceramic particles in alumina sol and water was analyzed and the interaction energy of system was estimated using D L V O theory. Dispersion of alumina and S i C particles has been substantially improved in alumina sol, as compared to pure water of similar acidity, i.e. the average agglomerate size is decreased by at least 50%. The particle size of alumina dispersed in water is larger than that dispersed in the alumina sol because of the strong steric repulsive interaction. The size of S i C particles dispersed in water is larger than that dispersed in the alumina sol, because the negatively charged S i C agglomerates adsorb oppositely charged alumina sol clusters on their surface, and produce strong electrostatic and steric repulsive forces. Based on these findings, it is expected that the alumina sol clusters can be used as an efficient, clean dispersant for single-phase and composite ceramics. Chapter 5 Experimental Results and Discussion 5.2 Alumina/Alumina and Alumina/Zirconia CSG Ceramics Sol-gel matrix composites reinforced with ceramic particles or whiskers have received much attention as high performance materials, for a wide range of engineering applications. The potential advantages of sol-gel processing for ceramic composites are fine scale mixing and low densification temperature, leading ultimately to improved properties. A variety of solid phases, such as fine powders or fibers, can be dispersed into a sol before gelation, leading to a composite with good homogeneity and intimate contact between the components. The composite slurry is typically dried at 25-250°C, and sintered at temperatures several hundred degrees lower than the corresponding calcined ceramics. The main objective for this section was to develop the composites in alumina/alumina and alumina/zirconia composite sol-gel systems sintered to full density below 1450°C. 5.2.1 T h e r m o g r a v i m e t r i c A n a l y s i s o f C S G C e r a m i c s Weight loss is one of the important parameters used to monitor structural changes and phase transformations during drying and sintering. During the drying stage, removal of solvent condenses the alumina gel towards a continuous matrix, providing mechanical strength to a pre-sintered body. When the composite green bodies were dried at 300 °C, almost all excess solvents escaped through interconnected porosity, accompanied by approximately 60% weight loss and 25% volume reduction. Figure 5.2-1 illustrates the typical weight loss of the C S G with 86.2 v o l % calcined alumina and of pure sol-gel phase as a function of temperature. The weight loss of pure alumina gel is approximately 10 wt% during evaporation of unstructured water and organic compounds from 100°C to 500 °C, but it is only 2 wt% for C S G , as shown in Region I in Figure 5.2-1. Heat treatment between 500 °C and 550 °C causes rapid dehydration of the gel as A l O O H transforms to V-AI2O3. This Chapter 5 Experimental Results and Discussion 90 results in a 14.3 wt% weight loss for a pure alumina sol-gel phase and 2.2 wt% weight loss for the composite sol-gel with 86.2 v o l % calcined alumina, as shown in Region II of Figure 5.2-1. The weight loss of both C S G and pure alumina gel as a function of sintering temperature is approximately constant in Region I E of Figure 5.2-1 because alumina gel has transformed to alumina. 25 20 --** 5, 15 --(0 (0 o —I .c 10 -O) '53 5 -0 -I* 100 300 Pure Alumina Sol-Gel I I I Composite Sol-Gel with 86.2 vol% AI2OJ 550 °C 900 500 700 Temperature, (°C) F i g u r e 5.2-1 Weight loss vs temperature for pure alumina sol-gel and composite sol-gel with 86.2 v o l % calcined alumina after drying at 100 °C for 20 hours 5.2.2 I n t e r a c t i o n b e t w e e n S o l - G e l M a t r i x a n d C e r a m i c F i l l e r s It is believed that alumina gels form by the step-wise linkage of A l O O H dimers via double chains to form a 3D-network [178]. This causes an increase in the polycondensation reaction rate, and gelation time becomes very short (1-5 min). The following phase transformations take place upon subsequent heat treatment of the resulting hydroxides: 7AIOOH 4 5 0 ° c - 5 5 0 ° c ) 7 A 1 2 0 3 8 S Q ° C ) 5Ai 2o 3 '°50°c >0M2o 3 1 2 0 0 ° c m A l 2 0 3 Chapter 5 Experimental Results and Discussion The calcined alumina/alumina sol slurry is stabilized because the positively charged small size (5-50 nm) colloidal clusters interact with the large A 1 2 0 3 particles (-300 nm) and increase the electrostatic repulsion force. During the gelation process, the hydrolyzed surface of alumina particles (=A10H) enters a polycondensation reaction with the alumina sol clusters: =A!OH + A l O O H - » = A l - 0 - A l = + H 2 0 Alumina Sol-Gel Matrix Phase Chemical Bonds Surface of and Hydro-bonds Hydro-Alumina at Interface K / -AI; Al o - A f A i - o - A f > I - O - A I = cr \ f Figure 5.2-2 Schematic of the proposed interaction of the alumina sol with hydrolyzed alumina particles during polycondensation reaction. The coordination number of aluminum ions is 6. The proposed structure resulting from this reaction is illustrated in Figure 5.2-2. It is expected that both hydrogen bonds and ionic/covalent bonds are formed at the interface of alumina particles/sol-gel matrix. Chapter 5 Experimental Results and Discussion 92 5.2.3 Properties of C S G Composites Properties of Alumina/Alumina Composite Figure 5.2-3 shows the relative density of alumina/alumina C S G with different sol-gel matrix contents as a function of sintering temperature. The increase of sol-gel matrix content in the composite results in the decrease of sintering temperature. The C S G composite with 24 v o l % sol-gel matrix was sintered to 98.5% densification at 1400°C for 1 hour. However, the calcined alumina without sol-gel matrix was sintered to only 95% densification at 1600 °C. 100 50 1100 1200 1300 1400 1500 Sintering Temperature, (°C) V m • 0.24 o 0.14 A 0.1 • 0.05 • 0 1600 1700 Figure 5.2-3 Relative density of alumina/alumina composite sol-gel ceramics with different sol-gel matrix content as a function of sintering temperature. Figure 5.2-4 shows the microhardness (HV) of the alumina-alumina C S G without additives, as a function of the sintering temperature. The microhardness of pure sol-gel derived alumina (0 v o l % calcined phase), although relatively higher at low sintering Chapter 5 Experimental Results and Discussion temperatures, becomes lower than that of the composite sol-gel after sintering at 1400 °C for 3 hours. It is believed that the pure sol-gel derived alumina experiences microcracks during sintering due to large densification strain. The final microhardness increases with addition of calcined alumina as a secondary phase, to reach a maximum of 17.5 G P a at about 86 vo l%, and decreases thereafter. The microhardness of the control sample of pure calcined alumina (100 v o l % calcined alumina line in Figure 5.2-4), sintered in parallel with C S G , reached 12.4 GPa at 1400 °C. This suggests a strong effect of the sol-gel-derived phase on the sinterability of the composite. C S G hardness is inversely proportional to porosity at different sintering temperatures, as illustrated in Figure 5.2-5. re CL o <n <n o c TJ >_ (0 o 1_ o 18 15 12 9 6 ^ 0 900 C a l c i n e d A lumina A 1 6 , (vol%) o- 0 -a-7 4 ,*' tS - • - 8 6 91 x s y ' - * - 9 5 J + 100 I I ' ' I I 1000 1100 1200 1300 Sintering Temperature, (°C) 1400 Figure 5.2-4 Microhardness ( H V 1 K g ) vs sintering temperature for alumina-alumina C S G (sintering time: 3 hours). Chapter 5 Experimental Results and Discussion 9 0 0 1 0 0 0 1 1 0 0 1 2 0 0 1 3 0 0 1 4 0 0 Sintering Temperature, (°C) Figure 5.2-5. Porosity as a function of sintering temperature for alumina-alumina C S G Properties of Alumina/Zirconia Sol-Gel Composite Similarly to the AI2O3 filler powder/Al203 sol-gel matrix composite, the Z r 0 2 filler powder /Al 2 0 3 sol-gel matrix composites were fabricated. The microhardness of the Zr0 2 /Al 2 03 composites increases with increasing alumina content. The highest microhardness of composites reached was 15 GPa, as shown in Figure 5.2-6. Meanwhile, the porosity of composite decreases with increasing sintering temperature and alumina sol-gel matrix content, as illustrated in Figure 5.2-7. Chapter 5 Experimental Results and Discussion 0 I -4 • 1 1 1 1000 1100 1200 1300 1400 Sintering Temperature, (°C) F i g u r e 5.2-6. Microhardness ( H V i K g ) vs sintering temperature of the composite of Z r 0 2 - 3 w t % Y 2 0 3 and alumina sol-gel matrix. 1000 1100 1200 1300 1400 Sintering Temperature, (°C) F i g u r e 5.2-7. The porosity of Z r 0 2 / A l u m i n a composites as a function of sintering temperature. 5.2.4 E f f e c t o f M g O o n S i n t e r a b i l i t y o f C S G C e r a m i c s Figures 5.2-8 and 5.2-9 show variation of microhardness and porosity of C S G as a function of M g O content in the alumina sol-gel derived phase. The microhardness of C S G Chapter 5 Experimental Results and Discussion increases and the porosity of G S G decreases with increasing M g O content in alumina sol-gel phase, as anticipated. However surprisingly, the microhardness of C S G with 2 mol% M g O in the sol-gel matrix phase appears to be in excess of 20 GPa, the value approaching that of a single crystal of alumina. Porosity of these C S G is less than 1 vo l%, as illustrated in Figure 5.2-9. Calcined Alumina 10 8+" 1300°C/91.4vol% frf3Wf3794vol% •+-0 0.5 1 2 MgO Content in Alumina Sol-Gel Matrix, (mol%) Figure 5.2-8. Microhardness ( H V l K g ) vs M g O content in alumina sol-gel matrix phase, for C S G sintered at 1300°C and 1400°C „ 6 a? o 5 £ 4 3 o o Q. (Calcined Alumina Content (vol%) of HI 86 .2 ^ Composite Sol-Gel [rrjj g-| ,6 ^ 9 4 + -+-0 0.5 1 2 MgO Content in Alumina Sol-Gel Matrix Phase, [mol%] Figure 5.2-9. Porosity vs M g O content in alumina sol-gel matrix phase,for C S G sintered at 1400°C Chapter 5 Experimental Results and Discussion 97 5.2.5 Microstructures of C S G Ceramics Figure 5.2-10 shows the microstructure of 86.2 v o l % calcined alumina C S G sintered at 1300°C and 1400°C for 3 hours. Two kinds of pores can be identified in these C S G . The "small" pores (< 0.1 pm), at the triple junctions of the grains, shrank rapidly at increased sintering temperature and with increased content of the sol-gel derived matrix phase. The "large" pores (~1 pm) were produced through imperfect consolidation of the C S G slurry. Most of the "large" pores were open and remained interconnected until C S G was sintered at 1400 °C. The average grain size at 1300 °C was about 0.4 pm, having grown slightly from the starting calcined alumina. The average grain size at 1400 °C was about 1.0 pm. Figure 5.2-10. S E M micrograph of C S G with 86.2 v o l % calcined alumina sintered at (a) 1300°C and (b) 1400°C for 3 hours. Bar=2 pm Figure 5.2-11 shows the microstructures of of 86.2 v o l % calcined alumina C S G with different M g O content in sol-gel matrix phase, fired at 1400 °C for 2 hours. The average pore size and content decreases with increasing M g O in the sol-gel matrix phase. Chapter 5 Experimental Results and Discussion 98 The X-ray map of the M g O distribution in the C S G composites is shown in Figure 5.2-12. The M g O is homogeneously distributed at the grain boundaries and triple junctions. Also, these results confirm indirectly that the sol-gel matrix uniformly covers the surface of the ceramic particles and fits into triple junctions to form a 3-D network. Chapter 5 Experimental Results and Discussion 99 MgO Distribution in Grain Boundaries and Triple Junct ions x30k 0000 £ 0 k V l P f f t Figure 5.2-12. The X-ray map of the M g O distribution in the alumina/alumina C S G Figure 5.2-13 shows the S E M morphology of (a) 6 4 v o l % Z r 0 2 - 3 6 v o l % A l 2 0 3 and (b) 50vol%ZrO 2 -50vol%Al 2 O3 composites fabricated by sol-gel processing. The Z r 0 2 (lighter phase) and alumina (darker phase) were homogeneously distributed into the sol-gel matrix. The microstructure of the composite with uniform sub-micrometer grains was successfully obtained. Chapter 5 Experimental Results and Discussion Figure 5.2-13. S E M morphology of alumina/zirconia composites: (a) 6 4 Z r 0 2 -3 6 A 1 2 0 3 and (b) 5 0 Z r O 2 - 5 0 A l 2 O 3 composites Figure 5.2-14 shows a Back Scatter image of Z r 0 2 / S o l - G e l matrix phase, indicating that alumina sol-gel phase resides at grain boundaries between Z r 0 2 particles and in triple junctions among Z r 0 2 particles. Chapter 5 Experimental Results and Discussion 101 Figure 5.2-14. Back Scatter picture of S E M of Zr02 /alumina sol-gel matrix. The lighter phase is alumina grains and sol-gel matrix phase and the darker phase is ZrC»2 5.2.6 Summary Alumina composite sol-gel (CSG) ceramics have been produced by dispersing calcined alumina powder in alumina sol. It has been found that the alumina sols acts as a dispersant and sintering accelerator for the calcined alumina. The sol-gel matrix phase of C S G bonded to calcined alumina forms a strong 3-D network. The microhardness and porosity of C S G depends on the sintering temperature and sol-gel phase content. When sol-gel matrix phase is 13.8 vo l%, the microhardness reaches the maximum value (17.5 GPa) for pure alumina-alumina C S G . The porosity decreases with increasing sol-gel matrix phase. The microhardness of composite sol-gel with M g O increases to more than 20 GPa and porosity decreases to less than 1 v o l % at a sintering temperature of 1400 °C. The sintering kinetics and final microstructure of C S G is strongly affected by M g O content in the sol-gel derived matrix phase. Chapter 5 Experimental Results and Discussion 5.3 SiC-Alumina Composite Sol-Gel Ceramics 102 Composites of SiC-Al2C»3 hold great promise for application as structural components and as wear-resistant elements, e.g. cutting tools and forming dies. The incorporation of a second phase (SiC) into the alumina ceramic matrix can lead to significant increases in fracture toughness, the level of toughening being strongly dependent on morphology of the second phase (i.e. particles, whiskers, or fibers). The increase of resistance to fracture initiation is believed to be the mechanism behind strengthening of ceramics as results of dispersion of the nanometer-size (-10 - 100 nm) secondary phase. Unfortunately, it is difficult to sinter to full density of the silicon carbide/alumina composite because of the covalent nature of S i C bond. The main objective of the present section was to investigate microstructure development in pressureless sintering of S i C - A l 2 0 3 C S G composites, where the fraction of alumina was replaced by hydrated alumina sol. 5.3.1 Mechanical Properties of SiC/Al 2 0 3 CSG The preparation of samples was identical to that described in Section 4.2.2. Microhardness of both types of A l 2 0 3 - S i C composites (i.e. composite sol-gel specimens C S G and composite water-dispersed C W D ) , sintered at 1850 °C for 1 hr, is shown in Figure 5.3-1 as a function of S i C content. The microhardness of A l 2 0 3 - S i C composites ( C S G , closed points) dispersed in alumina sol increases with increasing content of S i C up to 22.9 G P a at 50 vo l% S i C , and then starts to decrease. The microhardness of A l 2 0 3 - S i C composites dispersed in water ( C W D , open points) decreases continuously after reaching a maximum of 18.5 G P a at 5 v o l % S i C . Figure 5.3-2 shows microhardness of S i C - A l 2 0 3 sol-gel composites only (CSG) sintered at 1700°C, 1800°C, 1850°C and 1900°C for 1 hr, as a function of S i C content. Microhardness of Chapter 5 Experimental Results and Discussion 103 the composites sintered at 1700 °C increases with S i C content up to 12 vo l%, and then decreases. Similar data for 1800°C sintering shows a broad maximum at about 35 v o l % S i C . When the sintering temperature was further measured to 1850 °C or 1900 °C, the microhardness of composites continued to increase until reaching 22.9 G P a at 50 v o l % S i C . 24 10 -I 1 1 1 1 1 1 1 0 10 20 30 40 50 60 70 SiC Content, (vol%) Figure 5.3-1. Microhardness of both types of Al203-SiC composites: dispersed in alumina sol (CSG, closed points) and dispersed in water ( C W D , open points), as a function of SiC content, sintered at 1850 °C for 1 hr. SiC Content, (vol%) Figure 5.3-2. Microhardness of S i C - A l 2 0 3 sol-gel composites only (CSG), as a function of S i C content, sintering at 1700°C, 1800°C, 1850°C and 1900°C for 1 hr. Chapter 5 Experimental Results and Discussion \ 04 Figure 5.3-3 show the relative densities of SiC/Al203 composites as a function of S i C content. The relative densities of the composites with sol-gel matrix (CSG) decrease slightly from 0.99 to 0.972 when S i C content increases from 5 vo l% to 50 vo l%, and then decrease faster. The relative densities of the composites without sol-gel matrix ( C W D ) decrease sharply from 0.985 to 0.83 when S i C content increases from 5 vo l% to 50 vo l%. This result confirms the microhardness of the composites in Figure 5.3-1. SiC Content, (vol%) Figure 5.3-3. Relative density of both types of A l 2 0 3 - S i C composites: dispersed in alumina sol ( C S G , closed points) and dispersed in water ( C W D , open points), as a function of S i C content, sintered at 1850 °C for 1 hr. Alumina Sol-Gel Matrix Phase -A I ^^ I -O-A I^A I -O-A l ; Chemical Bonds S i0 2 Thin Film on the and Hydro-bonds Surface of SiC Particle H<g H<P H<° O - S * - A I ^ I - 0 _ A | ^ A l - 0 - A I ^ > l - 0 - S i C -H>0 H>0 , H P -S i 'c -AI; yy-o-Ai y y - O - A l ^ - O - S f e S i C Figure 5.3-4. Schematic interactions of hydrated silica f i lm on the surface of S i C particles with the alumina gel. Chapter 5 Experimental Results and Discussion 105 As shown previously in Section 5.1, dispersion of alumina and SiC particles is substantially improved in alumina sol, as compared to water of similar acidity, e.g. the average agglomerate size is decreased by at least 50%. The proposed mechanism involves absorption of the positively charged sol clusters (A104Ali 2(OH)24(H 20)i2 7 + ) on the surface of SiC particles, normally coated with ~1 nm film of hydrolyzed silica =Si(OH)\ The following chemical reaction may therefore occur at the interface between the particle and sol clusters: =Si(OH)" + A104Al l 2(OH)24(H20),2 7 + =Si-0-Al-A10 4 Al 1 i(OH)23(H 2 0) 1 2 6 + + H 2 0 As a result, a 2-35 nm thick layer may form on the surface of silicon carbide particles as proposed in the schematic Figure 5.3-4. Effectively, a continuous film of hydrated alumina coats both SiC and AI2O3 particles of the CSG composite dispersed in alumina sol. This film is absent in the CWD composites dispersed in water. It is anticipated that the film constitutes the fast diffusion path along boundaries between SiC/SiC and SiC/A^C^ grains. The active A1 2 0 3 sol-gel originating phase interacts with thin Si02 film on the surface of SiC, to form mullite phase (3Al203-2Si02), as illustrated in Figure 5.3-5, or eutectic liquid phase during sintering above ~1826°C. This fast mass transfer path is created for all kinds of boundaries in the system (i.e. SiC/SiC, SiC/ A1 2 0 3 and AI2O3/AI2O3 boundaries), aiding densification of CSG composites even with very high content of SiC, e.g. 50 vol%. The CWD composites appear to increase their hardness for small amounts of SiC, i.e. up to about 10 vol%. For this small content of SiC in CWD, about 90 % of the neighbors of the SiC grains are particles of alumina. The thin Si02 film on the surface of SiC particles may react with A1 2 0 3 to form mullite during sintering samples at 1850 °C, resulting in an increase of microhardness of the composite. The increase of number of SiC neighbors to SiC grains with increasing SiC content in CWD composites results in decreased densification rate due to lack of Chapter 5 Experimental Results and Discussion 106 mass transfer mechanisms at boundaries and volume of SiC grains. The small amount of intergranular SiC»2 does not apparently constitute a sufficient liquid phase assistance to sintering. Additionally, less effective dispersion of SiC/ A1203 in water, as compared to alumina sol, adds more SiC/SiC contacts which resist sintering. 2000 °c 1800 1600 1 4 0 4 0 20 40 60 80 100 SO, Figure 5.3-5. The phase diagram of Al203-SiC»2 system [167] SO, + Uul I I I I J 1 L 5.3.2 Microstructures of AhCVSiC CSG Figures 5.3-6 and 5.3-7 show the SEM morphology of the polished surface of 20vol%SiC-80vol%Al2O3 and 50vol%SiC-50vol%Al2O3 composites sintered at 1700°C for 1 hour, respectively. It is seen that the 20vol%SiC-80vol%Al2O3 composite is much denser than 50vol%SiC-50vol%Al2O3. Chapter 5 Experimental Results and Discussion 107 F i g u r e 5.3-6. The SEM morphology of 20vol%SiC-80vol%Al2O3 sol-gel composites sintered at 1700°C for 1 hour. F i g u r e 5.3-7. The SEM morphology of 50vol%SiC-50vol%Al2O3 sol-gel composites sintered at 1700°C for 1 hour. Figures 5.3-8 and 5.3-9 show the T E M morphology of 50vol%SiC-50vol%Al2O3 CSG composites sintered at 1850 °C. Figure 5.3-8 shows the triple junction of SiC-SiC-Al 2 0 3 grains. EDX analysis (ignoring carbon and oxygen) indicates 30wt% Si and 70wt % Al in this triple junction phase, confirming the proximity to mullite composition. A similar interlayer (-30 nm Chapter 5 Experimental Results and Discussion 108 thick) between S i C grains contains approximately 15% Si and 85% A l , confirming that S iC has been homogeneously dispersed into the alumina sol-gel matrix. This interlayer is believed to be liquid at 1850 °C sintering temperature to offer a fast mass diffusion path for densification. Figure 5.3-8. Triple junction morphology ( S i C / S i C / A l 2 0 3 ) for 50vol%SiO 2 -50vol%Al 2 C»3 composite dispersed in alumina sol, and sintered at 1850 °C for 1 hour. S G is alumina sol-originating intergranular phase. Figure 5.3-9. Alumina sol-originating coating on surface of S i C particles (sample from Fig . 5.3-8) Chapter 5 Experimental Results and Discussion 109 Figure 5.3-9 shows a micropore (-300x50 nm) between S i C and alumina grains. Again, a coating layer (-40 nm thick) on the S i C grain is clearly observed. The elemental composition of the layer (ignoring carbon and oxygen ) indicates 100 wt% A l at outside edge of layer, decreasing to approximately 55 wt% at the interface between coating layer and S i C grain. The concentration gradient of Si increases from 0 % in the middle of the layer to approximately 45 % at the interface between the coating layer and S i C grain. Therefore, it appears that the interlayer indeed originates from alumina sol-gel, which reacts with S i 0 2 on the surface of S i C to form mullite. The composite sintered at 1800 °C is illustrated in Figure 5.3-10 exhibiting relatively large pores in the triple junctions. Figure 5.3-10. Triple junction and micropore morphology for 50vol%SiO2-50vol%Al2O3 composite dispersed in alumina sol, and sintered at 1800 °C for 1 hour. S G is alumina sol-originating intergranular phase. The small S i C particles in intragranular positions give rise to strain contrasts in the surrounding AI2O3 lattice in Figure 5.3-11. Vitreous intergranular films present at internal interfaces accommodate the lattice mismatch between AI2O3 matrix and S i C inclusions. The facilitate strain relaxation around larger S i C particles in intergranular positions by particle Chapter 5 Experimental Results and Discussion 110 rearrangement during the sintering processing. However, because of considerable thermal expansion mismatches, radial tensile stresses exist around S i C particles upon cooling. Figure 5.3-12 shows the grain distribution of S i C and A 1 2 0 3 in the composites. i 800 nm F i g u r e 5.3-11. S i C particles in intragranular positions in alumina matrix F i g u r e 5.3-12. The grain distribution of S i C and AI2O3 in the composites Chapter 5 Experimental Results and Discussion 111 The microstructure of S iC-Al 2 03 composite without sol-gel phase (CWD) dispersed in water (i.e. free of sol-gel alumina) is shown in Figure 5.3-13 for of 50vol%SiC-50vol%Al2O3. Extensive agglomeration of S i C is evident, together with substantial porosity which appears to be linked to S i C grains. The interfacial composition between S i C grains indicates the absence of a mullite-type phase. A ~2 nm thick fi lm of S i 0 2 is present between S i C grains. This film is apparently not sufficient to offer fast mass transfer path for densification of the C W D composite. The boundary between S i C grains resembles a classical interparticle "neck" characteristic of the initial stage of sintering. This is in clear contrast to the intimate shape of the boundary between S i C grains sintered in the presence of alumina sol (refer to C S G compositions in Fig . 5.3-8 and 5.3-9), characteristic of grain boundaries in fully dense ceramics. Figure 5.3-13. Agglomerates of S i C particles joining through necks rather than grain boundaries, in 50vol%SiO 2 -50vol%Al 2 O3 composite, water-dispersed (CWD) , sintered at 1850 °C for 1 hr. 5.3.3 Summary The C S G alumina/silicon carbide composites with high content S i C have been successfully sintered by pressureless sintering at 1850 °C, using an active sol-gel alumina Chapter 5 Experimental Results and Discussion 112 additive. The additive is introduced to the system through dispersion of S i C and AI2O3 grains in alumina sol, which aids deagglomeration of the system. It is proposed that the sol-gel phase acts as a sintering additive through formation of a continuous intergranular f i lm around both alumina and silicon carbide grains. The active intergranular phase serves as a fast diffusion path during sintering of the composite. In particular, the intergranular f i lm of sol-gel originating alumina / silica offers a fast mass transfer path to assist in sintering of S i C / S i C grains. Absence of such a fast diffusion path in water-dispersed S i C - A l 2 0 3 results in poor densification of the composites for S i C content above 10 vo l%. The highest microhardness obtained was 22.9 GPa, for the composition of 5 0 v o l % S i C - 5 0 v o l % A l 2 O 3 . Chapter 5 Experimental Results and Discussion 5.4 Composite Sol-Gel Ceramic Coatings 113 It is difficult to make crack-free monolayer sol-gel coating films thicker than 1 \im by using alkoxide sol due to the large shrinkage during drying and sintering. This restricts the potential applications of these coatings for wear and corrosion protection since a thickness greater than 10 p:m is usually required. In this section the processing of composite sol-gel coating is described and the microstructure and properties of the coating are discussed. 5.4.1 Processing of the C S G Coatings The coatings were deposited by dip coating, spray coatings and electrophoretic deposition (EPD). In dip-coating, a substrate is usually withdrawn vertically from the coating bath at constant speed V. The moving substrate entrains the liquid in a fluid mechanical boundary layer that divides in two layers above the liquid bath surface, returning the outer layer to the bath. There are six main factors that govern the fi lm thickness [168]: (1) viscous drag upward on the liquid by the moving substrate; (2) force of gravity; (3) resultant force of surface tension in the meniscus; (4) inertial force of the boundary layer liquid arriving at the deposition region; (5) surface tension gradient; and (6) the disjoining pressure. When the liquid viscosity T] and substrate speed are high enough to lower the curvature of the gravitational meniscus, the deposited fi lm thickness ht is determined by balance of the viscous drag, proportional to rjV, and gravity force pgh [168]: h=-Z&- (5 4-1) where the constant c is about 0.8 for Newtonian liquids. Figure 5.4-1 shows the thickness of dip coatings as a function of withdrawal speed V of substrate, at constant viscosity of the liquid. Chapter 5 Experimental Results and Discussion 114 Figure 5.4-2 shows the thickness of dip coatings as a function of the viscosity of liquid at constant withdrawal speed. The thickness of coatings increases with increasing the viscosity of liquid and withdrawal speed of substrate. It appears that the thickness of composite sol-gel coatings calculated by using Eq.(5.4-1) (solid line) is higher than that of experimental data (points), especially for higher viscosity and withdrawal speed. It seems that the properties of liquid with high viscosity are different from that of Newtonian liquids. Therefore, the constant c of Eq.(5.4-1) is changed from 0.8 to 0.6 (dotted line) to better fit the data. 60 50 3, • Experimental Data a> c a> 20 0 0 2 4 6 8 Withdrawal Speed, (cm/min) 10 12 F igure 5.4-1. Thickness of C S G coatings as a function of withdrawal speed of the substrate at a constant viscosity (20 cP) of the liquid. The solid line presents the results calculated using Eq.(5.4-1). The points are the experimental data. The dotted line is the result of Eq.(5.4-1) after constant c was changed from 0.8 to 0.6 Chapter 5 Experimental Results and Discussion 115 45 of S 3 0 -• Experimental Data — Calculation 5 H 0 0 10 20 30 40 50 60 Viscosity, (cp) Figure 5.4-2. Thickness of C S G coatings as a function of the viscosity of liquid at constant withdrawal speed (4 cm/min) of the substrate. The solid line presents the results calculated using Eq.(5.4-1). The points present the experimental data. The dotted line is the calculation result of Eq.(5.4-1) after constant c was changed from 0.8 to 0.6 Coating deposition by spray coating was controlled by adjusting the air pressure and viscosity of the liquid. The electrophoretic deposition (EPD) was capable of achieving higher solids concentrations in the deposition fi lm, resulting in lower shrinkage during drying and sintering. For example, the f i lm coated by E P D was about 24 v o l % solid for a 1 M alumina sol (i.e. has 76% pores and liquid), compared with about 1.2 v o l % solid by dip coating and spray coating. In E P D , the ceramic particles coated by positively charged alumina sol clusters in liquid move to the cathode (metallic substrate) in the electric field applied between two electrodes. A uniform and 24 v o l % dense coating films was deposited on the surface of cathode. The deposition speed and quality of coatings were controlled by adjusting the voltage of the two electrodes, the viscosity of liquids, particle size, and concentration of ceramic particles. The C S G coatings were dried at room temperature and at 80°C for 20 - 100 min, depending on the coating thickness and concentration, and then sintered at 400°C-600°C for 15 Chapter 5 Experimental Results and Discussion 116 min. Figure 5.4-3 schematically illustrates the curing process for C S G of coatings. Figure 5.4-3a shows the C S G coating as deposited on the substrate before drying. The wet alumina sol-gel matrix forms a network after gelation and polymerization reactions are completed. Ceramic Filler Wet Alumina Gel Matrix Y-AI2O3 on Filler Surface Ceramic Filler Figure 5.4-3. Schematic diagram of curing process of C S G coatings: (a) the coating before drying; (b) the coating dried at 100°C for 1 hour; (c) the coating sintered at 550°C for 20 min. Chapter 5 Experimental Results and Discussion 117 Figure 5.4-3b shows that the ceramic particles were rearranged and the pore collapsed due to the evaporation of solvent and sol-gel matrix shrinkage. Figure 5.4-3c shows the C S G coatings sintered at 550°C. Aluminum hydroxide gel is transformed to Y-AI2O3 and the structural water is removed. It is believed that at this stage chemical bonds were formed between the sol-gel derived matrix phase, the dispersed ceramic powders, and the substrate. 5.4.2 P o s t D e p o s i t i o n T r e a t m e n t o f C S G C o a t i n g s The C S G coatings cannot be densified i f the sintering temperature is below 1000°C. Some metallic substrates, such as aluminum, and magnesium alloys, require curing temperature below 600°C or even lower. Additional alumina sol-gel sealing treatment was therefore carried out to improve the mechanical properties and to decrease the gas permeability of C S G coatings. The percentage of the porosity infiltrated using 0.5 M alumina sol at pressure 200 mTorr, as a function of infiltration time, is shown in Figure 5.4-4. The increase of infiltration time results in an increase of impregnation of the porosity of C S G , reaching 99 v o l % porosity filled by alumina sol after infiltrating for 25 min. 100 Vacuum: P=200 mTorr lnfiltrant:0.5 M Alumina Sol 0 0 10 20 30 40 Infiltrating Time, (min) F i g u r e 5.4-4. Infiltrated porosity of C S G coatings as a function of infiltrating time, using 0.5 M alumina sol and at 200 mTorr vacuum. Chapter 5 Experimental Results and Discussion 118 The infiltrated depth (/) of porous C S G coating is parabolic in time (r) and depends on fluid viscosity (77), surface tension (y), pore radius (r) and wetting angle (0) [169]: l>=^°-rt (5.4-2) 2T7 Therefore, good penetration of the sealant into the porous coating is assured through its low viscosity, high surface tension and wetting of the coating materials. Figure 5.4-5 shows the weight gained and infiltrated porosity percentage as a function of the concentration of alumina sol at the constant infiltration time (25 min) and pressure (200 mTorr). The results indicate that the weight gained increases with an increases of the concentration of alumina sol, but the infiltrated porosity percentage of C S G coatings decreased due to increases of the viscosity of alumina sol, as discussed in Section 5.1. Concentration of Alumina Sol, (M) Figure 5.4-5. The weight gained and fraction of infiltrated porosity as a function of the concentration of alumina sol at the constant infiltration time (20 min) and pressure (200 mTorr) Chapter 5 Experimental Results and Discussion 119 The results of gas permeability as a function of the concentration of alumina sol and the impregnation times are shown in Figure 5.4-6. The permeability decreased with increasing the concentration of alumina sol, reached the minimum value at a concentration of 1.5 M , and then starts to increase. The yield on decomposition from the dry alumina gel increased with increasing the concentration of alumina sol, so that the porosity of the C S G coating is expected to decrease. However, the viscosity of alumina sol increases when the concentration of alumina sol was more than 2 M . Thus the small infiltration depth caused the increases of the gas permeability for high concentration of alumina sol. The gas permeability of C S G coatings infiltrated by 1.5 M alumina sol decreased by 60% with single infiltration (K0 = 102.1xl0" 4 centi-darcys, K s i n g i e = 37.3xl0" 4 centi-darcys), by over 80 % with double infiltration (18.2xl0" 4 centi-darcys), and by over 90% with triple infiltration (9.7xl0" 4 centi-darcys). The significant reduction in the gas permeability confirms that the multiple infiltrations of C S G coatings with alumina sol sealant are effective in reducing the number of connected pores and the capillary diameters in the coatings. 0 0.5 1 1.5 2 2.5 3 Concentration of Alumina Sol, (M) Figure 5 . 4 - 6 . The gas permeability of C S G coating as a function of concentration of alumina sol with multiple inpregnations. Chapter 5 Experimental Results and Discussion 120 5.4.3 Micros t ructure and Properties of C S G Coatings Figures 5.4-7 and 5.4-8 show S E M morphology of the cross sections of C S G coatings fabricated using different methods, A to H . Figures 5.4-9 and 5.4-10 show the surface microhardness and the bonding strength between the coatings and substrates, as a function of the different processing methods for samples from A to H , respectively. Figure 5.4-7 The morphology of composite sol-gel coatings fabricated by the methods from A to D Chapter 5 Experimental results and Discussion 121 Figure 5.4-8. The morphology of the composite sol-gel coatings fabricated by the methods E to H Chapter 5 Experimental Results and Discussion 122 - 8 Ci o Processing methods Figure 5.4-9. The surface microhardness of composite sol-gel coatings on stainless steel substrates as a function of different processing methods B C D E F Processing Methods G H Figure 5.4-10. The bonding strength between composite sol-gel coatings and stainless steel substrates substrates as a function of different processing methods Chapter 5 Experimental Results and Discussion 123 (1) The deposition process A was to lower the metallic substrate into C S G slurry, and then withdraw it from the vessel at constant speed. The coating thickness (2-200 um) was controlled by the withdrawal speed (1-20 cm/min) and viscosity (2-30 mPa-s) of the slurry. Coating A has the highest porosity (40 vol%), lowest surface microhardness (0.5 GPa), and lowest bonding strength (3.7 MPa) . It cannot be used for corrosion protection and wear applications, but it may be used to produce ceramic membranes. (2) The deposition process B included infiltration of the coating A with 1.5 M alumina sol in vacuum and then withdrawal at a constant speed. Coating B has the same basic microstructure as the coating A but it has a 3 um thick dense layer on the surface. The surface microhardness (0.8 GPa) and bonding strength (3.9 MPa) slightly increase as compared to the coating A . It may be used for corrosion protection and dielectric coatings on metallic substrates. (3) The deposition process C was to infiltrate coating A with 1.5 M alumina sol using electrophoretic deposition (EPD). The sol clusters with positive charges were deposited throughout the connected pores and onto the metallic substrate. E P D processing gives high density close-packed structure prior to sintering. Coating C has a denser structure and thicker (~5um) non-permeable layer on the coating surface. The surface microhardness (1.1 GPa) and bonding strength (4.9 MPa) is increased. This coating may be used for corrosion protection of metallic substrates. (4) The chemically bonded coating D has approximately 12 v o l % porosity but the surface layer is dense. The surface microhardness (4.5 GPa) and bonding strength (20.3 MPa) are significantly increased. Coating D can be used for corrosion and wear protection of metallic substrates, especially, after additional surface heat treatment to close open porosity [170,171]. Chapter 5 Experimental Results and Discussion 124 (5) Coating E has two different layers: layer 1 is a thicker dense f i lm (30 urn thick and ~12vol% porosity) on the outside of the coatings; layer 2 is a soft interfacial coating fi lm. Also , coating E has a non-permeable surface coating layer. The surface microhardness (4.2 GPa) and bonding strength (15.1 MPa) are lower than that of coating D . (6) Coating F has a similar structure as the coating E but it is much denser (~8 v o l % porosity in outside layer) than that of coating E . The surface microhardness (6.8 GPa) and bonding strength (35 MPa) are significantly higher than that of coating E . Coating F can be used for corrosion and wear protection of metallic substrates. (7) Coating G is dense (~7 v o l % porosity), and so has the highest surface microhardness (7 GPa) and bonding strength (41.8 MPa) . It can be used for corrosion and wear protection of metallic substrates. (8) The structure and properties of coating H treated by chemically bonded C S G are similar to that of the coating G , but the processing is more stable. The results of corrosion test for T i protected by C S G coatings (method D) , spray formed on the titanium substrates, are shown in Figure 5.4-11. The titanium bars with and without C S G coatings were heat treated in the same furnace at 600°C and 800°C for different times, and then the weight gained was measured. In comparison with the titanium without C S G coatings, the oxidation rate of titanium with C S G coating at 800°C was significantly decreased from 23.5 mg/cm 2 day to 1.14 mg/cm 2 day. The oxidation rate of titanium without C S G coatings at 600°C was 0.121 mg/cm 2 day. However, the oxidation rate at 600°C decreased to zero for titanium coated with C S G . Chapter 5 Experimental Results and Discussion 125 20 -O-TJ/800C -«-Ti+CSG/800C oTi/600C -•-Ti+CSG/600C 0 5 10 15 Test Time, (Day) Figure 5.4-11. The weight gain for T i (coated and uncoated) as a function of test time at 600°C and 800°C. 5.4.4 Summary Novel composite sol-gel (CSG) technology has been successfully used to fabricate thick (up to 600 pm) non-permeable, and crack-free ceramic coatings on a variety of substrates, such as stainless steel, titanium and aluminum. The thickness of C S G coatings was a function of the viscosity of the slurry and withdrawal speed of the substrate in dip coating processing. Sol-gel sealing treatments were carried out to improve the microstructure of the coatings, for application in corrosive gaseous and high temperature environments. Post deposition effects were investigated through measuring gas permeability of C S G coatings as a function of concentration of alumina sol and infiltration processing. The optimum concentration of alumina sol for post deposition treatment was found to be 1.5 M . A significant reduction (over 90%) of gas permeability of C S G coatings has been achieved by multiple infiltration. Chapter 5 Experimental Results and Discussion 126 Correlations between the processing methods, microstructure, and mechanical properties of C S G coatings were investigated. The advantageous properties of C S G coatings on metallic substrates were successfully achieved by combining C S G and chemical bonding technologies. The bonding strength between the metallic substrates and coatings was more than 42 M P a . The surface microhardness of the C S G coatings was more than 6.5 GPa, and non-permeable to gases. Chapter 6 Sintering Model for CSG Ceramics 127 C H A P T E R 6 SINTERING MODEL FOR CSG Ceramics 6.1 Sintering of Alumina Gels Sintering is a densification process driven by decreasing interfacial energy. The sintering material moves by viscous flow or diffusion in such a way as to eliminate porosity and thereby reduce the solid/vapor interfacial area. The primary driving force for sintering is reduction of the free surface energy of the system due to densification, as illustrated in Figure 6.1-1. Simultaneously, grain growth (coarsening) decreases system energy without contribution to porosity elimination. Figure 6.1-1. Sintering is a process of microstructural change involving contributions from: densification and coarsening Chapter 6 Sintering Model for CSG Ceramics 128 In gels, surface area is large (i.e. 100-500 m 2/g), so the driving force is great enough to initiate sintering at relatively low temperatures. OC-AI2O3 particles or seeds doped into alumina sol may act as nucleation sites for the 0- to a - A l 2 0 3 transformation and for the development of an aggregate-free, ultrafine-grained microstructure after transformation to OC-AI2O3 [71]. It is believed that A104Al i2(OH ) 2 4 (H 2 0)i2 + 7 is transformed during gelation into pseudoboehmite ( A l O O H ) due to the loss of the tetrahedrally coodinated aluminum [30]. Hydrated alumina gel forms by the step-wise linkage of A l O O H dimers into double chains to form 3D-network [39]. This causes an increase in the polycondensation reaction rate, and gelation time becomes very short (1-5 min). The following phase transformations take place upon subsequent heat treatment of the resulting hydroxides: y A I O O H 4 5 0 ° c - 5 5 0 O c ) y A l 2 0 3 8 5 0 ° c ) <5A1203 1 0 5 0 ° c )0A1 2 O 3 1 2 0 0 ° c ) O A 1 2 0 3 Large shrinkage (i.e. 30% linear) of alumina gel occurs during sintering to full density because the relative green density of dry alumina gel is low, approximately 0.33. The dried alumina gels doped with 2 wt% OC-AI2O3 as nucleation seeds were preheated at 550 °C for 2 hours, then sintered at different temperatures for 30 min. Linear shrinkage and relative density of alumina gel versus sintering temperature are shown in Figure 6.1-2 and Figure 6.1-3, respectively. The seeded alumina gels have fine uniform grains with pores only at grain boundaries during sintering. However, the unseeded alumina gels develop large vermicular "worm-like" grains with internal pores [71]. The vermicular grains densify slowly, because they are single crystals, so the pore volume is not intersected by many grain boundaries. The alumina gels seeded with small amount of CC-AI2O3 have sufficiently large surface energy, for the mass to be transported rather by the surface and volume diffusion rather than viscous flow. This result agrees well with the results reported by Kumagai [69,71]. However, when the amount of the filler phase increases, e.g. to 70 % to 95% range in C S G , the system does appear to follow the viscous sintering behavior. This phenomenon is elaborated in the sintering model for C S G , introduced in the next Section. Chapter 6 Sintering Model for CSG Ceramics 550 650 750 850 950 1050 1150 1250 1350 Sintering Temperature, (°C) Figure 6.1-2. Linear shrinkage of alumina gel doped with 2vol% (X-AI2O3 as a function of sintering temperature (30 min hold at each temperature). CL in c a> a a> > 0) rr 550 650 750 850 950 1050 1150 1250 1350 Sintering Temperature, (°C) Figure 6.1-3. Relative density of alumina gel doped with 2vo l% OC-AI2O3 as function of sintering temperature (30 min hold at each temperature). Chapter 6 Sintering Model for CSG Ceramics 6.2 Sintering Model for Composite Sol-Gel (SMCSG) 130 6.2.1 Idealized Composite Sol -Gel System In composite sol-gel (CSG) systems, ceramic particles are dispersed into alumina sol. It is assumed that ideally, the alumina sol-gel matrix phase homogeneously coats the surface of ceramic particles. The coated particles form the compact green body, as schematically illustrated in Figure 6.2-la. The micrometer-size ceramic particles are further considered to be non-sintering and incompressible bodies at temperatures below 1350°C. However, the large shrinkage of the active sol-gel matrix drives ceramic particle rearrangement, repacking and collapse of large pores during sintering temperatures from 1100°C to 1350°C. It is believed that in this temperature range mass transport in C S G can be globally discussed as viscous flow. The volume fraction of sol-gel matrix Vm in C S G can be obtained assuming uniform coating thickness (R - r) of sol-gel phase on the particles, as illustrated in Figure 6.2-la: V = R3 R3-r3 = 1- (6.2-1) For example, i f Vm = 0.14, (a value typical in this research) then: = 0.86, and thus, — = 0.95 R (6.2-2) The maximum linear shrinkage between two particles is limited by the thickness of the sol-gel f i lm and is given by: R-r R 1--R (6.2-3) Chapter 6 Sintering Model for CSG Ceramics 131 Ideally, = 0.05 for Vm = 0.14. In reality the linear shrinkage of the compact body of C S G is larger than 0.05 because of additional effects of rearrangement and repacking during sintering. Sol-Gel Matrix ,. _ a b F i g u r e 6.2-1 Two-sphere sintering model of composite sol-gel particles with the development of the interparticle bond during sintering, a) starting with a point contact at sol-gel coating layers of inclusions, b) neck growth creates a new grain boundary at particle contact of sol-gel matrix layers. The neck growth between the contacting particles, as illustrated in Figure 6.2-lb, is the key aspect of sintering. The neck size X is an indicator of sintering progress, and can be expressed as [172]: X"1 = Act (6.2-4) where X is the neck diameter and Ac is a constant which accounts for surface energy, atomic volume, Boltzman constant, absolute temperature, and the appropriate coefficients of diffusion. The parameter "m" can be linked to the sintering mechanism, i.e.: viscous or plastic flow X2 -< t (6.2-5) evaporation/condensation X3 oc t (6.2-6) Chapter 6 Sintering Model for CSG Ceramics 132 volume diffusion surface diffusion X5 oct X7 -ct (6.2-7) (6.2-8) The compact density (p«) is often expressed as a fraction of theoretical solid density (pT), in terms of relative density p = p«/pr. The relationship between relative density p and linear shrinkage AL/L0 can be obtained as: Po (6.2-9) where p0 is the initial relative green density of compact. According to Taylor's series expansion, the Eq.(6.2-9) can be expressed as: 2! J U 1 + 3 ^ + 3 ( 3 + 1 ) Po h AL 3(3 + l)(3 + 2) 3! + • (6.2-10) When the linear shrinkage — « 1, the Eq.(6.2-10) can be approximately expressed as: L A ~ l + 3 ^ Po K (6.2-11) AL For CSG ceramics, —is smaller than 0.15, and thus approximately linear relationship between relative density and linear shrinkage can be expressed as: i AL P-Po =lPo — (6.2-12) The validity of Eq.(6.2-12) is confirmed by Figure 6.2-2 which shows a linear relationship between relative density and linear shrinkage at the different initial densities. The error of linear approximation is about 10% at = 0.15. Chapter 6 Sintering Model for CSG Ceramics 133 Q_ 5^ •«—' '(/) c 0 D 0 > « 0 DC 0.05 0.1 Linear Shrinkage, (AL/L0) 0.15 Figure 6.2-2. Relative density of composites with different initial relative density as a function of linear shrinkage. The points were calculated using Eq.(6.2-9) and lines were calculated using Eq.(6.2-12). Although sintering models utilize the neck size X in Eq.(6.2-5) to (6.2-12), it is easier to measure the compact dimensional change rather than the neck size. Shrinkage is approximately related to the neck size by [173]: AL ( X V V 2 D , (6.2-13) B y combining Eq.(6.2-12) and Eq.(6.2-13), the relative density as a function of the neck size ratio is given by: P-PO=3PO V 2 D (6.2-14) 6.2.2 Viscous Sintering of C S G In the following discussion, it is assumed that densification of C S G systems may be approximated through viscous sintering. Frenkel [62] presented an analysis of the coalescence of Chapter 6 Sintering Model for CSG Ceramics 134 a pair of spheres, through viscous sintering, as shown in Figure 6.2-1. The centers of the spheres approach one another as the neck between them widens. The change in distance between the centers of two spheres is assumed to be equal to the linear contraction of a compact of such particles. This geometry is not as simple as it may seem, because the shape of the neck between the particles changes considerably by the time the centers move significantly. The mathematical description of the shape of the neck is therefore difficult, and is usually presented with severe simplifying assumptions. Frenkel obtained a relationship for the growth in diameter X, of the neck between spheres of diameter D [62]: f x } 2 = 3ysut/8riD (6.2-15) where ysu is free surface energy and r\ is the material viscosity, dependent on temperature [173]: rj=T]0 exp xkT, (6.2-16) where Q is the activation energy for viscous flow, k is Boltzmann's constant, T is absolute temperature, and r\0 is the reference viscosity. For the C S G , r\0 is a function of sol-gel phase content and porosity. Viscosity of a solid-liquid system rj0j depends on the volume fraction of solid Vs, approximately given as [173]: "--o^tf ( 6 2" 1 7 ) where r\0i is the reference viscosity of pure liquid phase and c is a constant which is related to solid grain size, geometry, and fractional coverage of the grain boundaries by liquid. Moreover, viscosity of porous materials depends on the relative density p0 and can be given by [85]: rio2=riofpop{l-poyX (6.2-18) Chapter 6 Sintering Model for CSG Ceramics \ 35 where rj0f is the reference viscosity of fully dense material, p and X are constant coefficients ( for example, p ~ 0.05 and X =1.67 for alumina system [85]). B y combining Eq.(6.2-17) and Eq(6.2-18), the reference viscosity of C S G can be approximated as: "•=»"-"• Vcv.yp.-'.i-p.r < 6 - 2 " 1 9 ) where h is a constant. Finally, replacing term r\0 in Eq.(6.2-16), the viscosity of C S G can be expressed as a function of relative density pQ and filler phase content Vs: hr\oLr\of KkT , (6.2-20) 6.2.3 F o r m u l a t i o n o f S i n t e r i n g M o d e l f o r C S G The sintering process of C S G may involve not only viscous sintering, but also crystallization and grain growth. Therefore, the progress of sintering may deviate from that predicted by Equations (6.2-13) and (6.2-15). This is taken into account by formulating a general parametric equation for density change. Combining Eqs. (6.2-4), (6.2-12), (6.2-15) and (6.2-20), the relative density of C S G can be expressed as a function of sintering time t and sintering temperature T through the following equation: p-po+ Bt exp (6.2-21) kT where b is a coefficient dependent on the volume fraction of the sol-gel matrix (refer to the following Section 6.3), Vm = 1 - Vs, and B is a coefficient given by: B = 9 y „ A ( l - c V , ) ' p : - ( l - p , ) ' 1 2 h D n A Chapter 6 Sintering Model for CSG Ceramics 136 where A is a constant, and p0 is initial relative density. The coefficient B in Eq.(6.2-22) is a function of the initial relative density of the composite and the volume fraction of the sol-gel matrix: B = f(ym,P„)= Cf, (Vm ) / 2 (p.) (6.2-23) where C is a constant, Vm is the volume fraction of sol-gel matrix, Vm = 1 - Vs in C S G . The term f\(Vm) is a function of the volume fraction of the sol-gel matrix given by: MVm)=(cVm+l-cf (6.2-24) a n d / 2 ( p j is a function of the relative density of the sol-gel matrix given by: fi{p.)=P.l-'(l-PoY . (6-2-25) The hypothesis is that this model can be used to predict the sinterability of the composites, and to analyze hydrostatic sintering stresses of the composite sol-gel ceramic system. Ultimately, it is believed that the model may help to understand sintering mechanism of C S G , although further future work is required to reach this objective (refer to Chapter 8). Chapter 6 Sintering Model for CSG Ceramics 6.3 Experimental Verification of the Sintering Model for CSG 137 6.3.1 Model Coefficients Calculated from Experimental Data A number of isothermal sintering experiments have been executed to verify the S M C S G model. The small size samples (10x10x2 mm 3 ) were used to avoid the influence of heat transfer on the sintering model. The samples were preheated at 550 °C for 1 hour, and then sintered isothermally at 1350 °C for variable time. The volume fraction Vm of the AI2O3 sol-gel matrix of alumina/alumina composite was varied from 0.05 to 0.24. The detailed experimental procedures have been described in Section 4.2.1. The relative densities of the composites are plotted as a function of the sintering time in Figure 6.3-1. For higher sol-gel matrix content in the composites, higher relative densities of the composite sol-gel ceramics are obtained at the same sintering time. In order to calculate the activation energy from Eq.(6.2-21), the C S G samples were sintered isothermally at 1250, 1300, and 1350°C for 1 hour. In order to determine coefficients b, B, and Q, all samples were prepared with the same initial relative densities of 0.56. 0.9 V, 0.85 ^ 0.55 " 1 1 1 1 1 0 50 100 150 200 250 Sintering Time, (min) 7=1350°C po=0.56 Figure 6.3-1. The relative densities of AI2O3/AI2O3 sol-gel composites sintered isothermally at 1350°C as a function of sintering time. The lines were calculated using S M C S G using parameters from Table 6.3-1 and the points are experimental data. Chapter 6 Sintering Model for CSG Ceramics 138 -1.2 j -1.6 --2 -o Q. 1 -2 4 -c _l -2.8 --3.2 --3.6 -2.2 2.6 3 3.4 3.8 4.2 4.6 5 5.4 Ln(0 Figure 6.3-2. A plot of Ln(p — po) as a function of Ln(t); experimental data are same as in Fig.6.3-1. The S M C S G Eq . (6.2-21) can be expressed in logarithm form: Ln(p-po) = bLn(t)--0- + Ln(B) (6.3-1) kT E q (6.3-1) is plotted in Figure 6.3-2, indicating good linear correlation between Ln(p - p o ) a n d Ln(t) as shown in Table 6.3-1: Table 6.3-l.The experimental data fit into S M C S G model at 1350°C. Volume Fraction of Sol-Gel Matrix, (Vm) Linear Equations R2 0.24 Ln(p - p „ ) = 0.336Ln(r) - 3.0641 0.998 0.14 Ln(p-p0) = 0.423Ln(f) - 3.625 0.999 0.1 Ln(p - p J = 0.49\Ln{t) - 4.1578 0.999 0.05 Ln(p - p0) = 0.561L«(r) - 4.709 0.995 Chapter 6 Sintering Model for CSG Ceramics 139 -1 -r -1.2 --1.4 --1.6 -1 -1.8 -• Q . -2 -p o =0.56 vm t = 60 min • 0.24 • 0.14 • 0.1 r — • — . ° 0 0 5 6.1 6.2 6.3 6.4 6.5 6.6 104/T, (1/K) Figure 6.3-3. A plot of Ln(p - po) as a function of 1/T. Values of Ln(p - po) are plotted in Figure 6.3-3 as a function of 1/T for a constant sintering time, and results are listed in Table 6.3-2 Table 6.3-2. Ln(p- po)as a function of 1/T for C S G samples sintered for 1 hour Volume Fraction of Sol-Gel Matrix, (Vm) Linear Equations R 2 0.24 Ln(p-po) = -0.6903 / T + 2.63 0.974 0.14 Ln{p -p0) = -0.7044IT + 2.486 0.990 0.1 Ln(p -p0) = -0 .7082 /T + 2.212 0.999 0.05 Ln(p -pB) = -0.6929 IT +1.823 0.996 The results in Table 6.3-2 indicate that the activation energy (Q) for viscous sintering of C S G is independent of the volume fraction of sol-gel matrix and the sintering temperature. The activation energy for C S G is approximately 58 kJ/mole, which is comparable to the activation energy of glass sintering, i.e. about 65 kJ/mol [177]. The coefficients B and b in Eq.(6.2-21) are Chapter 6 Sintering Model for CSG Ceramics 140 calculated using Eq.(6.3-1) and the data in Table 6.3-1 and Table 6.3-2. The results are listed in the following Table 6.3-3. Table 6.3-3. The coefficients b and B as a function of volume fraction of sol-gel matrix in Al 20 3/Al203 C S G vm B b 0.24 3.486 0.336 0.14 1.989 0.423 0.10 1.168 0.491 0.05 0.673 0.561 Reduction of volume fraction of sol-gel matrix in C S G results in a decrease of the coefficient B and an increase of the coefficient b, in Figure 6.3-4 and Figure 6.3-5, respectively. A n exponential function was chosen to empirically fit the coefficient b(Vm): b = 0.637 exp(-2.265V m ) R 2 = 0.992 (6.3-2) It can be also shown that b is independent of the initial relative density pQ. po=0.56 T=1350°C 0.05 0.1 0.15 0.2 0.25 Volume Fraction of Matrix, (Vm) Figure 6.3-4. Coefficient Ln(b) of the S M C S G as a function of volume fraction of sol-gel matrix of in A1203/A1203 composites Chapter 6 Sintering Model for CSG Ceramics 141 t -po=0.56 3 - r=1350°C 1 n 0 0.05 0.1 0.15 0.2 0.25 0.3 Volume Fraction of Sol-Gel Matrix, (Vm) Figure 6.3-5. Coefficient B of the S M C S G as a function of volume fraction of sol-gel matrix of composites The coefficient B was calculated using Eq.(6.2-23) and data from Figures (6.3-6) to (6.3-7), given by (refer to Appendix I for details): B = (3.15p 0 2 - 7.145p 0 + 4.00)- (23.42Vm 2 + 8.269Vm + 0.2085) (6.3-3) 0.9 vm T=1300°C 0.24 0.85 " ' 'Z-- 0 °-14 sity, 0.8 sity, , 0.10 c 0) Q 0.75 _ _ . o 0.05 / / ' . . . . • • - • ' > lati 0.7 .-•-^  0) rr 0.65 / - •6 ' ' 0.6 • i I I 0 40 80 120 160 200 Sintering Time, (min) Figure 6.3-6. The relative densities of AI2O3/AI2O3 sol-gel composites with different initial relative densities sintered isothermally at 1350°C as a function of sintering time. The lines were calculated using S M C S G and the points are experimental data. Chapter 6 Sintering Model for CSG Ceramics 142 -1.5 -2 g -2.5 -3.5 -4 vm • 0.24 o 0.14 • 0.1 o 0.05 ^^^X^ T=1350°C 2.5 5.5 3 3.5 4 4.5 5 Ln(t) Figure 6.3-7. A plot of Ln(p — p0) as a function of Ln(t); experimental data are same as in Fig.6.3-6. NO l O © O CL fN 0.55 0.6 0.65 0.7 Initial Relative Density of CSG, p 0 0.75 Figure 6.3-8 Coefficient B as a function of the initial relative density of C S G . For composites that do not contain any sol-gel matrix, Vm = 0, and Equations (6.3-3) and (6.3-2) reduce to: B = 0.2085(3.15p 0 2 - 7 . 1 4 5 p 0 +4.00) (6.3-4) b = b0 = 0.637 (6.3-5) Chapter 6 Sintering Model for CSG Ceramics 143 This means that due to the small value coefficient B, densification rate is reduced. According to Eq(6.2-20), viscosity of this hypothetical C S G reaches the maximum value. When the volume fraction of sol-gel matrix increases, the viscosity of C S G decreases and coefficient B increases, and thus sintering accelerates. When the composites contain 100% sol-gel matrix, Vm = 1, and the solutions of Equations (6.3-2) and (6.3-3) are given by: In this case, B reaches maximum and b minimum. The lowest viscosity of the composites results in high densification rate and low sintering temperature. 6.3.2 Validation of Sintering Model for C S G The first step to verify the model is to show that the coefficients b and B are independent of the sintering temperature. To prove this, the volume fraction of the AI2O3 sol-gel matrix of alumina/alumina composite was varied from 0.05 to 0.24, the samples were preheated at 550 °C for 1 hour, and then sintered at 1250, 1300, and 1450°C for different times. A l l initial relative densities of the samples were 0.56. The experimental data (points) and model calculation results (lines) are plotted in Figures (6.3-9) to (6.3-11). Figure 6.3-9 and Figure 6.3-10, in which the coefficients B and b in Eq.(6.2-23) remain constant for both sintering temperatures (as seen table 6.3-3), show that the sintering model for C S G ( S M C S G ) fits well to the experimental data. This also indicates that coefficients of the S M C S G are not affected by the sintering temperature at 1250°C to 1300°C. However, Figure 6.3-11 shows that calculated relative density (using b and B as in Table 6.3-3) is lower than that measured in experiments. Apparently, the coefficients b and B are not independent of B = 31.89(3.15p o 2 - 7 . 1 4 5 p o +4.0o) (: maximum value) (6.3-6) b = 0.066 (6.3-7) Chapter 6 Sintering Model for CSG Ceramics 144 sintering temperature T at, or possibly above 1450 °C. It is expected that, in this case, the mass transport is not only limited to small grains of the sol-gel matrix, but also proceeds through grain boundary diffusion, volume diffusion, and surface diffusion within the larger grains of alumina filler. Therefore, the sintering model for C S G is not valid at sintering temperature 1450°C. 0.8 -£0 .75 0.7 'SS c o> D 0> I 0.65 4) CC 0.6 0.55 7 = 1250 °C p o=0.56 vm • 0.24 0.14 ...»0.10 .,-0"' _ , P 0.05 " _..-*""""'" / / / /'•• 20 40 60 80 100 120 140 Sintering Time, (min) Figure 6.3-9. The relative density of AI2O3/AI2O3 sol-gel composites sintered isothermally at 1250°C as a function of sintering time. The lines were calculated using S M C S G and the points are experimental data. 0.75 5K 0.7 w c a> a Qi > « J OJ rr 0.65 0.6 0.55 T = 1300°C po=0.56 ^__j0.24 ___.--a0.14 .JD-""" .^ 0.10 -0 0.05 / s / ' ••' • '"" //•'' 0 20 40 60 80 100 120 140 Sintering Time, (min) Figure 6.3-10. The relative densities of AI2O3/AI2O3 sol-gel composites sintered at 1300°C as a function of sintering time. The lines were calculated using S M C S G and the points are experimental data. Chapter 6 Sintering Model for CSG Ceramics 145 1 -i 0.9 -•*—< Densi 0.8 -> 0.7 -(0 0) QC 0.6 -0.5 7=1450°C p 0=0.56 Vm=0.24 — Calculation • Experimental Data 50 100 150 200 Sintering Time, (min) 250 Figure 6.3-11. The relative densities of AI2O3/AI2O3 sol-gel composites sintered at 1450°C as a function of sintering time. The line was calculated using S M C S G and the points are experimental data. To further verify the infuence of volume fraction of matrix on coefficients B and b, a 1 mm thick f i lm of alumina gel with 2 v o l % of OC-AI2O3 was reheated at 550°C for 1 hour, and then sintered isothermally at 1200°C for 30 min. Figure 6.3-12 shows that calculations using the S M C S G model and the experimental points agree very well . 1 T 0.9 -0.8 -& in c 0.7 -0) 0 0 _> 0.6 -n o> cc 0.5 -0.4 -0.3 -20 40 60 80 S i n t e r i n g T i m e , ( m i n ) 100 120 Figure 6.3-12. The relative densities of alumina gel doped with 2wt% OC-AI2O3 sintered at isothermal 1200°C as a function of sintering time. The line was calculated using S M C S G and the points are experimental data. Chapter 6 Sintering Model for CSG Ceramics 146 0) o > « Q) CC 1 0.9 | 0.3 0.7 0.6 0.3 6>° -7 . . J O — - ° -' m - ^ 0 . 9 8 . . 0 - 0 . 8 - O - - 0 . 5 II I I I - • - 0 i 0 20 100 120 40 60 80 Sintering Time, (min) Figure 6.3-13. The relative densities of composites with different sol-gel matrix contents sintered at isothermal 1200°C as a function of sintering time. The relative density of sol-gel composites with sol-gel matrix content from 0.98 to 0 as a function of sintering time was calculated using the S M C S G and plotted in Figure 6.3-13. The composite without sol-gel matrix showed very little sintering when heat treated isothermally at 1200°C, while the composite with 98 v o l % of sol-gel matrix was fully dense after 2 hrs. Chapter 6 Sintering Model for Composite Sol-gel 6.4 Discussion of Sintering Model for CSG (SMCSG) 147 6.4.1 Analysis of S M C S G The densification rate ( p = — ) of C S G can be obtained from the sintering model for dt C S G in Eq.(6.2-21): p = Bbtb'1 expl Q_ kT + =- exp kT2 Q_ kT dt dT where — is the heating rate. It is evident that the densification rate of C S G is determined by dt sintering time, sintering temperature, volume fraction of sol-gel matrix, initial density of C S G , and heating rate. The effects of these variables on densification rate of C S G is analyzed below. Figure 6.4-1 shows the influences of heating rate on the densification rate of C S G samples at a constant volume fraction of sol-gel matrix (Vm = 0.24) and initial relative density of C S G (p0 = 0.56). It is assumed that the C S G samples are sintered with different heating rates, starting from T = 1000°C (i.e. the samples are "dropped" into a furnace at 1000°C giving an initial heating rate at 1000°C/min). It is shown that all of the densification rates drop below 0.007/min from 0.04/min after 2 min of sintering. Subsequently, the heating rate drops to a "typical" heating rate which is used for heating the sample to 1350°C, i.e. l-25°C/min. The densification rate is approximately constant at a heating rate of 5°C/min. For these conditions, a decrease of the densification rate due to decrease of surface area of powder is approximately balanced by an increase of densification rate due to increasing sintering temperature. The densification rate increases quickly with increasing sintering time for heating rate above 10°C/min. Chapter 6 Sintering Model for Composite Sol-gel 148 0 "I 1 n 1 1 1 1 0 5 10 15 20 25 30 Sintering Time, (min) Figure 6.4-1. The densification rates of C S G as a function of sintering time for different heating rate at constant initial density and volume fraction of sol-gel matrix. Figure 6.4-2 shows the densification rate as a function of heating rate for C S G with different volume fractions of sol-gel matrix, when samples are put directly into a furnace at 1000°C for 2 min. The densification rate of C S G increases with increasing volume fraction of sol-gel matrix at the same heating rate. For higher sol-gel matrix content in C S G , the densification rate increases faster with increasing heating rate. Chapter 6 Sintering Model for Composite Sol-gel 149 o o 0.7 0.6 0.5 & 0.4 0.3 4> ts cc c o "•3 a u £ 0.2 w c v Q 0.1 -0.24 po=0.56 t = 2 min O-0OOOOO' O-OO-O-OO-<3-o-oo-ooooo- o-o-oo-o- o-o-o 10 15 20 Heating Rate, (°C/min) 25 30 Figure 6.4-2. The densification rate of C S G as a function of heating rate for different sol-gel matrix content at a constant initial density after the samples were sintered from 1000°C for 2 min. 0 -I 1 1 1 1 1000 1100 1200 1300 1400 Sintering Temperature, (°C) Figure 6.4-3. The densification rate of C S G as a function of sintering temperature for different sol-gel matrix content at constant initial density and heating rate. Figure 6.4-3 shows the densification rate of the C S G with different sol-gel matrix content as a function of sintering temperature at a constant heating rate (5 °C/min) and initial relative density of C S G (0.56). The densification rate increases with increasing sintering temperature Chapter 6 Sintering Model for Composite Sol-gel 150 after dropping between 1000 °C to 1040 °C because the samples are placed directly into a furnace at 1000°C. The sample with higher sol-gel matrix content retains a higher densification rate at a constant heating rate. Figure 6.4-4 shows the influence of the initial relative density on densification rate of C S G at a constant volume fraction of sol-gel matrix (Vm = 0.24) and heating rate (5 °C/min). The results indicate that C S G samples with lower initial densities have higher densification rates. 3 c E o t 2 (0 CC c o ™ O I c 0) Q Heating Rate: 5 °C/min Po Vm=0.24 0.8 \ \ 0.56 0.4 0 1 2 3 4 5 Sintering Time, (min) Figure 6.4-4. The densification rate of C S G as a function of sintering time for different initial densities at constant heating rate and volume fraction of sol-gel matrix. dT For isothermal sintering processes, the heating rate — = 0 ; and, Eq.(6.4-1) reduces to: dt p = Bbtb~1 exp Q_ kT (6.4-2) A plot of densification rates of composite sol-gel ceramics sintered at the isothermal condition of 1350°C as a function of sintering time is shown in Figure 6.4-5. The results indicate that densification rates decrease with increasing sintering time. The increase of volume fraction of sol-gel matrix results in an increase of densification rates of C S G . Chapter 6 Sintering Model for Composite Sol-gel 151 0.2 j E 0.15 -V— QT •*-> « ? 0.1-o ^ ro o »^  1 0.05 -a 0 • 0 0.1 0.2 0.3 0.4 0.5 Sintering Time, (min) Figure 6.4-5. The densification rate of C S G as a function of sintering time for different sol-gel matrix content at constant initial density and constant sintering temperature. In accordance with the discussion above, the influence of parameters of Eq.(6.4-1) on the densification rate are summarized in Table 6.4-1: Table 6.4-1. Influence of parameters of Eq.(6.4-1) on densification rate of C S G Parameter of Eq . (6.4-1) Densification Rate (p ) Sintering Temperature (T) rr It Sintering Time, (r) rr II Initial density of C S G , ( p 0 ) rr Ii Volume Fraction of Sol-Gel Matrix, (Vm) ir rr Heating Rate, ( ) dt rr rr Figure 6.4-6 shows the relative density of C S G as a function of sintering temperature, with different heating rates at constant volume fraction of sol-gel matrix (0.24) and initial relative density (0.56). Although the densification rate of C S G is lower at lower heating rates, lower heating rates result in higher relative density of C S G . This is because the sintering time is very short (14 min) for higher heating rate (25 °C/min) when the C S G is sintered from 1000 °C Chapter 6 Sintering Model for Composite Sol-gel 152 to 1350°C, so that there is not enough time for particle rearrangement and repacking, or micropore collapse during initial and intermediate sintering stages. 0.95 o (/> o o If in c 0) Q o > ra o CC 0.85 0.75 0.65 0.55 Heating Rate, (°C/min) VM=0.24 — 1 p 0=0.56 / 5 10 — 25 1000 1100 1200 1300 Sinter ing T e m p e r a t u r e , ( ° C ) 1400 Figure 6.4-6. The relative density of C S G as a function of sintering temperature for different heating rates at constant initial density and volume fraction of sol-gel matrix. 0.8 O 0.75 o o >. 0.7 '5 c 0 ) Q a> > 0.65 H ra EC 0.6 0.55 vm = Heating Rate: 5 °C/min 0.24 P O=o.56 y 0.14 0.1 — 0.05 1000 1100 1200 1300 Sintering Temperature, (°C) 1400 Figure 6.4-7. The relative density of C S G as a function of sintering temperature for different volume fractions of sol-gel matrix at constant initial density and heating rate. Figure 6.4-7 shows the relative density of C S G as a function of sintering temperature with different volume fractions of sol-gel matrix at constant heating rate (5 °C/min) and initial relative density (0.56). Higher volume fraction of sol-gel matrix results in the higher relative density of C S G sintered from 1000 °C to 1350°C. This result indicates that the sol-gel matrix of Chapter 6 Sintering Model for Composite Sol-gel 153 C S G brings about a sintering driving force for densification. Figure 6.4-8 shows the relative density of C S G as a function of sintering time with different initial relative densities at constant heating rate (5 °C/min) and volume fraction of sol-gel matrix (0.24). The C S G with the highest initial relative density (0.8) requires the shortest time (100 min) to be sintered to 86% of theoretical density. 0.9 (3 </) £ 0.8 o & 1 0.7 0) T3 W 0.6 m a> CC 0.5 0.4 Heating Rate: 5 °C/min Vm=0.24 Po 0.8 0.56 0.4 20 40 60 Sintering Time, (min) 80 100 Figure 6.4-8. The relative density of C S G as a function of sintering time for different initial densities at constant volume fraction of sol-gel matrix and constant heating rate. 6.4-2: The influences of parameter on the relative density of C S G can be summarized in Table Table 6.4-2. Influences of parameters of Eq.(6.2-21) on relative density of C S G Parameter of Eq . (6.2-21) Relative Density ( p ) Sintering Temperature (7) tl 11 Sintering Time, (t) tl It Initial density of C S G , (p0) tl tl Volume Fraction of Sol-Gel Matrix, (V m ) 11 tl Heating Rate, ( ) dt tl It Chapter 6 Sintering Model for Composite Sol-gel 154 6.4.2 Micros t ructure of C S G at Intermediate Sintering Stage The S E M morphology of the composite sol-gel ceramics shows that the particle size of Figure 6.4-9. S E M morphology of AI2O3/AI2O3 sol-gel composites with 14vol% sol-gel matrix sintered at 1350°C for 10 min Figure 6.4-10. S E M morphology of AI2O3/AI2O3 sol-gel composites with 14vol% sol-gel matrix sintered at 1350°C for 30 min Chapter 6 Sintering Model for Composite Sol-gel 155 Figure 6.4-11. S E M morphology of AI2O3/AI2O3 sol-gel composites with 14vol% sol-gel matrix sintered at 1350°C for 60 min Figure 6.4-12. S E M morphology of AI2O3/AI2O3 sol-gel composites with 14vol% sol-gel matrix sintered at 1350°C for 120 min This result confirms the basic assumption of the S M C S G that the ceramic filler was considered as composed of incompressible and non-sintering inclusions. The micrographs show Chapter 6 Sintering Model for Composite Sol-gel . 156 that the micro-pores decrease with increasing sintering time during isothermal sintering of the samples at 1350°C. This is because the sol-gel matrix, sintered at low temperature, wi l l experience a large shrinkage to drive ceramic particle filler repacking, rearrangement, and collapse of the micropores. When the ceramic particles form a continuous network, the steric force of the ceramic particles is large enough to prevent further shrinkage of the sol-gel matrix and may cause sintering damage, such as cracks. If the initial densification rate is very fast, there is not enough time for ceramic fillers to re-pack and rearrange. Therefore, a lot of micropores still exist in the composite structure, acting like a hard agglomeration. This impacts on the densification of the composite sol-gel ceramics in the final sintering stage. The results of the analysis of S M C S G (Section 6.4.2) have shown that the relative density of C S G sintered by a slower heating rate is higher than that by the faster heating rate. These results have been confirmed by experimental data listed in Table 6.4-3, which gives the density of two identical samples with different initial densities, sintered isothermally at 1350°C. The samples with higher initial relative density showed higher final relative density than that of the samples with lower initial relative density. Table 6.4-3. Relative density as a function of initial relative density of C S G sintered isothermally at 1350°C for 3 hrs. vm 0.24 0.24 0.14 0.14 0.10 0.10 0.05 0.05 Po 0.56 0.70 0.56 0.67 0.56 0.65 0.56 0.62 p 0.826 0.848 0.797 0.822 0.760 0.782 0.717 0.744 Chapter 6 Sintering Model for CSG Ceramics 6.5 Final Stage of Sintering CSG 157 During sintering, the nano-sized grains in the sol-gel matrix grow and coalesce to become small grains located at the triple junctions at the grain boundaries between larger ceramic particles of the filler phase, as illustrated in Figure 6.5-1. This result has been confirmed by X -ray map in Figure 5.2-12. Figure 6.5-1. Morphology of AI2O3/AI2O3 sol-gel composites with 14vol% sol-gel matrix sintered at 1450°C for 30 min. The smaller grains of sol-gel matrix are located at the triple junctions and/or grain boundaries of large ceramic particles. In the final sintering stage, as a consequence of densification, the pores become isolated, mainly at three or four-grain corners. Some pores lie on grain boundaries, depending on behavior of grain growth. Compared with the initial and intermediate stage, the final sintering stage is a slow process. Chapter 6 Sintering Model for CSG Ceramics 158 The model of grain growth considers the movement of a single grain boundary in a pure, dense material [66]. There is a free energy difference AU across a curved grain boundary having a surface energy / a n d principal radii of curvature ry and ry. AU=yVn r, (6.5-1) The molar volume of atoms moving across the boundary is V„. This free energy difference provides the driving force for the boundary to move toward its center of curvature. The rate of boundary movement is proportional to the curvature, thus inversely proportional to the average grain size, G , and proportional to the ability of the atoms to cross the grain boundary, Dgb. The rate of grain growth is then [66]: (6.5-2) dt G and G2-G0 2oct (6.5-3) Figures 6.5-1, 6.5-3 to 6.5-5 are S E M photographs of the microstructure of AI2O3/AI2O3 composite sol-gel ceramics sintered isothermally at 1450°C for different times. The grain size was measured by linear intersection methods and plotted according to Eq.(6.5-3) in Figure 6.5-2. The plot shows two distinctive regions. It is believed that while the smaller grains of sol-gel matrix grow, at the same time the larger alumina particles engulf the majority of these smaller grains, as illustrated in region I of Figure 6.3-3. A t this time of about 100 min, the densification of the composite reaches 99%. When the sintering time is longer than 120 min, the diffusion and grain growth occur predominantly between the larger alumina particles, thus the initial grain size G0 = 0.31 Jim in Eq.(6.3-3) should be modified, as shown in region II of Figure 6.3-3. Chapter 6 Sintering Model for CSG Ceramics 159 2.5 | , 5 O C3 1 I CO 0.5 j - II I Alumina/Alumina CSG T = 1450°C 1 G 0 = 0.31 nm 100 200 300 400 Sintering Time, (min) 500 Figure 6.5-2. G2 - Gg2 as a function of sintering time, where G is average grain size and GQ (0.31 (im) is average initial grain size. 033655 £9KV xEQ.BK 1.56utn Figure 6.5-3. Morphology of Al20 3 /Al 2 03 sol-gel composites with 14vol% sol-gel matrix sintered at 1450°C for 50 min. It is shown that the content of the smaller grains of the sol-gel matrix has decreased, as compared to F ig . 6.5-1. Chapter 6 Sintering Model for CSG Ceramics Figure 6.5-4. Morphology of AI2O3/AI2O3 sol-gel composites with 14vol% sol-gel matrix sintered at 1450°C for 100 min. The content of the smaller grains of sol-gel matrix further decreases. Figure 6.5-5. Morphology of AI2O3/AI2O3 sol-gel composites with 14vol% sol-gel matrix sintered at 1450°C for 420 min. The smaller grains of sol-gel matrix disappear and the larger grains grow larger. Chapter 6 Sintering Model for CSG Ceramics 6.6 Summary 161 1. The sintering processes of alumina gels have been investigated. The alumina gel doped with 2vol% a - A l 2 0 3 as seeds was sintered to 97% densification at a sintering temperature of 1200°C; the linear shrinkage of alumina gel was near 30%. 2. The semi-empirical model sintering model for C S G ( S M C S G ) was proposed based on the viscous sintering model for the sol-gel matrix with rigid inclusions. The model predicts variation of density p for C S G , as a function of time and temperature according to the general expression: r n \ p = po+ Bt exp Q_ kT where p0 is the initial relative density, B and b are parameters depending on pQ and Vm (volume fraction of sol-gel phase in C S G composites) 3. The coefficients b and B of the S M C S G model were calculated from experimental data where the samples of composite sol-gel ceramics were sintered isothermally at 1350°C. The following equations were determined through best fit of S M C S G to experimental data: b = 0.637 exp(-2.265V m ) B = (3.15p 0 2 - 7 . 1 4 5 p 0 +4.00)(23.42V m 2 + 8.269Vm +0.2085) where Vm is the volume fraction of sol-gel matrix, and p0 is the initial relative density of the composites. 4. When the sintering temperature is below 1350°C for AI2O3/AI2O3 system, the S M C S G model fits the experimental data very well . However, when sintering temperature is 1450°C, the coefficients B and b are no longer independent of the sintering temperature T. The Chapter 6 Sintering Model for CSG Ceramics 162 verification results indicate that S M C S G model is suitable for the entire range of volume fraction of sol-gel matrix Vm = 0 to 1. 5. The sintering stress was estimated using the S M C S G model. The results show that samples with a higher content of sol-gel matrix produces higher sintering stresses. Heating rate and schedules that avoid excessive sintering stress were explored. 6. The S M C S G model was used to optimize the parameters of the sintering process, to eliminate the cracks caused by sintering stresses and to provide the best method to sinter the sol-gel composite to full density. 7. The microstructures of composite sol-gel ceramics were studied using S E M during the initial and intermediate sintering stages. The occurrence of micropores within composites sintered isothermally at 1350°C decreases with increasing sintering time and is accompanied by grain growth. 8. The grain growth in the final sintering stage of the composites was well simulated using the model: G2 -Gg2 °= t. A grain growth discontinuity is observed when C S G approaches full density. Chapter 7 Summary and Conclusions 163 CHAPTER 7 SUMMARY AND CONCLUSIONS 7.1 Summary Hydrated alumina sols are known to have high reactivity due to their small particle size, generally less than 10 nm at p H blow 4.5. The sols have therefore been used in the well known sol-gel process to produce ceramic bodies and coatings after suitable shaping, drying and sintering operations. The basic deficiency of the sol-gel processing is the large shrinkage of the gel upon conversion to the final ceramics. Although it has been proposed to disperse ceramic fillers into a sol to develop the composite sol-gel (CSG) ceramics and coatings, there are many problems still limiting the applications of C S G technology, such as agglomerations, cracks, sintering stresses, delamination of coatings, and excessive porosity of the composites sintered at low temperature. Thus, an understanding of the mechanisms of dispersion, gelation, drying, sinterability, and interfacial bonding of C S G could markedly facilitate the development of optimum C S G processes. The general research objective of this work was to study and develop an understanding of the mechanisms of dispersion, gelation, drying, interfacial bonding and densification of C S G ceramics. The general technological objective was to fabricate high performance C S G ceramics and to produce novel non-permeable, adherent C S G coatings on stainless steel and other metallic alloy substrates for high temperature corrosion and wear protection. In order to approach these goals, a series of investigations were conducted to develop the C S G technologies and to fabricate the bulk ceramics and coatings. These investigations were divided into four categories: (1) basic Chapter 7 Summary and Conclusions 164 CSG processes, (2) alumina/alumina and alumina/zirconia CSG ceramics, (3) SiC-alumina CSG ceramics, and (4) CSG ceramic coatings. In investigating the CSG processes, the properties of alumina sol were studied by measuring viscosity, conductivity, ionic strength, and pH of the sol. The dispersion and stability of ceramic particles in alumina sol were investigated by measuring particle size distributions, zeta potential, and calculating their interaction energy according to D L V O theory. In the alumina/alumina and alumina/zirconia CSG ceramics, hydrated alumina sols were used as a sintering and dispersion additive. The composites were fabricated through dispersing ceramic fillers (i.e. calcined alumina and stabilized zirconia) into alumina sol, gelcasting, drying, and sintering. The correlations between mechanical properties and microstructures of composites were studied by measuring microhardness, porosity, and SEM morphology. An additional sintering additive (MgO) was used to enhance the densification kinetics of the composites. The elemental M g distribution in composites was determined using X-ray mapping and EDX analysis. SiC and calcined alumina were dispersed into alumina sol to fabricate high performance SiC-alumina CSG ceramics containing from 5 vol% to 60 vol % SiC. The composites were sintered using a tungsten element furnace in argon, at 1600 °C to 1900 °C for 1 hour. The correlations between mechanical properties and microstructures of composites were studied by measuring microhardness, porosity, T E M , EDX, and SEM morphology. The grain boundaries of SiC/SiC and SiC/Al 2 03 were analyzed by measuring composition and morphology under SEM. The CSG ceramic coatings were deposited by dipping, spraying, and electrophoretic deposition. A post deposition treatment was carried out to improve microstructures, and enhance densities and mechanical properties of CSG coatings. The correlations between mechanical Chapter 7 Summary and Conclusions 165 properties and microstructures of the coatings were investigated by measuring the surface hardness, bonding strength between coatings and substrates, gas permeability, and S E M studies. A generalized model to predict the sinterability of C S G , to analyze sintering stress of C S G and to optimize the processing technologies of C S G has been proposed. The sintering model has been based on the viscous sintering theory and mechanism of mass transfer. The model was verified by the experimental results, and used to design the optimum processing conditions for C S G . 7.2 Conclusions Based on the experimental and theoretical work carried out in this study for the processes, composites, and coatings of the C S G , the following conclusions are drawn: 1. B y analysis of preparing alumina sol processing, it has been determined that p H value and temperature dominate the time of sol preparation. 2. Viscosities of alumina sols increase slowly with increasing concentration at p H = 4 for concentrations below 2.8 M . However, the viscosity increased rapidly when the sol concentration was greater than 3 M . When the p H exceeded 5.6, the viscosity rapidly increases from 1.2 cP to more than 30 cP. The electrical conductivity of 1 M alumina sols decreased sharply above pH=2 because the simple species, such as A l ( O H ) 2 + , A l ( O H ) 2 + , A l i 3 0 4 ( O H ) 2 4 ( H 2 0) i2 7 + , dimer and/or trimer in solution may condense to large polyvalent ions. 3. The zeta potential of AI2O3 in 1 M alumina sol increases slightly, whereas that of S i C reverses its sign over a wide range of p H , as compared to that of AI2O3 in water. It is believed that a coating layer on the S i C particles has been created by alumina sol clusters. The stability of ceramic particles in alumina sol was analyzed and the interaction energy of Chapter 7 Summary and Conclusions \ 66 system was estimated using D L V O theory. Dispersion of alumina and S i C particles has been substantially improved in alumina sol, as compared to pure water of similar acidity, i.e. the average agglomerate size was decreased by at least 50%. Based on these findings, it is expected that the alumina sol clusters can be used as an efficient, clean dispersant for single-phase and composite ceramics. 4. Alumina composite sol-gel (CSG) ceramics have been produced by dispersing calcined alumina powder in alumina sol. It has been found that the alumina sol acts as a dispersant and sintering accelerator for the calcined alumina. The sol-gel matrix phase of C S G bonded calcined alumina to form a strong 3-D network. The microhardness and porosity of C S G depends on the sintering temperature and sol-gel phase content. When sol-gel matrix phase content is about 14 vo l%, the microhardness reaches the maximum value (17.5 GPa) for pure alumina-alumina C S G . Porosity of the composites decreases with increasing sol-gel matrix phase content. The microhardness of composite sol-gel containing M g O sintering additive further increases above 20 G P a and porosity decreases to less than l v o l % , for a sintering temperature of 1400 °C. The sintering kinetics and final microstructure of C S G is strongly affected by M g O content. 5. The alumina/silicon carbide C S G with high S i C content (50 vol%) have been pressureless sintered at 1850 °C to 98% density and microhardness of 22.9GPa, using active sol-gel alumina additive. The additive was introduced to the system through dispersion of S i C and A 1 2 0 3 grains in alumina sol, which aids deagglomeration of the components. It is proposed that the sol-gel phase acts as a sintering additive through formation of a continuous intergranular f i lm around both alumina and silicon carbide grains. The active intergranular phase serves as a fast diffusion path during sintering of the composite. In particular, the intergranular f i lm of sol-gel originating alumina offers a fast mass transfer path to assist in Chapter 7 Summary and Conclusions 167 sintering of S i C / S i C grains. Absence of such fast diffusion path in water - dispersed S i C -AI2O3 results in poor densification of the composites for S i C content above 10 v o l % 6. The novel composite sol-gel (CSG) technology has been successfully used to fabricate non-permeable, crack-free, up to 600 pm thick ceramic coatings on stainless steel, titanium, and aluminum alloy substrates. The thickness of C S G coatings was controlled by the viscosity of the slurry and withdrawal speed of the substrate in dip coating processing. 7. The sol-gel sealing treatments were carried out to improve the microstructure and protective characteristics of C S G coatings in corrosive gaseous and high temperature environments. The effects of post deposition were investigated through measuring gas permeability of C S G coatings as a function of concentration of alumina sol and multiple infiltration. The optimum concentration of alumina sol for post deposition treatment is 1.5 M . The significant reduction (over 90%) of gas permeability of C S G coatings has been achieved by multiple infiltration. 8. The correlations between the processing methods, microstructure, and mechanical properties of C S G coatings were investigated by studying coatings' morphology, and measuring their bonding strength and microhardness. The best C S G coatings on metallic substrates were successfully fabricated by combination of C S G and chemical bonding technologies. The bonding strength between the substrates and coatings is about 42 M P a . The surface microhardness of the coatings is about 6.5 GPa, after sintering at 550°C. 9. The phase transformation and sintering processes of alumina gel have been investigated. The alumina gel doped with 2 v o l % a - A l 2 0 3 seeds was sintered to 97% density at 1200°C. The linear shrinkage of alumina gel was approximately 30%. 10. The sintering model for C S G ( S M C S G ) was built based on the viscous mass transport process. The model describes variations of relative density of C S G p as a function of sintering time t and temperature T: Chapter 7 Summary and Conclusions 168 ( Q_ kT p = p0+ Btb exp -J where b and B are coefficients depended On the volume fraction of sol-gel matrix and initial density, and Q is activation energy for viscous flow. 11. The S M C S G was validated by the experimental data. For sintering temperatures below 1350°C in AI2O3/AI2O3 system, the S M C S G model fitted very well the experimental data. However, when sintering temperature is 1450°C, the coefficients B and b are no longer independent of the sintering temperature T. The coefficients b, B, and Q of the S M C S G model were obtained from experimental data for the samples of composite sol-gel ceramics sintered isothermally. Q = 58 kJ/mole and is comparable to the activation energy for glass sintering [177]. 12. The hydrostatic sintering stress was analyzed by elastic constitutive equation and calculated from the S M C S G model. The sintering stress was in the range of 0.1-10 M P a . The results show that the samples with higher content of sol-gel matrix produces the higher sintering stresses. The S M C S G model was used to optimize C S G processing. The optimized sintering processes avoid crack formation due to excessive sintering stresses. 13. The microstructures of composite sol-gel ceramics were studied using S E M after initial and intermediate sintering stages. The content of micropores in composites sintered isothermally at 1350°C decreases with sintering time, but the grain size remains approximately constant. The grain growth in the final sintering stage of C G S conformed to a parabolic model within two distinctive regions. Chapter 8 Recommendations for Future Work 169 C H A P T E R 8 RECOMMENDATIONS FOR FUTURE WORK Although is has been shown that the sol-gel matrix enhances the sinterability of C S G , the kinetics of sintering processes have not been clearly resolved. Further investigations, both theoretical and experimental, are needed to quantify the influence of sol-gel matrix on the correlation between the sintering stresses and sinterability of C S G . The sintering model for C S G should be extended to other composite sol-gel system, such as ceramic filler/zirconia sol or/and ceramic filler/titania sol, to enhance fundamental understanding of sintering mechanisms and the mass transport of C S G . The further investigation of structure of alumina sol clusters and the transition from alumina gel to Y - A 1 2 0 3 should be carried out by using N M R . The C S G technology can be used to develop high toughness nano-size alumina/zirconia composites. In one of the promising variants of the methods, alumina and zirconia powders (0-40vol%) may be dispersed in zirconia sol. The samples should be sintered at 1400-1500°C after gelcasting and drying. The grains of alumina powder coated by zirconia gel w i l l not grow during sintering at high temperature, preserving nano-size of the grains. The zirconia sol-gel phase wi l l be homogeneously distributed at the grain boundaries of alumina and zirconia. It is expected that the composites would have a uniform nano-grain size, accompanied by high fracture toughness and high microhardness. The C S G technology also can be used to fabricate the SiC/alumina and TiC/alumina composite powders. 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Appendix I 179 APPENDIX I PARAMETERS OF SINTERING MODEL FOR CSG For coefficient B in Eq(6.2-21), f{po) is a constant when all initial relative densities of the C S G are 0.56. The experimental data of the coefficient B was fit using Eq.(6.2-23) and (6.2-24), given by: / • ( 0 = f * _ l (23 .42V m 2 + 8.269Vm +0.2085) (1) C/ 2 (0.56) J In order to determine the influence of initial density on the coefficients B, the volume fraction of the AI2O3 sol-gel matrix of the alumina/alumina composite studied was from 0.05 to 0.24. The samples were preheated at 550 °C for 1 hour, and then sintered isothermally at 1250 °C for 1 hour. Finally, the samples with different initial relative densities were sintered isothermally again at 1350°C for different sintering times. For comparison to the C S G samples of the initial relative density 0.56, sintered at 1350°C, experimental data was calculated using the same procedure as that above. Figure 6.3-6 shows the relative density of sol-gel composites with different initial relative density sintered isothermally at 1350°C. Figure 6.3-7 shows a plot of Ln(p - pa) as a function of sintering time. According to Eq.(6.2-23), the coefficient B = Cf (Vm ) / 2 (p o ) was obtained using the experimental data from Figure 6.3-7. The coefficient B0.56 was calculated using Eq.(6.3-3) in which all the initial densities were 0.56. The ratio of ^ 2 ^ ° ^ is plotted as a function of initial A (0.56) relative density in Figure 6.3-8. The / 2 ( p o ) can be expressed as: fiiPo) = fi (0.56)(3.15p o 2 - 7.15p 0 + 4.00) (2) Appendix I 180 B y combining the Equations (6.2-23), (6.3-3), and (6.3-4), coefficient B can be expressed as: B = (3.15p 0 2 - 7.145p o + 4.00)(23.42Vm 2 + 8.269Vm + 0.2085) (3) Appendix II 181 APPENDIX II SINTERING STRESSES IN CSG Ceramic fillers added into sol-gel matrix cause a reduction of the densification rate of the composite. In particular, stresses develop due to differential shrinkage characteristics of the filler (treated here as heterogeneity). The stresses may cause sintering damage, such as cracks, flaws, planar arrays of void, and isolated pores [85,174]. In this section, S M C S G is used to estimate the magnitude of these stresses in composite sol-gel ceramics. 1. Constitutive Equation The radial and hoop stresses with in the spherical inner zone subject to the surface traction, cr, are given by [175]: o'r(r) = cr,e(r) = el(r) = o (1) The corresponding stresses in the outer zone: c>'r'(r)=r>a3/r3 (2) <r"(r) = ej''(r) = -e>ai/2r3 (3) where a is the inner zone radius and r is distance from the center of the zone, as illustrated in Figure 1. The volume change rate, V , is related to relative density change rate p , V p u= ( 4 ) V p So the densification rate of the composite is p / p = - 3 e (5) where ec the strain rate of composite. Appendix II 182 Or Figure 1. A schematic of a spherical inclusion indicating the coordinate system. The constitutive equation for elastic strain is [175] ex=ef +E'1[Gx-v(cry+C7z)] (6) where ex is the strain in the direction, £/ is the free strain, E is Young's modulus, v is Poisson's ratio, and ax, <Jy, and <7Z are the stresses in the x, y, and z directions, respectively. In a thermal stress problem, Ef = aAT, where a is the liner thermal expansion coefficient and AT is the temperature change. The essential point is that the observed strain (e*) is a linear combination of free strain (£/) and the strain caused by stresses. The analogous result for a viscous material is [175]: ex =ef +(3t1r1[ax-(l/2)(<7y +az)] (7) where ex and ef are strain rates; 77 is the viscosity and v = 1/2, indicating that the material is incompressible. The sintering matrix is not incompressible because it contain pores, so Eq . (7) becomes: Appendix II 183 =£f +Em-\ax-vm(oy+Gz)} (8) where Em is the "apparent Young's modulus" and v m is Poisson ration of porous matrix. Em represents the viscous response of porous material to a uniaxial stress, so it w i l l be called the uniaxial viscosity of porous matrix. According to this model, i f 0 is the load-bearing fraction of the cross sectional area of porous body, then Em = 3rj(j) and v m = $11 [175]. The strain rate depends on the magnitude of the pressure and the bulk modulus of the body. If the applied stress is hydrostatic, ox = ay = oz = P, then Eq.(8) becomes ex =ef+P(l-2vm)/Em = ef+P/(3Km) (9) where Km is the "apparent bulk modulus" defined by Km=EJ[3(\-2vJ] (10) by analogy to the elastic bulk modulus. Km represents the viscous response of porous body to hydrostatic stress, so it w i l l be called the bulk viscosity of the porous matrix. The viscous or elastic response of an isotropic material to stress or strain can be described by two independent functions. The constitutive equation can be written in term of Em and v m in Eq.(8) or equivalent expressions as Km, the "apparent shear modulus", Gm Gm=EJ[2(l+vm)] (11) Again, Gm is not a true modulus; it represents the viscous response of the porous materials to shear stress, so it w i l l be called the shear viscosity of the porous matrix. 2. Stress on the Inclusion In spherical coordinates, the elastic constitutive equation reduces to [176]: er =ef + £ - ' ( c T r - 2 v c T e ) (12) Appendix II 184 e9=ef+E-x[ae-v{cr+ae)] (13) The free strain of the inclusion core is zero (e^ = 0). The elastic strain of the inclusion core is very small compared to the densification strain of strain of the matrix, so the inclusion core can be regarded as incompressible (i.e., its bulk modulus Ki—> °°), thus, the elastic solution is [176]: Gri=Oa =oi={\-Vi)Klefin (14) on>=l*lr'f-vk EaeJ> (15) ^=4(l /2X«/r) 3 +V , .k^ (16) Kl^lfaZyV^Klf (17) where GmE and KmE are the elastic shear and bulk modulus of the matrix, respectively. The volume fraction V, is the current value, which increases as the matrix density. The radial displacement (w) at the surface of the composite sphere is K . ( r ) / r = ( l - V J * * £ ^ / ( 3 * * ) (18) The constitution equations of the sintering matrix are in spherical coordinates er=ef+Ef-\or-2vmoe) (19) £e =ef +Em-l[og -vm(ar +ag)] (20) Since the constitutive equations and the boundary conditions of the sintering problem are analogous to those of the thermal stress problem, the solutions to the later can be directly adopted. The solutions for the sintering problem are obtained by replacing GmE and KmE with Gm and Km (defined by Eqs. (10) and (11)) and replacing the strains with the respective strain rates. Thus, the stresses in the sintering composite sphere are given by ff^d-Vi)^ (21) (ym={ialrf-Vi}KJfm (22) Appendix II 185 aBm=-[{\l2){alr)'+Vi}KJfm (23) ^ =[1/ (40 . ) + V,/(3^jr 1 (24) The circumferential stress in the matrix is tensile (since <0) , so cracks tend to propagate radially from the rigid inclusion [101]. The liner strain rate of the composite is derived from Eq.(18): ic=um(r)/r = (\-Vi)KJftn/(4Gm) (25) The densification rate of matrix Pm'Pm = - ( 2 e t e + e r m ) (26) Using Eqs. (19) and (20) this lead to PjPm = -Ufin-OmIKm (27) where om is the hydrostatic stress in the matrix, given by Om=(orm+2aem)/3 (28) Thus the sintering rate is affected by the hydrostatic component of the stress in the matrix, which is independent of radial position. From Eqs. (21) to (23) °m=-ViKJfm=-Vioil(\-Vi) (29) Since hydrostatic compression develops in nonsintering inclusion, the matrix must be subject to hydrostatic tension. 3. Hydrostatic Stresses Calculated by SMCSG The hydrostatic stress in the element, which represents any region of matrix, is found by applying the viscous analogy to Selsing's solution [176]: strains are replaced by strain rates and the respective modulus of the island and continuum are replaced by Km and G c . Thus, Appendix II 186 a = (e - e, ) m \ c fin ' 1 1 3tf„ 4G„ 1 1 • + -3 £ f 4 G , (30) (31) Since e . = 0 , Eq.(31) reduces to CT, = £ , 1 1 - + • s - l (32) 3K, AGC where Gc is an "apparent shear modulus" of porous C S G and K is an "apparent bulk modulus" of inclusions (ceramic filler), expressing viscosity of the matrix and inclusion, respectively. Replacing the linear strain rate £ by Eq.(5), the sintering stress can be expressed as: 3p 1 1 - + -3 £ . 4G„ (33) The hydrostatic sintering stress on the sol-gel matrix can be expressed as: CT = V, 1-V, 3p 1 1 3K, 4G„ (34) The modulus M (shear or bulk) of a porous system has been determined as [85]: f T 7 / 1 - \ \ M=M. l + -7 ( 1 - p ) (35) \ l - ( T + l ) ( l - p ) J where Y is a constant (Y ~ - 4 for AI2O3 system [85]) and Ma is the modulus at full density The sintering hydrostatic stresses on ceramic fillers and sol-gel matrix as a function of sintering time at 1350°C are plotted in Figure 2 and Figure 3 using Eqs. (6.2-21), (33), (34) and (35), respectively. The initial relative density of all composites was 0.56. The sintering stress on ceramic fillers decreases with sintering time because the densification rate continues to drop with sintering time. The stress on ceramic fillers is compressive. The increase of sintering stress with content of sol-gel phase Vm indicates that the driving force for densification increases. However, Appendix II 187 the sintering stress on sol-gel matrix is tension stress in Figure 3. The high sintering stress on sol-gel matrix during processing may cause some damage, such as cracks and/or voids [85]. The sintering stress model for AI2O3/AI2O3 system was valid at sintering temperature between 500°C - 1400°C. In o iZ o E (0 0) O ^ 0) V) O) c *c V c '(/> 0 -1 -2 -3 -4 -5 -6 -7 -8 -9 -10 / / If vm — 0.24 I T=1350°C — 0.14 0.1 ' p o =0 .56 — 0.05 0.1 0.2 0.3 0.4 Sintering Time, (min) 0.5 Figure 2. The sintering stress on ceramic fillers for C S G as a function of sintering time for different sol-gel matrix contents at a constant initial density and isothermal conditions. I 5 H •3 9 4 o <f) _ c * 0 £ 3 in g, in a K 2 u> c S 1 c T = 1350°C p 0 =0.56 vm 0.24 \ — 0.14 \ » V 0.1 V \ \ — 0.05 \ 0.1 0.2 0.3 Sintering Time, (min) 0.4 0.5 Figure 3 The sintering stress on sol-gel matrix for C S G as a function of sintering time for different sol-gel matrix contents at a constant initial density and isothermal conditions. Appendix II 188 Figure 4 shows the influence of different initial relative densities on the sintering stresses on ceramic fillers during isothermal sintering according to the predictions of S M C S G . The results indicate that the samples with higher initial relative density experience lower sintering stresses. 0 -5 4-^ , 1 , 1 1 0 0.2 0.4 0.6 0.8 1 Sintering time, (min) Figure 4. The sintering stress on ceramic fillers for C S G as a function of sintering time for different initial densities at constant sol-gel matrix contents and isothermal conditions. 0 -0.6 -H , , 1 1 1000 1100 1200 1300 1400 Sintering Temperature, (°C) Figure 5. The sintering stress on ceramic fillers for C S G as a function of sintering temperature for different heating rates at constant sol-gel matrix content and initial density. Appendix II 189 Figure 5 shows sintering stress on ceramic fillers as a function of sintering temperature for different heating rates at a constant initial relative density (0.56) and volume fraction of sol-gel matrix (0.24). In all cases, the sintering stress reduces to minimum at 1040 °C. This is because placing the samples directly into the furnace at 1000°C is equivalent to heating the samples to 1000°C from room temperature at very fast rate (~1000°C/min). Subsequently, the sintering stress slowly increases with increasing sintering temperature. The sintering stress for higher heating rate increases faster because of the higher densification rate. The results indicate that the heating rate strongly affects the sintering stress of C S G . 4. Simulation of Optimized Sintering Experiments Assuming arbitrarily that the maximum stress that does not cause damage in composite sol-gel ceramics during sintering processing is 1.2 M P a , the optimal process parameters may be predicted using the sintering model for C S G ( S M C S G ) . Figure 2 shows the sintering stresses on ceramic fillers for C S G with different sol-gel matrix content, sintered isothermally at 1350°C as a function of sintering time. In order to keep the sintering stress below 1.2 M P a , i f a higher sol-gel matrix content is present, then a lower initial sintering temperature must be used. The S M C S G model was used to design sintering processes which maintain sintering stress below 1.2 M P a . There are two basic options for sintering the C S G : multiple-step sintering and constant heating rate sintering. The multiple-step sintering is simulated using the following method. In the first step, the maximum temperature (Ti) at which sintering stress is approximately equal to 1.2 M P a is calculated using Eq.(33). During sintering at temperature Tjy the sintering stresses is allowed to drop until 0.1 M P a at which point the sintering driving force is assumed to be too small to sinter to full densification after sintering time, t\. Subsequently, the first step is repeated until the Appendix II 190 sintering temperature reaches 1350°C. Alternatively, the initial sintering temperature and a constant heating rate, which maintains the sintering stress below 1.2 M P a , are selected and densification rate compared with the multi-step sintering. 1400 3 £ 1000 co 900 — Multiple-Step Sintering — Constant Heating Rate, (13.4C/min) 20 40 60 Sintering Time, (min) 80 Figure 6. Sintering temperature of C S G as a function of sintering time, designed for sintering stress below 1.2 M P a . Figure 6 shows the two sintering profiles as a function of sintering time. The sintering time for the second method was chosen to be the same as that of the first method, such that the resulting heating rate for the second method is 13.4 °C/min. In the multiple-step sintering route, AI2O3/AI2O3 sol-gel composite with 40 v o l % of sol-gel matrix is sintered in seven steps. The sample is first sintered isothermally at 950, 1020, 1070, 1130, 1200, 1280 °C for 5 min, respectively. Finally, the sample is sintered isothermally at 1350°C for 50 min. For the constant heating rate experiment, the sample is heated from 950 °C to 1350 °C at a heating rate of 13.4 °C/min, and then sintered isothermally at 1350°C for 50 min. The sintering stresses calculated from S M C S G for the two routes are illustrated in Figure 7. In the multiple-step sintering, the sintering stress oscillates between 0.03 and 1.2 M P a each time the Appendix II 191 sintering temperature increases. For constant heating rate, the sintering stress initially drops, and then continues to increase with increasing sintering temperature to 1350°C at a constant heating rate, and then drops to 0.04 M P a during isothermal sintering at 1350 °C for 30 min. Thus, the sintering process through the constant heating rate is unlikely to cause sintering damage. 10 20 30 40 Sintering Time, (min) 50 60 Figure 7. Predicated sintering stress for C S G as a function of sintering time for multiple-step sintering and a constant heating rate sintering 0.95 (3 8 0.85 o >» '</> g 0.75 ^ a> > «° a> 0.65 H 0.55 p O = 0 . 5 6 V m = 0 . 4 0 — Multiple-Step Sintering — Constant heating Rate,(13.4 C/min) 20 4 0 60 Sintering Time, (min) 8 0 Figure 8. Relative density of C S G as a function of sintering time for multiple-step sintering and at a constant heating rate sintering. Appendix II 192 Figure 8 shows the predicted relative densities of C S G as a function of sintering time for the two different sintering methods. In the multiple-step sintering, the relative density of C S G increases step-wise between sintering temperatures of 950 °C to 1350 °C. The relative density of C S G sintered at a constant heating rate increases continuously from 950 °C to 1350 °C. Finally, the relative density of C S G obtained in multiple-step sintering is higher than that of C S G sintered at a constant heating 

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