© Copyright 2011
Steven R. Mercurio Jr.
All Rights Reserved
EFFECT OF COPRECIPITATION OF SINTERING AIDS ON THE
MICROSTRUCTURE AND GRAIN BOUNDARY DEVELOPMENT OF SINTERED
SILICON CARBIDE
by
STEVEN R. MERCURIO JR.
A dissertation submitted to the
Graduate School-New Brunswick
Rutgers, The State University of New Jersey
In partial fulfillment of the requirements
For the degree of
Doctor of Philosophy
Graduate Program in Materials Science and Engineering
Written under the direction of
Professor Richard A. Haber
And approved by
______
______
______
______
______
New Brunswick, New Jersey
October, 2011
Abstract of the Dissertation
Effect of Coprecipitation on the Microstructure and Grain Boundary Development of
Sintered Silicon Carbide
By Steven R. Mercurio Jr.
Dissertation Director:
Richard Haber
Coprecipitation was examined as a method of introducing sintering aids into silicon carbide (SiC) as a fine, reactive coating. The improved sinterability and mixedness of coprecipitated samples, when coupled with advanced densification methods, developed fine grained SiC with varied microstructures. Coprecipitation imparted additional process control and influenced the phase, crystallinity, and properties of the grain boundaries.
A simple coprecipitation process was developed to introduce aluminum and rare earth sintering aids into SiC. Early samples yielded low densities so the process was modified to address the dispersion of the SiC particles and breakdown of agglomerates before coating. The modifications improved the densification and influenced the structure and properties.
Powders were prepared with varying weight percents of sintering aids and several rare earths in order to study the grain boundary structure and properties. Samples were densified using hot pressing and spark plasma sintering to better utilize the enhanced
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sinterability of coprecipitated powders. These samples were compared to conventional ball mill processing.
Scanning electron microscopy was utilized to examine the microstructure and determine the grain size and presence of defects. The degree of mixedness of the additives was investigated through fluorescence measurements. X-ray diffraction was used to determine the polytype and phase distribution. Selected mechanical properties were measured and compared between the different samples. Hardness was studied extensively, including analysis of load-hardness curves over a range of loads. The hardness data and indents were examined in order to explore the fracture behavior and defect effects.
Liquid phase sintered SiC prepared using coprecipitation exhibited very different phase content and crystallinity than ball milled samples. Fluorescence measurements for coprecipitated samples showed longer decay lifetimes indicating improved mixedness.
Samples with amorphous grain boundaries and triple points were developed, where XRD results displayed a lower amount of yttrium aluminum garnet than other methods. The formation of a crystalline mullite phase and absence of excess alumina were observed.
These overall results indicated the possibility of different fracture behavior.
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Acknowledgments
Zero score and 6 or so years ago, I began forth at Rutgers, this thesis, conceived in struggle and dedicated to the proposition that I could see this through. Now, I’m finished engaging in this great task which can long endure. I have come to dedicate these thank yous as a final resting place for those who gave their hard work, time, effort, and care so that this thesis may live. It is altogether fitting and proper that I should do this.
But, in a larger sense, I cannot dedicate – I cannot consecrate – I cannot hallow – this thesis. The brave men and women who struggled here with me have consecrated it far above my poor power to add or detract with thanks. The world will little note, nor long remember what I say here, but I will never forget what all my undergraduate technicians, group members (even Stevie B), friends, family, faculty, and others did here.
It is for us to be dedicated here to the excellent work which they who have helped me have thus so nobly advanced. It is rather for us to be here thanked for the great work finished – that from these hard working friends, department staff, John Yaniero, we take the increased devotion to the cause, the thesis, for which they gave their last full measure of devotion. I here highly resolve that these people shall not have worked in vain, that this thesis now has a new birth of freedom, and that it shall not perish from the earth.
Thanks to Abe Lincoln for inspiring me to write some of my acknowledgments as such.
I would like to give special thanks to my advisor, Dr. Haber, for your support, guidance, and entertaining moments and conversations throughout the years. I also have the deepest gratitude to Dr. Todd Jessen for providing me with support, additional funding, and especially your technical approach and wisdom. Your belief in this work during some of my less optimistic moments helped me see it through. I’d like to give a
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hearty thank you to Dr. Greenhut for your valuable conversations, both technical and non-technical. I’m also grateful to Dr. Niesz for serving on my committee. Thank you to
Dr. Joseph Michael of Sandia National Laboratory for the assistance with TEM sample preparation and Dr. F. Cosandey for his guidance on the TEM imaging and analysis. The technical expertise of Dr. Bill Mayo on x-ray diffraction was invaluable in moving this work forward. Thanks to the US Army Research Laboratory, Materials Center of
Excellence, and Center for Ceramics Research at Rutgers for their financial support.
I’d like to thank Dr. Mihaela Jitianu for her help and expertise that added greatly to this research. Very special thanks to Dr. Steve Miller for both his contributions to this research and my overall sanity during my time at Rutgers, Doug Slusark for being my necessary “nemesis”, and to Dan Maiorano for walking much of this long and winding path with me….We will always share a bond that can only be forged in moments when hot presses are melting before your eyes and the floor is stained green with antifreeze.
My deepest gratitude goes to my parents for raising me to always strive for learning and for giving me the traits that made this journey possible. Thanks to my little sister for the laughs, support, and the pleasure of watching you strive and achieve as well.
To Maddie, who can’t even read this yet…Thank you for raising my spirits during the tough parts, and for appearing in that dream when I was thinking of quitting and telling me to “Finish what you started.” Becky, my wonderful wife, I can never thank you enough for helping me, supporting me, and most of all, putting up with me during this journey. I would not and could not have done this without you.
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Table of Contents Abstract of Dissertation………………………………………………………………..….ii
Acknowledgments…………………………………………………………………………v
Table of Contents…………………………………………………………………...……vii
List of Tables……………………………………………………………………………xiii
List of Figures………………………………………………………………………..…..xv
Chapter 1 Introduction to Silicon Carbide…………………………………………….....1
1.1 Motivation and Background……………………………………………………1
1.2 Silicon Carbide………………………………………………………………….5
1.2.1 Crystal Structure………………………………………………………..5
1.2.2 Stacking Faults………………………...………………………………10
1.3 Green Processing of SiC………………………………...…………………….13
1.3.1 Production of SiC Powders……………………………………………13
1.3.1.1 Importance of Raw SiC Powders…………………………….....15
1.3.2 Conventional SiC Processing………………………………………….16
1.3.3 Colloidal Processing……………..……………………………………18
1.3.4 Sintering…………………………………………………………….…19
1.3.4.1 Sintering Theory………………………………………………..19
1.3.4.2 Assisted Sintering Methods…………………………………….24
1.3.4.2.1 Hot Pressing……………………………………….24
1.3.4.2.2 Spark Plasma Sintering……………………………25
1.3.4.3 Solid State Sintering…………………………...……………….26
1.3.4.4 Liquid Phase Sintering………………………………………….28
1.3.5 SiC Sintering Studies…………………………………………….……31
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1.3.5.1 Solid State Sintering Studies……………………………………31
1.3.5.2 Liquid Phase Sintering Studies……………………...………….33
Chapter 2 Method of Attack……………………...…………………………………….35
Chapter 3 Improved Coprecipitation Processing as a Method of Removing Defects….39
3.1 Introduction…………………………………………………..………………..…39
3.2 Processing – Structure – Properties Relationship………………………………40
3.2.1 Impact of Defects on Mechanical Properties………………………..….40
3.2.2 Influence of Processing on the Population of Defects………………...41
3.3 Improved Processing Methodology………………………………………….…44
3.3.1 Coprecipitation………………………………………………………...45
3.4 Procedures……………………………………………………………………....50
3.4.1 Initial Coprecipitation Procedure…………………………………..….50
3.4.1.1 Raw Materials……………………………………………………..50
3.4.1.1.1 Silicon Carbide……………………………………..50
3.4.1.1.2 Sintering Aids………………………………..……..51
3.4.2 Composition Selection………………………………………………….52
3.4.2.1 Titration Curves……………...…………………………………52
3.4.3 Coprecipitation Procedure ……………………………………………..53
3.4.4 Processing Improvements………………………………………………56
3.4.4.1 Milling Treatment………………………………………………56
3.4.5 Comparison Samples………………………………………...…………57
3.4.6 Note on Sample Naming and Nomenclature………….………………..57
3.4.7 Sintering…………………………………...……………………………58
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3.4.7.1 Hot Pressing…………………………………………………….58
3.4.7.2 Spark Plasma Sintering…………………………………………60
3.5 Characterization Techniques……………………………….…………………….62
3.5.1 BET……………………………………………………………………..62
3.5.2 FTIR…………………………………………………………………….62
3.5.3 Electrokinetic Analysis / Zeta Potential….……………………………..63
3.5.4 Density………………………………………………………………….63
3.5.5 Grinding, Cutting, Polishing……………………………………………64
3.5.6 Etching………………………………………………………………….64
3.5.7 SEM…………………………………………………………………….65
3.5.8 Grain Size and Porosity………………………………………………...66
3.5.9 Hardness………………………………………………………………..66
3.5.10 Ultrasonic Nondestructive Testing……………………………………..67
3.5.11 XRD………………………….…………………………………………68
3.5.12 TEM………………….…………………………………………………69
3.5.13 Characterizing Mixedness Changes…………………………………….69
3.6 Results and Discussion…………………………………………………………..70
3.6.1 Processing………………………………………………………………70
3.6.1.1 Coprecipitation………………………………………………….70
3.6.2 Initial Characterization………………………………………………….71
3.6.2.1 Zeta Potential……………………...……………………………71
3.6.2.2 FTIR………………………………………...……………….....74
3.6.2.3 Density………………………………………………………….77
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3.6.3 Improved Coprecipitation Process……………………………………...78
3.6.3.1 Milling……………………………………...…………………..79
3.6.3.2 Coprecipitation in Isopropanol…………………………………80
3.6.3.3 Calcination……………………………………………………...82
3.6.4 Characterization of Coprecipitated Samples after Improvements……..85
3.6.4.1 Zeta Potential…………………………………...………………85
3.6.4.2 TEM…………………………………………………………….87
3.6.4.3 BET……………………………………………………………..88
3.6.5 Comparison of Coprecipitated versus Conventional Mixing Processes..89
3.6.5.1 Hot Pressed Samples……………………………………………89
3.6.5.1.1 Density……………...………………………………90
3.6.5.1.2 SEM………………………………………………...91
3.6.5.1.3 XRD…………………………………………...……95
3.6.5.1.4 Mechanical Properties………………………………97
3.6.5.2 Spark Plasma Sintered Samples……………………………...…99
3.6.5.2.1 Density………………………...……………………99
3.6.5.2.2 Mechanical Properties……………….…………….100
3.6.6 The Use of Fluorescence Analysis to Compare Mixedness……...……102
3.6.7 EDS/TEM analysis of Grain Boundaries…………………………...…108
3.7 Summary………………………………………………………………………..111
Chapter 4 - Coprecipitation as a Method of Controlling the Microstructure of SiC…...112
4.1 Introduction……………………………………………………………………112
4.2 Microstructure Development…………………………………………………..112
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4.2.1 Grain Boundary Engineering…………………………………………..113
4.2.2 Grain Size………………………………………………………………116
4.3 Results…………………………………………………………………………..118
4.3.1 Microstructure Control………………………………………………….118
4.3.2 Microstructure…………………………………………………………..118
4.3.2.1 Influence of Mixing – Coprecipitation versus Ball Milling………...118
4.3.2.1.1 Densification Results…………………...…………………..119
4.3.2.1.2 XRD………………………………………………………..122
4.3.2.1.3 EDS Mapping in SEM……………………………………...131
4.3.2.1.4 Presence and Effect of Grain Boundary Glass………….…..134
4.3.2.2 Influence of Sintering Aid Percentage………………………….….137
4.3.2.2.1 Densification……………………………………….……….137
4.3.2.2.1.1 2 weight percent………………………….…………..139
4.3.2.2.1.2 5 weight percent………………………….…………..141
4.3.2.2.1.3 1 weight percent……………………………………...143
4.3.2.2.2 Discrepancy in density calculations……………………...…145
4.3.2.2.2.1 Reassessment Theoretical Density Calculations…….….148
4.3.2.3 Influence of Sintering Method……………………………………..153
4.3.2.3.1 Microstructure Changes with Hot Press Heating Schedule...153
4.3.2.3.2 Differences Developed between HP and SPS………………157
4.3.2.3.2.1 Density, Grain Size, Microstructure………………...….157
4.3.2.3.2.2 XRD…………………………………………………….158
4.3.2.4 Grain Boundary Modification through Varying Rare Dopants…….162
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4.3.2.4.1 Hot Pressed Samples………………………………………..162
4.3.2.4.1.1 Gadolinium Samples……………………………...…….162
4.3.2.4.1.2 Lanthanum Samples…………………………………….163
4.3.2.4.2 Spark Plasma Sintered Samples…………………………….164
4.3.2.4.2.1 Densification……………………………………………165
4.3.2.4.2.2 Elastic, Bulk, and Shear Moduli………………………..173
4.3.3 Summary………………………………………………………………..175
Chapter 5 – Effect of Coprecipitation on Selected Mechanical Properties………...…..179
5.1 Introduction…………………………………………………………….………179
5.2 Importance of Plasticity……………………………………………….……….179
5.2.1 Measuring Plasticity………………………………………….…………181
5.3 Procedures………………………………………………………………………184
5.3.1 Hardness…………………………………………………………….…..184
5.3.2 Indentation versus Load Profiling………………………………………184
5.4 Results………………………………………………………………………….185
5.4.1 Plasticity Analysis………………………………………………………185
5.4.2 Indent Analysis………………………………...……………………….189
5.5 Summary……………………………………………………………………….197
Chapter 6 – Conclusions and Future Work……………………………………………..199
6.1 Conclusions…………...……………………………...………………………..199
6.2 Future Work………………………………………………………...…………202
References………………………………………………………………………………205
Curriculum Vita…………………………………………………...……………………215
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List of Tables
Table 3.1 Ionic radii of selected rare earth elements………………………..…………..51
Table 3.2 Density improvement with powder calcination and atmosphere control….…83
Table 3.3 Density values for coprecipitated (C1), ball milled (B1), and propeller mixed (P1) samples in comparison study…………………………………………………….…90
Table 3.4 Mechanical property data returned from ultrasonic non-destructive evaluation of coprecipitated and ball milled samples………………………………………………..98
Table 3.5 Density values for spark plasma sintered samples at 2 and 5 wt% sintering aids, prepared through coprecipitated and ball milling as a comparison……………….100
Table 3.6 Comparison of density and elastic modulus values for coprecipitated and ball milled samples in SPS comparison section……………………………………………..102
Table 3.7 Samples prepared for fluorescence analysis……………………………..…103
Table 4.1 Phase amounts in 10 wt. % coprecipitated sample R61 as estimated by Easy Quantitative function from peak fit in Jade program……………………………...……126
Table 4.2 Phase amounts in 10 wt. % additives ball milled sample as estimated by Easy Quantitative function from peak fit in Jade program…………………………………...129
Table 4.3 Density and moduli values for Al2O3-Y2O3 samples……………….……....150
Table 4.4 Reassessment of theoretical density values for yttrium containing samples.152
Table 4.5 Measured and expected elastic modulus values if theoretical density, phase differences, and porosity are evaluated…………………………………………………152
Table 4.6 Density and grain size results from coprecipitated powder batch split into 3 smaller batches and densified using different hot pressing cycles…………………..…154
Table 4.7 Comparison of HP and SPS Methods………………...………...…………..157
Table 4.8 Density and grain size values for 5 wt. % samples………...……………….158
Table 4.9 Matrix of coprecipitated compositions studied in spark plasma sintering of rare earth modified samples………….……………………………………………..…..166
Table 4.10 Density results for matrix of coprecipitated samples………………...……166
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Table 4.11 Differences in densification and elastic properties with small changes in final SPS holding temperature………………………………………………………………..168
Table 4.12 Density and moduli values for various samples throughout this thesis……174
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List of Figures
Figure 1.1 Parallel (a) and anti-parallel (b) stacking of SiC tetrahedra………….………6
Figure 1.2 Atomic stacking sequences in the common polytypes of SiC………………..7
Figure 1.3 Impact of aluminum on polytype transformations in SiC. At low aluminum contents, type II (6H) is favored. With an increase to 0.05 % aluminum, type 1 (15R) is observed, and upon further increase above 0.10% aluminum, type III (4H) formation is encouraged...... 9
Figure 1.4 Difference in stacking sequence between HCP and FCC sequences……….11
Figure 1.5 (a) Random mixing of additives, likely to develop sintering aid clusters and inclusions (b) Idealized mixing with reduced clustering……………………………...…17
Figure 1.6 Two-particle model showing neck development and important radius of curvature considerations for the driving force of sintering……………………………....21
Figure 1.7 Possible atomic transport mechanisms operative during initial stages of sintering……………………………………………………………………………….….23
Figure 1.8 The stages of liquid phase sintering………………………………...………30
Figure 3.1 Strength distribution and variability in commercial SiC materials showing high variability in fracture strengths. The wide distribution in strengths is indicative of a strong dependence on flaw size and distribution………………………………………...41
Figure 3.2 Carbon inclusion (C) and surrounding interface grains (I) on fracture surface of SiC…………………………………………………………………………………….43
Figure 3.3 Replica micrograph showing carbon inclusions at triple points and grain junctions………………………………………………………………………………….43
Figure 3.4 Formation of electrical double layer due to charge behavior of ions in solution. Curve indicative of changes in zeta potential with distance and ion adsorption near charged surface…………………………………………………………………..…46
Figure 3.5 Net interaction curve showing competition between repulsion and attraction that governs agglomeration or dispersion of particles in solution…………………….....48
Figure 3.6 Highly porous region in liquid phase sintered SiC prepared through coprecipitation. Region caused from lack of sintering aid penetration due to agglomeration of starting powder…………………………………………………….....49
Figure 3.7 Titration curves for selected rare earth systems…………………………….53
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Figure 3.8 Flow chart of processing steps for preparation of coprecipitated samples….55
Figure 3.9 Zeta potential graphs for raw SiC powder and Al/Gd, Al/La, and Al/Y systems at a molar ratio of 4:1 Al to RE. Isoelectric point of raw SiC powders is similar to that of SiO2, as expected. After coprecipitation, isoelectric point values observed to shift to values representative of aluminum and rare earth oxides………………………..72
Figure 3.10 Zeta potential curves showing shifts for Al:RE systems at higher RE content of 3:2 Al to RE. Increase in shift compared to Figure 5.1 at 4:1 Al:RE is due to higher isoelectric point of RE hydroxides compared to aluminum hydroxide……………….....74
Figure 3.11 FTIR Spectra of hydroxide precipitates, uncoated SiC, and SiC - Al/La 3/2 before and after sintering at 1800°C and 1900°C………………………………………..75
Figure 3.12 Micrograph of lower density coprecipitated sample from early powder batch……………………………………………………………………………………...78
Figure 3.13 Particle size graph for raw SiC powder versus post milling treatment powder with measurement in H2O. Decrease in average particle size observed, as well as reduction in 1 to 1.5 micron particles……………………………………………………80
Figure 3.14 Particle size graph for raw SiC powder versus post milling treatment powder with dispersion in isopropanol. Isopropanol when combined with milling treatment reduces particle size and eliminates particles larger than 750 nm………….…81
Figure 3.15 Comparison of all particle size graphs, indicating improved dispersability of SiC with milling treatment and mixing in isopropanol…………………………………..82
Figure 3.16 X-ray diffraction pattern and phase matches for coprecipitated powder before calcination treatment………………………………………………………….…..84
Figure 3.17 X-ray diffraction pattern from coprecipitated sample after calcination in argon at 500°C cycle……………………………………………………………………..85
Figure 3.18 Zeta potential curves for 2 wt% additive coprecipitated samples after milling treatment and with coprecipitation in isopropanol…………………………………….....86
Figure 3.19 Zeta potential curves for 5 wt% additive coprecipitated samples after milling treatment and with coprecipitation in isopropanol………………………………86
Figure 3.20 TEM image of uncoated SiC powder……………………………………...88
Figure 3.21 TEM image of coated SiC powder after coprecipitation process………….88
Figure 3.22 BET results on coprecipitated powders at 2 and 5 wt% sintering aids….….89
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Figure 3.23 Secondary electron SEM image of coprecipitated sample C1……………..91
Figure 3.24 In-lens SEM image of coprecipitated sample C1………………………….92
Figure 3.25 Secondary electron SEM image of propeller mixed sample P1………..…..92
Figure 3.26 In-lens SEM image of propeller mixed sample P1……………….………..93
Figure 3.27 XRD pattern and matches to SiC polytypes for coprecipitated sample C1..96
Figure 3.28 Improved resolution XRD pattern and phase matches for ball milled (B1) and coprecipitated (C1) samples in comparison study……………………………….….97
Figure 3.29 Knoop hardness curve for three comparison samples……………………..98
Figure 3.30 Knoop hardness curves comparing spark plasma sintered samples prepared through coprecipitation (CS) and ball milling (BS) of samples at 2 and 5 wt. % sintering aids…………………………………………………………………………………...…101
Figure 3.31 Emission spectra from coprecipitated samples at 5 (R54) and 2 (R62) weight percent sintering additives……………………………………………………...104
Figure 3.32 Fluorescence decay curves from coprecipitated samples at 5 (R54) and 2 (R62) weight percent sintering additives……………………………………………….105
Figure 3.33 Emission spectra comparing coprecipitated and ball milled samples at 2 weight percent additives total………………………………………………………..…106
Figure 3.34 Fluorescence decay spectra showing possible structural gradients in ball milled samples at 2 weight percent additives total………………………………..……106
Figure 3.35 Fluorescence decay spectra showing improved fluorescence decay in coprecipitated versus ball milled samples at 2 weight percent additives total…………107
Figure 3.36 Elemental analysis and map of aluminum, yttrium, and oxygen at grain boundaries in sample prepared from ball milling. Note aluminum rich and yttrium deficient area in upper left of images…………………………………………………...109
Figure 3.37 Elemental analysis and map of aluminum, yttrium, and oxygen at grain boundaries in sample prepared from coprecipitation showing improved mixing and distribution of elements…………………………………………………………………110
Figure 4.1 Example of core/rim microstructure observed in LPS-SiC. Solid, red lines and circles indicate core regions. Blue, dotted lines outline rim regions………...……114
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Figure 4.2 Variations in density and elastic modulus values with different SPS sintering cycles for 10 wt% additive samples………………………………………………….....120
Figure 4.3 SEM micrograph of coprecipitated sample R61, 90 wt% SiC, 6.4 wt% Al2O3, 3.6 wt% Y2O3. The density appears to be much higher than theoretical density calculations would indicate……………………………………………………………..120
Figure 4.4 SEM micrograph of coprecipitated sample R61, 90 wt% SiC, 6.4 wt% Al2O3, 3.6 wt% Y2O3…………………………………………………………………………...122
Figure 4.5 XRD pattern of 10 wt% additive sample. Only 6H, 4H, and 15R polytypes of SiC identifiable……………………………………………………………………...….123
Figure 4.6 High resolution XRD scan on coprecipitated sample R61………………...124
Figure 4.7 XRD pattern of R61, zoomed in on area of interest. 2:1 mullite phase now identifiable in improved scan…………………………………………………………..125
Figure 4.8 High resolution XRD pattern from ball milled 10 wt. % additive sample...127
Figure 4.9 XRD region showing appearance of Y2Si2O7 phase in ball milled sample..127
Figure 4.10 XRD pattern of 10 wt% additive sample prepared through ball milling. Highlighting the peaks attributed to the excess alumina……………………………….128
Figure 4.11 XRD patterns from 10 wt. % sintering aids for both coprecipitated and ball milled samples. Note excess Al2O3 in ball milled sample……………………………..128
Figure 4.12 XRD pattern comparison of 10 wt% additive samples from coprecipitation and ball milling………………………………………………………………………....129
Figure 4.13 Area of interest for EDS mapping in FESEM…………………………....132
Figure 4.14 Individual elemental maps from EDS scan of sample R61………………133
Figure 4.15 Mixed color map from EDS data showing association of specific elements………………………………………………………………………………...134
Figure 4.16 XRD pattern obtained for 100 percent glass frit……………………….....136
Figure 4.17 XRD patterns from raw SiC powder, overlayed with scans of the powder mixed with 5 and 10 wt. % of the glass frit……………………………………..……...136
Figure 4.18 Grain size comparison for SPS cycle with no intermediate low temperature hold (1900°C for 10 min.) versus cycle with lower temperature intermediate hold (1800°C for 5 minutes, 1900°C for 10 minutes)………………………………………..139
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Figure 4.19 SEM Micrographs of R48, 98 wt. % SiC, 1.29 wt. % Al2O3, 0.71 wt. % Y2O3. Note vacant boundaries and triple points…………………………………….....140
Figure 4.20 In-lens image of R48, 98 wt. % SiC, 1.29 % Al2O3, 0.71 % Y2O3…..…141
Figure 4.21 SEM images of R52, 95 wt. % SiC, 3.22 % Al2O3, 1.78 % Y2O3………142
Figure 4.22 Figure 3.22: In-lens image of R52, 95 wt. % SiC, 3.22 % Al2O3, 1.78 % Y2O3……………………………………………………………………………………142
Figure 4.23 SEM micrographs showing low contrast and vacated triple points in 1 wt. % sintering aid sample…………………………………………………………………….144
Figure 4.24 SEM images of coprecipitated sample R43-1, 98 wt. % SiC, 1.71 % Al2O3, 0.29 % Y2O3. Sample was hot pressed at 1950°C for 30 minutes…………………..…156
Figure 4.25 SEM images of coprecipitated sample R43-2, 98 wt. % SiC, 1.71 % Al2O3, 0.29 % Y2O3. Sample was hot pressed at 1900°C for 1 hour………………………….156
Figure 4.26 SEM images of coprecipitated sample R43-3, 98 wt. % SiC, 1.71 % Al2O3, 0.29 % Y2O3. Sample was hot pressed at 1850°C for 1 hour………………………….156
Figure 4.27 XRD scans comparing 5 wt. % additive samples from coprecipitation and ball milling…………………………………………………………………………..….159
Figure 4.28 XRD scans comparing coprecipitated 5 wt. % additive samples relative to heating method, SPS versus hot pressing………………………………………………160
Figure 4.29 XRD scans comparing 5 wt. % additive samples relative to processing and densification method…………………………………………………………………....161
Figure 4.30 SEM image of R32, 95 wt. % SiC, 2.65 % Al2O3, and 2.35 % Gd2O3...163
Figure 4.31 SEM image of R62, 98 wt. % SiC, 1.08 % Al2O3, 0.92 % Sm2O3…...... 168
Figure 4.32 SEM image of R62, 98 wt. % SiC, 1.08 % Al2O3, 0.92 % Sm2O3……….169
Figure 4.33 In-lens image of R62, 98 wt. % SiC, 1.08 % Al2O3, 0.92 % Sm2O3……..169
Figure 4.34 SEM images of R55, 95 wt. % SiC, 2.65 % Al2O3, 2.35 % Gd2O3...... …170
Figure 4.35 In-lens image of R55, 95 wt. % SiC, 2.65% Al2O3, 2.35 % Gd2O3…...... 170
Figure 4.36 In-lens image of R55, 95 wt. % SiC, 2.65 % Al2O3, 2.35 % Gd2O3. Infrequent regions of large grains such as this observed in some SPS samples………..171
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Figure 4.37 SEM image of R54, 95 wt. % SiC, 2.69 % Al2O3, 2.31 % Sm2O3…….…171
Figure 4.38 SEM image of R41, 95 wt. % SiC, 2.78 % Al2O3, 2.22 % La2O3…….….172
Figure 4.39 In-lens image of R41, 95 wt. % SiC, 2.78 % Al2O3, 2.22 % La2O3…...... 172
Figure 4.40 In-lens image of R41, 95 wt. % SiC, 2.78 % Al2O3, 2.22 % La2O3…...... 173
Figure 5.1 Load-hardness curves for selected samples……………………………...…186
Figure 5.2 Log-log plot of load and hardness…………………………………………187
Figure 5.3 Plasticity values for a variety of samples…………………………….……189
Figure 5.4 Plasticity values for 5 weight percent additive systems………………...…189
Figure 5.5 Comparing 2 kg. indents in SiC-N (left) and coprecipitated sample R62 (right)…………………………………………………………………………………...190
Figure 5.6 10kg indents in coprecipitated hot pressed sample R31 – 5 weight percent additives, Al,Y……………………………………………………………………….…192
Figure 5.7: Indents in coprecipitated spark plasma sintered sample R52 – 5 weight percent additives, Al,Y…………………………………………………………………193
Figure 5.8: 10kg indents in coprecipitated SPS R41 – 5 wt. % additives, Al, La…….194
Figure 5.9: Indents in coprecipitated SPS R56 – 1 weight percent additives, Al,Y…..195
Figure 5.10: Indents in coprecipitated spark plasma sintered R63 – 2 weight percent additives, Al, Gd………………………………………………………………………..196
Figure 5.11: 2 kg. indent in coprecipitated sample SPS R62 – 2 weight percent additives, Al, Sm……………………………………………………………………..…197
xix 1
Chapter 1
Introduction to Silicon Carbide
1.1. Motivation and Background
The continued growth and needs of our advancing civilization has fed the desire for structural materials that could enable the manufacture of taller buildings, engines and turbines that can operate at higher temperatures but are also strong and light, and materials capable of withstanding the environments of space and deep undersea applications. While applications like these could be advanced through the use of engineered structural ceramics such as silicon carbide, a number of technical challenges have continued to hinder their implementation.
Silicon Carbide (SiC) has been an industrially relevant ceramic material since
1893 when Edward Acheson patented the process for mass producing SiC powder. SiC has been utilized in a wide range of applications; as an abrasive, a refractory, wear resistant bearings, semiconductors, electronic devices, and high temperature applications like heating elements and heat exchangers [1, 2]. The ability to perform in such a range of application is due to the inherent properties of silicon carbide: high hardness, low density, high oxidation and corrosion resistance, high thermal shock resistance, good wear resistance, and low thermal expansion [1]. Even with these favorable properties however, the widespread usage of SiC as a structural ceramic has not yet been realized.
The primary cause of this is the inherent brittleness of SiC. Many structural applications still prefer metallic components because of the ability of metals to plastically deform before fracture, even if the corrosion resistance or other properties of a ceramic
2 would be preferred. Ceramics typically show a linear stress-strain response ending abruptly when the fracture stress is exceeded [3, 4]. Metals, on the other hand, when stressed beyond the yield point, display a large region of plastic deformation before fracturing [4]. The inability of brittle ceramics to plastically deform to an appreciable amount means fast fracture is the only mechanism through which they can release stress concentrations, and fracture is catastrophic and severe [3]. Few structural applications can tolerate the risk and damage of fast fracture so improving the plastic response of ceramics is essential in their further development.
The basic pathway for improving the brittle strength of ceramics was indicated by the work of Griffith back in the 1920’s, but the practical challenges of ceramic processing have prevented its complete realization. Griffith’s equation indicates the strength of a ceramic as:
-1/2 σ = Y *KIc * c (Equation 1)
where σ is the fracture strength, KIc is the Mode I fracture toughness, c is the size of the
critical flaw, and Y is a constant related to the shape and location of the flaw [5].
Examination of Griffith’s equation reveals two methodologies to improving the brittle
response of SiC: 1) Modify the c parameter by reducing the size and severity of
microstructural flaws and 2) Increase K1C; improve the fracture toughness of the material.
Much research has been done into both approaches.
When examining SiC microstructures, it is not uncommon to observe a number of
flaws. These features are often undesired residual effects from processing such as large
pores, unreacted sintering aids, impurities, residual phases and material from processing
like milling media, and even features like large grains [6, 7]. The highly covalent nature
3 of the bonding in silicon carbide, while providing many of the excellent properties listed above, makes sintering silicon carbide to high densities difficult. Starting from the pioneering work of Prochazka with solid state sintering of SiC using boron and carbon sintering aids highly dense silicon carbide was possible [8]. Achieving high densities through solid state sintering necessitated high temperatures however and often required sintering additives in amounts that lead to defects. Solid state sintering also only allowed for somewhat limited control of microstructure and properties.
Researchers later discovered that SiC can be liquid phase sintered utilizing combinations of oxides such as alumina and rare earth oxides. During sintering these additives react with the native silica layer present on the SiC powder to form a liquid capable of promoting densification at lower temperatures than solid state sintering [9,
10]. Upon cooling this liquid phase typically remains as a residual phase at the grain boundaries and can have a profound effect on mechanical properties [10]. Becher et al. have shown the influence of modifying the grain boundary phase in Si3N4 and observed
drastic changes in composition and properties through small additive changes [11]. By
modifying the processing, additives, and other features, liquid phase sintering allows for
the development of many new variants of SiC with a wide range of microstructures with
varying grain sizes, grain shapes, compositions, and grain boundaries [12-14].
SiC with improved fracture toughness through microstructure control via liquid
phase sintering is commonly observed in literature. A popular approach to improve
strength in sintered silicon carbide is to develop microstructures with enhanced fracture
toughness through the growth of elongated grains that encourage longer crack paths [15,
16]. For many applications, the improvement in fracture toughness from this type of
4 microstructure is beneficial. But, on a more critical level, this type of solution only offers marginal improvements under the right situations. Under certain loading conditions these large grains could function as the critical flaw, fracture could initiate at a lower stress level, and the material may perform poorly [17]. Some of the processing steps required to develop this type of microstructure may lead to the development of other detrimental defects such as increased porosity from improper annealing or a change in the grain boundary phases [18].
For silicon carbide to advance further as a structural material for higher demand applications both approaches revealed by the Griffith equation must be applied concurrently. Independent improvements in either aspect can mask the real performance benefits for some applications; i.e. the large acicular grains that enhance fracture toughness can also be the largest defect in a material and decrease the true strength.
Flaws must be removed from the material to increase strength and improvements in fracture behavior needed to be gained from other microstructural features besides crack path modification from elongated grains. Conventional processing using milling or mixing to add the sintering aids can result in poor mixedness and regions deficient in sintering aids. This ultimately limits the sinterability of the silicon carbide and forms defects in the final microstructure.
This thesis will examine the use of coprecipitation as an advanced processing technique with the goal of developing a fine grained, liquid phase sintered (LPS) SiC.
The improved mixing and scale of the sintering aids will enhance the homogeneity of the phases in the sintered products as well as improve sinterability, leading to greater control of both the microstructure and grain boundary phases. The reduction of defects is
5 expected to allow the grain boundaries to be more influential in the fracture of LPS-SiC.
Further coupling the advanced processing with novel densification methods such as spark plasma sintering can enable rapid sintering cycles and the development of grain boundaries with improved mixedness and interesting phases. It is believed that the fine grained liquid phase sintered SiC with modified grain boundary characteristics may display interesting fracture behaviors versus more conventionally prepared SiC materials.
1.2. Silicon Carbide
1.2.1. Crystal Structure
Silicon carbide has great potential in many applications due to its unique combination of processing methods, structures, and properties. Many of these benefits are afforded directly or indirectly by the bonding and structure of SiC. The short range order of the crystal structure can be depicted by stacking alternating layers of hexagonal atoms where a layer of carbon atoms sits atop a layer of silicon atoms with the carbon displaced into the close packed positions above the silicon. This stacking can be viewed in a more simple and representative form by instead treating the basic unit as tetrahedra of Si4C, or similarly C4Si, joined at the corners. By picturing layers of joined tetrahedra
stacked along the basal plane direction with their apex pointing up, the structure is
synonymous to individual atoms in a close-packed lattice. If the directionality of the next
layer of tetrahedra is varied relative to the previous, as shown below in Figure 1.1, it is
possible to develop a nearly infinite number of stacking sequences in one dimension by
altering the number of layers before repeat. Although this results in a different structure
it also preserves the stoichiometry of the system. This ability to exist in a number of
6 differently ordered structures with the same stoichiometry is known as polytypism.
Whereas many materials are polytypic to a limited extent, i.e. α-iron, γ-iron, α-Al2O3, γ-
Al2O3, the polytypism of SiC is extensive with over 200 polytypes having been observed
[19-21], and the existence of many others not out of the question.
Figure 1.1: Parallel (a) and anti-parallel (b) stacking of SiC tetrahedra [19]
A number of notations have been developed to describe the different structures of
polytype materials. The most common notation, Ramsdell’s, labels the polytypes as nL,
where n is a number indicating the periodicity in the stacking of the tetrahedra layers
along the c-axis and L is a letter indicating the general crystal symmetry. For example,
3C is indicative of cubic symmetry with 3 layers before repeat. This is the only cubic
polytype for SiC and is designated as β-SiC. Other symmetries and stacking sequences
are common and possible. Many of the most common polytypes of SiC have hexagonal
symmetry, notably the common polytypes of 2H, 4H, and 6H. There is one common
rhombohedral polytype, 15R, and countless other less common and more exotic
combinations like 33R or 1200R. All of the non-cubic polytypes although different, are
grouped together and considered as α-SiC Many other notations and designations for
polytypes have been developed as well as theories behind their stability and formation,
7 but this single notation and cursory explanation should be sufficient in the scope of this work. Examples of many of the common polytypes of SiC are shown in Figure 1.2.
Figure 1.2: Atomic stacking sequences in the common polytypes of SiC [22]
X-ray diffraction is the most common way to characterize the presence of specific
polytypes in SiC. While it is often simple to qualitatively observe the presence of a
certain polytype in an XRD pattern by finding certain characteristic peaks, it is not nearly
as straightforward to quantitatively determine all of the polytypes present because of
overlapping peaks. Although many researchers do know the general polytype grouping
of their initial powder, α or β, they often fail to be more specific as to the amount of 2H,
4H, 6H, etc. present in the starting powder because of the effort required to be specific.
A number of researchers have developed simple relationships between peak intensities in
order to evaluate the polytype differences [23, 24]. These methods can be simple to use
mathematically, but their accuracy is questioned by some researchers, and they can give
indeterminate results in many cases. More accurate analysis can be performed using
8
XRD whole pattern fitting methods such as Rietveld refinement or techniques such as
Raman or nuclear magnetic resonance. These procedures can be difficult and costly to implement however.
The polytypes of SiC can not be ignored during the development of microstructure. Depending upon the temperature profile employed during densification as well as other factors such as impurities, a given polytype can transform into a more stable one during sintering and this can be accompanied by desirable or undesirable grain growth and changes in densification, properties, and performance [19, 25-27]. For example, it is common to improve the fracture toughness of SiC by exploiting the anisotropic grain growth that occurs when polytypes transform. The high sintering temperatures required for densifying SiC promote the transformation of β grains to
α grains and this can develop large, elongated platelet grains. By seeding an α SiC powder with β grains before sintering, microstructures with improved fracture toughness can be designed by taking advantage of the increased crack paths around the elongated grains. In other cases, these large grains can act as critical flaws and lead to decreased mechanical properties such as strength [28, 29].
Impurities and sintering additives play an important role in the development and transformation of polytypes. In 1948, Lundqvist researched the observations that different polytypes were often associated with varying colors of SiC crystals in certain powders: 6H were green, 15R were yellow, and the appearance of 4H grains or mixed samples made samples appear black. Through the use of careful x-ray examination of over 200 powders from a variety of locations and spectrochemical analysis on a number of these powders, large variations in aluminum content and smaller variations in iron
9 content were noted. With low aluminum contents the 6H polytype appeared to be favored, where as 0.05 – 0.06 wt. % aluminum encouraged 15R, and a transition to 4H formation above 0.10 wt. % . These polytype changes with aluminum content are shown in Figure 1.3. He also examined the inclusions microscopically noting few inclusions in the clear to light green samples, most of which were unreacted graphite. An increased number of fine carbon inclusions were detected in the darker green samples. In the darkest, black samples obvious large defects and compounds of aluminum and iron, were often present as well as disturbances in the nearby crystal structure. This indicates a low solubility for these impurities in the crystal structure [30]. A wide range of other impurity elements and sintering aids also have influence over the temperature at which the polytype transformations occur and the exact sequence of the transformations [31-35].
Figure 1.3: Impact of aluminum on polytype transformations in SiC [30]. At low aluminum contents, type II (6H) is favored. With an increase to 0.05 % aluminum, type 1 (15R) is observed, and upon further increase above 0.10% aluminum, type III (4H) formation is encouraged.
The temperatures required for densification also have influence over the transformation of polytypes in SiC. Densifying SiC at temperatures above 1900°C will
10 cause β grains to transform into α polytypes and this is accompanied by rapid, anisotropic grain growth [25]. If the initial material is instead an α powder, sintering at or above 1900°C will result in a fine, equiaxed microstructure. Certain conditions can also cause α grains to reverse transform back into β, or encourage the stability of 4H polytypes over 6H. Careful control over purity, sintering aids, and processing is clearly required in order to modify the microstructure of silicon carbide relative to polytypes.
1.2.2. Stacking Faults
It is convenient to describe the stacking sequence of an entire crystal as a specific polytype of SiC such as to indicate a grain of 4H polytype SiC within a sample of otherwise 6H polytype grains. This is not the only way to consider the importance of the atomic positions in SiC however. It is also possible to have a small, localized change in stacking sequence within any specific grain. In this case, the specific grain would still be a majority of 6H, but there can be local regions where the stacking sequence has changed to 4H positions for a few layers and then back to 6H. These changes in stacking sequence within a single crystal are known as stacking faults.
The understanding of stacking faults and their connection with plastic deformation behavior has been developed in the study of metals for many years now. A normal face centered cubic metal (fcc) can be described by the stacking sequence of the close packed planes of atoms: ABC ABC. The hexagonal closed packed (hcp) crystal structure can be visualized by shearing the C plane back into the A position, creating a sequence of ABAB [4, 36]. The representative atomic positions for each stacking sequence are shown in Figure 1.4. The process of slip on a close packed plane can
11 produce the same shift in stacking sequence for a number of layers in a crystal; this shifted region is known as a stacking fault.
Figure 1.4: Difference in stacking sequence between HCP and FCC, redrawn from [36]
The stacking fault itself can be vectorially defined as an extended dislocation that
is bounded by two partial dislocations. Like all dislocations, the stacking fault has an
energy associated with its creation and this energy is typically very different than that of
the parent crystal. These energies can also differ greatly between materials. Low
stacking fault energy (LSFE) materials readily form many stacking faults and have large
separations between partial dislocations. High stacking fault energy materials (HSFE)
require more energy for creation and therefore form fewer and narrower faults. Silicon
carbide has a low stacking fault energy and it is not uncommon to find many stacking
faults present through crystals.
The multiple polytypes of SiC each have their own stacking fault energy, but regardless of which polytype is the majority, the SFE of SiC is low in comparison to metals, with values ranging from -5.80 meV for 3C to 13.56 meV for 6H [37]. The
12 predominant source of stacking fault research on SiC is from investigations of electronic applications. Heavy nitrogen doping and voltage cycling have been observed to lead to an increase in stacking faults in SiC and a decrease in electronic performance [27, 38].
Overall however, SiC typically has a low SFE even when the system is modified by processing or dopants. With such a low SFE, SiC is expected to favor the formation of stacking faults and twins under normal processing conditions. Rubble from ballistic impact experiments do indeed show a considerable amount of stacking faults and twins
[39, 40]. It is believed the slip of the partial dislocations may be an important source of plastic deformation behavior for SiC.
Stacking faults can be an important, but overlooked feature in the microstructure and mechanical behavior of many materials. Ceramics are by nature brittle materials because of their strong bonding and restricted conditions for slip, and little plasticity is measured in most systems under general stress and strain conditions [4]. Under high loads and strain rates, such as those encountered in some structural applications or impact events however, plastic behavior has been witnessed and is believed to be important in the bulk material response. Low stacking fault energy materials twin readily under shock loading also as the presence of greater numbers of stacking faults provides locations for twins to form easily [41]. The movement of stacking faults can be a potential source of plastic deformation for ceramics and influence the inherent brittle nature of SiC.
13
1.3. Green Processing of SiC
Like most ceramic systems, SiC materials begin as a raw powder batch that requires a series of processes to obtain final products with the size, shape, density, and performance that is desired. These processing steps may entail the addition of sintering aids to the starting powder, mixing or milling steps to distribute these additives, additional binders or pressing aids, drying steps, if wet milling is used, and a pressing step that lightly assembles the powders into the desired size and shape. This process of forming an unsintered body, called a green body, ends with a low density, porous SiC body that requires a densification or sintering step to eliminate the large pores between powder particles and develop the microstructure as desired. Green processing is also the source of the majority of defects observed in SiC materials. The forming steps that bring the loose powder particles together before sintering are also essential in limiting the development of defects; hard agglomerates, poor dispersion of sintering aids, new impurities, and pressing defects can be introduced during green body formation and can add to the variability and failure of a ceramic [42]. The following introductory sections serve as a walkthrough of many of SiC processing.
1.3.1. Production of SiC Powders
A majority of the silicon carbide produced today is manufactured through static bed carbothermal synthesis also known as the Acheson process. Discovered by Edward
Acheson in 1891, the process creates silicon carbide through the reaction of quartz (SiO2)
and coke (C) in a heated resistance furnace. The process is typically carried out in very
large resistance furnaces capable of producing 10 to 250 tons of material per firing. The
14 furnace is loaded with the ground raw materials, typically quartz and coke, surrounding the electrode. The furnace charge is protected from reacting with the surrounding containment by a layer of unreacted SiC. The furnace reaction is heated by passing large current through graphite electrodes surrounded by the precursor silica and coke. As current is passed through, a large amount of heat is generated at the electrode surface.
This heat is then transferred through the furnace charge. The exact temperature generated during the process can be imprecise because of differences in raw material, impurities, and furnace properties [43]. However, temperatures generally range from 1500°C to
2300°C. Above 1500°C, various reactions between the quartz and coke can begin and
SiC starts to form, as shown by the most common reaction path:
SiO2 + 3C Æ SiC + 2 CO (Equation 2)
After a long firing time of 40 hours or more, the system is allowed to cool. The outside layer is broken away from the furnace and this unreacted material is removed.
The inner sections of SiC are then also broken apart and manually separated by physical examination of color and other factors. These pieces are selected into the various product types, ranging from impure SiC for abrasives to the purer powder close to the electrode that is used for advanced ceramics. This powder is segregated, depending upon the intended applications, for further crushing, milling, and acid washing to make ceramic powders. This relatively manual and non-scientific method of powder development has very important ramifications. Because of the scale of the process, impurities in the starting materials, excess additions of graphite and impure SiC, and the powder selection method, the material yielded from a large Acheson furnace can have highly variable quality relative to both impurity content and polytype consistency across a batch.
15
There are quite a few other methods of producing SiC, often resulting in powders with much better uniformity, purity, and quality. SiC has been prepared through chemical methods such as sol-gel processing as well as the decomposition of polymers and chemical processing from biological materials such as wood [44-46]. SiC has also been produced from innovative laser methods and chemical vapor deposition [47, 48].
The drawback for most of these methods is in the much smaller batch size, the need for specialized equipment, and high cost.
1.3.1.1. Importance of the Raw Powder
The microstructure, mechanical properties, and performance of a dense ceramic are ultimately dictated by careful control of each and every processing step from the raw powder to the green body forming stage to the final firing and finishing. Properties of the raw starting material, such as size, shape, impurities, and surface character need to be considered in all following steps. Sintering will not eliminate defects previously introduced into the material; it can only serve to exacerbate or create other defects [49,
50]. Therefore any attempt to remove defects in SiC must begin with control and improvement in the initial powders and green processing, and not just in sintering.
Like most ceramic systems, SiC processing begins from a raw powder which has been beneficiated by crushing and milling and has been processed through washing or other steps into a clean, relatively impurity free starting powder of a sufficient size. The properties of the starting powder are important in the control of the final product.
Powders tend to have a high surface area to volume ratio which impacts the absorption behavior, agglomeration, and other important processing variables. The structure and
16 energy conditions of a free powder make it favorable to absorb species through chemical or physical bonding [42]. The polytype distribution, particle size and morphology, and impurity level are also strongly influenced by the initial raw powder characteristics.
Controlling the raw powder properties requires entirely new beneficiation or powder synthesis steps, and the approach for instituting this research can be more costly and long term although certainly worthy of research. However, this is not the focus of this thesis.
1.3.2. Conventional SiC Processing
Historically, most silicon carbide was processed through simple mechanical mixing of silicon carbide powder with the sintering aids added as powders. During the earliest years of production additives such as boron carbide and carbon black were mixed into batches of SiC as dry powders through milling [51-53]. Other methods were later designed with a wide range of additives from boron nitride to silicon nitride to alumina
[8, 13, 54, 55]. Not surprisingly, silicon carbide produced in this way often had regions with too much sintering aid and regions deficient in additives, resulting in poor final properties. Boron or boron carbide added as powders can form inclusions in the final microstructure due to poor mixing or agglomeration. Carbon, as a powder additive, was often a larger problem because of its poor grain boundary diffusivity in SiC [56]. Studies on ideal and random mixing have shown that adding sintering aids as powders can only result in random mixing at best, as opposed to the preferred ideal case as shown below in
Figure 1.5
17
(a) (b)
Figure 1.5: (a) Random mixing of additives, likely to develop sintering aid clusters and inclusions (b) Idealized mixing with reduced clustering. Redrawn from Carlstrom [56]
Later, researchers began processing SiC with polymer based carbon sources, such as phenolic resins, alone or in combination with other powders [57-59]. After dissolution into a solvent, the resin can be easily mixed with SiC powder. When the solvent is evaporated, a coating of the polymer remains. The coating ability of the resin additives improved the homogeneity and distribution of sintering aids leading to improved sintered
products. The mixing and coating of resins is still far from optimal however in most
systems, and SiC processed in this matter still shows numerous defects from poor mixing.
As the understanding of the sintering process and computer modeling has advanced, it is now possible to calculate the exact amount of sintering additives required to yield high density. This understanding can help minimize the need to add excess sintering aids [21].
Research into engineered surfactants, paints, and other coatings is also ongoing
[60]. These represent further improvement, but these do not fully address the issues of poor mixing, agglomeration, and defect formation.
18
1.3.3. Colloidal Processing
Colloidal processing has replaced dry mixing in most modern green processing procedures for SiC and has resulted in the decreased observation of defects [61, 62].
Controlling the colloidal processing environment allows for improved dispersion and increased homogeneity of a green body. Modification of the surface charge and character, as well as the application of shear forces during colloidal mixing, can break apart agglomerates and improve particle separation. This can result in the sintering aids attaching to favored surfaces and avoiding uncoated regions within agglomerates.
Researchers have investigated numerous methods of improving the processing of SiC through colloidal processing.
Li et al. added tetramethylammonium hydroxide (TMAH) as a dispersant in an aqueous SiC slurry in order to improve the rheological behavior of SiC. The addition of the TMAH promoted the disassociation of the silanol groups on the SiC surface, imparting a more negative charge in aqueous solution. Samples prepared this way had an increased dispersability and lower viscosity of the slurry as well as a higher green density
[61]. Rao et al. studied the dispersability of various SiC powders in water versus the pH.
They observed the optimal dispersion occurred at a pH of 10 [50]. SiC slurry dispersion was also improved by the addition of polyethylene imine (PEI) [63]. Vasanthakumari developed a non-aqueous tape casting procedure with high green densities using phosphate ether as a dispersant and an isopropanol-toluene azeotopic mixture as the solvent system [64] . Krishnan similarly modified the dispersion of SiC in aqueous solution to obtain better tape castings using a number of dispersants. Defect free tapes could be formed even at higher solids loading when properly dispersed [65].
19
Many of these processes improve the rheological properties of the slurry through the use of additional dispersants or additives. While these methods do improve the processing and reduce defects, these methods can also result in the introduction of excess carbon or other unwanted features into the systems and still rely on the initial quality of the additives in term of particle size and distribution. Methods introducing additives on a finer scale with more intimate chemical homogeneity would result in a further reduction of defects and possible changes in microstructure and behavior.
1.3.4. Sintering
Regardless of the method of green processing a substantial portion of the considerations made during those steps are to prepare the powder for sintering where the majority of strength and mechanical properties are physically developed. It cannot be emphasized enough that processing and sintering are intimately related; improper development and/or execution in either area will doom the other. Poor processing can not be eliminated through advanced sintering, and the best process developed to date can be undone by poor sintering design and performance. Therefore a complete discussion on advanced processing must also consider the advancement of sintering.
1.3.4.1. Sintering Theory
Sintering is the formation of mechanical bonds between loose powder particles by thermally activated diffusional processes, resulting in a dense final body. The elimination of porosity is essential in developing high elastic moduli and other properties.
Sintering can be carried out at high temperatures, just under the melting point of the
20 primary phase, through solid state atomic diffusion, or can be accelerated through the formation of a liquid phase which can fill pore space, enhance atomic transport, and greatly modify properties [10]. Separate sections that follow will address the relevant differences between the solid and liquid phase sintering processes. Regardless of the method, sinter bonds develop as necks form between contacting particles and porosity is slowly eliminated. This greatly increases the strength of the material versus the porous, unfired green body state.
The driving force for these bonds forming is the lowering of the high surface energy associated with the high curvature of loosely contacted powder particles. The elimination of solid-vapor interfaces, i.e. pores, by forming solid-solid interfaces, i.e particle necks and eventually grain boundaries, is favorable for lowering the surface energy of the system. The specific mechanisms that drive the atomic motion that moves material to necks is also surface energy dependent [66]. The mathematical treatment of
Young and LaPlace on the effect of surface tension and curvature on the stress of a surface reveals that curved surfaces are in a higher stress state. Moving towards a flatter surface lessens the stress and energy of the system [67]. The equation is shown below where Δp is the pressure difference between two surfaces, γ is the surface energy, and r1 and r2 are the principle radii of curvature for the surfaces involved.
⎛ 1 1 ⎞ Δp = γ ⎜ + ⎟ (Equation 3) ⎝ r1 r2 ⎠
The work of Young and LaPlace was later refined by Gibbs, Thomson, and
Freundlich into a relationship more exact to sintering. Considering the region between
two particles, as shown below in Figure 1.6, the Gibbs-Thomson-Freundlich equation
21 indicates that surface curvature is the driving force for sintering. The curvature at the sintering necks has a drastic effect on the chemical potential difference between the two surfaces which is easily visualized through the Gibbs-Thomson-Freundlich equation:
⎡γV ⎛ 1 1 ⎞⎤ m ⎜ ⎟ c = c0 exp⎢ ⎜ + ⎟⎥ (Equation 4) ⎣ RT ⎝ r1 r2 ⎠⎦
Where c is representative of the vapor pressure, c0 is the equilibrium vapor pressure, r 1
and r are the radii of curvature for the curved surfaces, V is the molar volume, γ is the 2 m
surface energy per unit area of the surface, R is the gas constant, and T is the absolute
temperature.
Figure 1.6: Two-particle model showing neck development and important radius of curvature considerations for the driving force of sintering, redrawn from [67]
As the equation shows, a curved surface will have a higher chemical potential for
its components and therefore the vapor pressure above the surface will also be greater
than a flat or concave surface [67]. Mass will therefore be driven from the large
positively curved surface with radius R and into the sharply negative regions, radius r in
Figure 1.6, where the grains meet and necks form. The consideration of curvature also
defines a capillary pressure which exists because of the curvature of a forming neck; this
capillary pressure serves to draw the particles together. These energy considerations
22 drive the movement of atoms from the particles into the formations of necks and provide pressure to draw the particles together into a sintered body. They further indicate the role of particle size in sintering as finer powders will have greater curvature and therefore more driving force, and could lead to quicker sintering and higher densities [66].
More specifically, the neck formation is driven by six atomic transport mechanisms as shown below in Figure 1.7. Although all of the mechanisms are operating concurrently in most systems, which specific mechanisms are dominating the process will highly influence the densification. The high temperatures of sintering, often near the melting point for solid state systems, provide the needed energy to promote the motion of atoms and vacancies. How and where the atomic transport occurs determines the densification and grain growth behaviors of a sintering material [10, 20, 67].
23
Figure 1.7: Possible atomic transport mechanisms operative during initial stages of sintering [66, 67]
The formation of a neck between two particles can draw the particles closer together if the movement of the atoms is from the bulk or grain boundaries. This movement can also draw any other connected particles together. This means that as the necks form and grow, the porosity of the part will decrease as the centers of the particles move closer together and accordingly, the overall size of the part should also decrease slightly. This occurrence is called shrinkage and is an indicator of efficient sintering and an important consideration of pressureless sintering as the part can change shape and size considerably. The amount of shrinkage that occurs during sintering can differ widely over a range of materials due to the different mechanisms that can operate [10, 66].
24
If the atoms and vacancies are moving from the surface of one particle to another, as in surface or lattice diffusion, necks will grow, but this is not really a densification process. This type of neck growth does not draw the particle centers closer together nor reduce the porosity. This situation can exacerbate grain growth without limiting the final pore volume of a material. A system that favors surface or vapor diffusion processes will not sinter efficiently without sintering aids or careful control. If instead the transport is from the bulk of the material to the neck, as in bulk or grain boundary diffusion, material shrinkage will be observed and true densification is occurring. Higher temperatures favor bulk diffusion mechanisms; hence why high temperature isothermal holds are important for the sintering of many materials. In most materials there is observation of parallel mechanisms, but the use of sintering aids, external pressures, and correct sintering cycles is essential in promoting the necessary mechanisms and developing the desired microstructure. [10, 20, 66][5].
1.3.4.2. Assisted Sintering Methods
1.3.4.2.1. Hot Pressing
The discussion of the basics of sintering above indicated a capillary pressure exists at the necks of a green compact and that this force aids in the process of sintering.
This force is small in most cases, on the order of less than 1 MPa, and unless the powder particles are very fine, it is not surprising that most materials, particularly covalently bonded ceramics, like SiC and B4C, do not sinter easily without additives or very high
applied temperatures [67]. In cases where the amount of additives is preferred to be low
or there are limits on particle size reduction, it is often most practical to increase the
25 driving force for sintering by providing an external sintering pressure. Hot pressing is the application of an external compressive force on the particles during sintering. Although this can increase the cost of production and requires more complicated furnaces, dies, and shape control, it can be a very practical way to sinter materials where low additive contents and very high densities are necessary [21, 67].
1.3.4.2.2. Spark Plasma Sintering
Spark plasma sintering (SPS) has shown itself to be an effective method of sintering materials to high densities with limited grain growth and no binders [68]. It has been utilized in the densification of a wide range of ceramics including SiC, and can produce high densities at temperatures considerably lower than hot pressing or pressureless sintering [69].
Densification using SPS differs from conventional furnace sintering methods in that the heat is being created directly from the passage of high current through the sample as opposed to transferred from the outside of the sample via induction from a die or from surrounding heating elements. There are a number of different theories as to the exact heating mechanism employed in SPS, and whether an actual “spark” or “plasma” exists.
One of the more widely believed theories is that the application of the high current, when pulsed quickly between on and off, generates extremely high temperatures at the contacts between particles, enough to generate a small local spark plasma. This leads to very rapid and efficient heating as the current discharges along the intergrain pathways in the material dispersing with the constant cycling. The generation of the sparks and/or plasma is thought to develop temperatures in excess of thousands of degrees, leading to the
26 removal of impurities, vaporization of surface material, and local liquid formation. The presence of the liquid and capillary pressure, in addition to the applied pressure, serves to draw the particles together and sinter them. The entire heating process is more efficient than standard methods and leads to reduced temperatures and cycle times [68, 69].
Researchers have observed differing sintering kinetics and interesting effects from the SPS of SiC and other ceramics. Xu et al. studied the spark plasma sintering of SiC and observed little sintering at temperatures under 1600°C but rapid neck formation at
1850°C [70]. Hojo et al. spark plasma sintered nano-SiC with AlN and Y2O3 and
obtained very fine grained, dense microstructures. They found that the sintering additives
were necessary in order to retard abnormal grain growth even while sintering in the SPS
[71]. Guillard et al. performed an SPS study on SiC focused on determining the different effects of temperature, pressure, and time. Samples without additives were sintered to
densities greater than 92 percent at 1850°C [69].
1.3.4.3. Solid State Sintering
Solid state sintering occurs in a series of stages during which both the
mechanisms of the sintering process and the resultant pore structure undergo changes.
Although the features of the stages can be discussed, in reality they are not as clearly
distinct as presented. During the initial stage neck growth occurs quickly because the
majority of the mass transport that is happening is surface transport between the two
contacting particles. This is occurring at reasonably low temperatures and involves
transport over a short distance between the particles. This results in no shrinkage and
little change in density occurring during this stage. The porosity remaining at the end of
27 this stage is continuous and open; the grain size is the same as the initial particles. The initial stage typically begins from a green density of 50 to 60 percent of theoretical, and ends at 70 percent [20, 66, 72].
In the intermediate stage of sintering the mechanisms of transport change from surface to bulk and grain boundary transport and the density increases substantially from
70 percent to roughly 90 percent of theoretical. Diffusing atoms now no longer originate from the surface of the powder particle but from within the interior of the particles.
Diffusion now occurs over a longer distance beginning in the inside of the particle and terminating at the neck. Considerable shrinkage therefore accompanies the densification during this stage. The pore structure smoothens during the intermediate stage, but remains open and interconnected until the late intermediate stage. At this point the pores will begin to close leaving isolated, spherical pores. This will be accompanied by a decrease in the sintering rate and some limited grain growth as the small spherical pores are not as effective in limiting growth [20, 66, 72].
Densification during the final stage will be slow and occurs in constant competition with grain growth. Bulk diffusion is the only mechanism capable of closing individual isolated pores while grain boundary diffusion can modify only pores located on a boundary. These processes are slower and more sensitive to sintering conditions so the removal of the last bit of porosity can be difficult. The shape and curvature of pores can now cause some to grow larger at the expense of others. Gases can be trapped in the pores, limiting the achievement of high densities unless vacuum processes are used. If sintering is performed properly however, it is possible to develop densities near theoretical [20, 66, 72].
28
The addition of certain elements as sintering aids can have a large effect on improving the sintering behavior. The presence of certain solute atoms can promote the boundary or lattice diffusion mechanisms which most effectively reduce porosity. Other sintering aids are necessary to remove rate-limiting oxide species. Sintering aids which promote the formation of liquid phases, as in liquid phase sintering, are discussed later
[20, 66, 72].
1.3.4.4. Liquid Phase Sintering
Issues with high temperatures, poor microstructure control, phase transformations, and lower than theoretical final densities have resulted in many researchers seeking alternatives to solid state sintering [20, 21]. In many cases, solid state sintering using a hot press or more innovative methods like spark plasma sintering of pure powders still fail to deliver the high densities or desirable microstructures. In many cases, liquid phase sintering (LPS) is an attractive alternative method.
As opposed to relying on the atomic transport mechanisms native to the starting powder, liquid phase sintering involves choosing sintering aids and cycle conditions that encourage the formation of a liquid phase that serves to enhance transport during sintering. The presence of a liquid during sintering increases the atomic diffusivity and capillary forces resulting in faster sintering. The capillary force in the presence of the liquid also functions as a lubricant, allowing for rapid rearrangement of the packed powder into a higher density configuration. This is without considering some of the other changes and benefits offered by the LPS process. Just as with solid state sintering, liquid phase sintering has a number of modifications and stages which can be loosely defined.
29
For a system to be considered liquid phase sintered it does not need to meet all the conditions outlined or follow the full sequences of steps, it is only required to have a liquid form during the sintering process; the liquid does not necessarily have to be permanently visible in the final microstructure [10, 72].
Liquid phase sintering is often introduced through the mixing of sintering aids and the melting of either a single component or a eutectic reaction among multiple components. In many oxide ceramics or systems with a native oxidation layer, like the silica layer present on SiC, liquid phase sintering of these systems can involve reactions of the sintering aids with this oxide. To be an effective sintering mechanism, sufficient liquid must be present and it must be able to wet the solid phase so compositional control and temperature conditions are vital in LPS as well as in solid state sintering [10, 72].
The primary variations of liquid phase sintering are defined by the solubility of the various phases for each other. In one case, it is possible for the liquid to have a high solubility in the solid phase. Although this case is generally unfavorable for typical liquid phase sintering, if the liquid formation process is well controlled, the sintering can be boosted by the presence of the liquid through rearrangement and capillary effects only.
It is even possible in this case for the liquid to completely diffuse into the solid during the process, resulting in no visible liquid at the end of the sintering cycle; this is a very special case of LPS known as transient liquid phase sintering. This is uncommon, but can be advantageous for some systems. In more typical liquid phase sintering systems the solid phase will have a high solubility in the liquid phase [10, 72, 73].
There are three stages in a typical liquid phase sintering process: rearrangement, solution-reprecipitation, and solid state sintering. These stages are shown in Figure 1.8.
30
As discussed earlier, rearrangement is the most beneficial stage of liquid phase sintering, and much of the boost in densification rate is from this occurrence. During rearrangement, the wetting liquid fills the open pore space between powders and boosts the capillary pressure for sintering, moving the particle centers closer together and quickly increasing the packing and density [10].
Figure 1.8: The stages of liquid phase sintering [10]
In the second stage, the presence of the surrounding liquid causes finer grains to go into solution as well as making all grains more rounded in shape. This establishes a concentration gradient in the liquid and the dissolved phase can transport quickly through the liquid and reprecipitate out on the larger grains. This is a coarsening process and contributes to grain growth during this stage, but it also increases the densification because of the improved packing of the more uniformly sized, rounder grains [10]. The solution-reprecipitation process has important ramifications in microstructural design,
31 and depending upon the sintering conditions, can result in core-rim or core-shell development in the final microstructure [20, 74]. This feature will be discussed later.
The final stage of liquid phase sintering is referred to as solid state sintering. The mechanisms that operate during this stage are analogous to those of the late stages of solid state sintering: the closing of remaining pores. This stage is slow and unfavorable for microstructural development as grain growth can be exacerbated during [10].
1.3.5. SiC Sintering Studies
1.3.5.1. Solid State Sintering
The highly covalent bonding of SiC and subsequent low self diffusion coefficients of Si and C make sintering very difficult without the use of high temperatures and pressures [75]. As a result, SiC is typically processed with various sintering additives and a wide range of methods. Starting from the pioneering work of Alliegro in the
1960’s with a number of metallic additives [76] and Prochazka in the 1970’s using boron and carbon as sintering aids to pressureless sinter nearly dense SiC [8], many researchers have developed different additives, combinations, precursors, and conditions for the sintering of SiC. This section is not meant to be an all inclusive summary, but will briefly introduce some of the more promising methods while focusing on those most relevant to structural applications.
The majority of silicon carbide is still processed similarly to the work of
Prochazka using 1 wt. % boron and carbon additives [8]. While Prochazka was able to process SiC to above 96 % dense without pressure using these additives, it is common today to use a similar composition but instead hot press the components to remove the
32 residual porosity and ensure high density. It is generally believed the addition of carbon serves to react with the native SiO2 layer present on the SiC powder. This oxide layer
greatly inhibits diffusion along the grain boundaries and its removal is key in successful
sintering [58]. Researchers indicate the role of boron is in preventing ineffective
densification methods and increasing self diffusion coefficients through the creation of
vacancies [77]. Ermer utilized advanced characterization techniques to determine the
role of boron was to enhance grain boundary diffusion while carbon served to form
inclusions which greatly slowed grain growth [78]. Since the additives are believed to
improve sintering through surface reaction and vacancy creation routes, SiC containing
boron and carbon is generally known as solid state sintered SiC although some
researchers today question both the mechanisms through which boron and carbon serve to
densify SiC and if there is or is not a transient liquid phase formed in these compositions
[59, 79, 80]. The “lack” or transience of a liquid phase in these materials has important
repercussions in the fracture behavior of these materials as their grain boundaries are
clean, leading to a predominance of transgranular fracture in solid state sintered SiC [54].
It is believed that changing the fracture path to a more circuitous intergranular one may
absorb more energy.
Carefully controlling and modifying the composition of the sintering additives can
have a marked effect on the microstructure and properties. Alliegro processed SiC using
a wide range of additives including B, Be, Al, Ca, Li, and Cr. Small additions of Al were
observed to function as a solid state sintering additive much like boron. Above 1 wt. %
however, the addition of aluminum or aluminum compounds results in a transition from
solid state to liquid phase sintered SiC [76]. Zhang and De Jonghe processed SiC
33 containing aluminum foils and were able to denote the transition from clean grain boundaries showing transgranular fracture at low Al contents to increasingly intergranular fracture at high Al regions [81].
1.3.5.2. Liquid Phase Sintering Studies
Liquid phase sintering of silicon carbide is advantageous for multiple reasons and is the preferred method of processing SiC for many applications. Besides modifying the grain boundary fracture behavior, it also lowers the required sintering temperatures and times and can reduce the need for pressure-assisted sintering. This allows for more precise control of grain size and grain boundary composition as well as lower costs.
Certain mechanical properties, such as fracture toughness, can be enhanced using liquid phase sintering [82]. Other properties however, such as hardness, tend to suffer.
Controlling this tradeoff between mechanical properties and finding the correct correlation between properties are believed to be key steps in developing structural SiC.
A large amount of research has been done on liquid phase sintering SiC using yttria and aluminum based sintering aids [15, 83-89]. Although yttrium is not technically a rare earth, its low cost, availability, similar behavior, and 3+ valence makes it very
similar to many rare earths. As such yttrium has been established as a model material for
rare earth and liquid phase sintering studies. Huang et al. systematically characterized
SiC with AlN-Y2O3 additives. The work showed that starting with α-SiC was necessary
in order to develop fine equiaxed microstructures. High resolution TEM characterization
and electron energy loss spectroscopy (EELS) revealed the presence and composition of
34 intergranular phases and established the importance of the liquid phase in microstructural development and properties [90].
Liquid phase sintered materials can be modified further through the addition of other rare earth elements [14, 16, 91, 92]. Besides creating a more refractory grain boundary glassy phase suitable for high temperature applications, many of the mechanical properties can be altered by controlled additions [31]. Becher et al. did an extensive study on the segregation of specific rare earths to crystallographic surfaces in
Si3N4 and the subsequent effect of these elements on the grain growth [11]. Trace rare
earth impurities in Al2O3-5 % SiC composites increased the fracture toughness through crack deflection encouraged by intergranular fracture. Chemical analysis showed that the
fracture change was influenced by the segregation of rare earths to grain boundaries [93].
Terao et al. examined the effect of Sm2O3, La2O3, and Y2O3 on AlN and found Sm and
La increased the fracture toughness and changed the fracture mode. The increase in
fracture toughness was attributed to strengthened grain boundaries [94].
Zhou showed that LPS-SiC using Al2O3 and certain rare earth oxides displayed
different mechanical properties even though the microstructures appeared generally
similar [14]. Balog liquid phase sintered SiC using combinations of Y, Yb, and Sm
additives and found similar microstructures after low temperature sintering. Upon further
annealing however, the different combinations of additives showed varied grain growth
and noticeable trends in hardness and fracture toughness relative to the choice of rare
earth and its ionic radius [95, 96].
35
Chapter 2
Method of Attack
The fundamental goal of this thesis is to pursue and develop advanced processing methods capable of yielding silicon carbide with features and properties encouraging potential application as a structural ceramic. Secondarily, a major goal of this thesis was to gain insight and understanding into the influence of processing on the microstructure and grain boundary development of SiC and the related influence of the different microstructures on the performance of SiC. In order to achieve these goals, three objectives have been established. Each objective will be covered in its own chapter as outlined below. The general literature review of Chapter 1 is applicable to all sections while specific and relevant literature to each objective can be found at the start of each remaining chapter. Brief conclusions are described in each chapter with an overall summary and future work suggestions in Chapter 6.
Objective I: Develop a novel processing method for controlling the microstructure and grain boundaries of SiC that allows for reduction in large defects.
It is hypothesized by this author that the pathway to understanding the grain boundary behaviors of SiC can not be studied without reducing the population of large and anomalous defects in order to remove their overwhelming effects on mechanical properties and behaviors. This necessitates the addition of sintering aids in a well-mixed and reactive way to enable the best sintering and most control over the microstructure.
36
In addressing Objective I, coprecipitation will be developed as a method to introduce grain boundary phases into silicon carbide while improving the mixing of sintering aids and reducing the size and frequency of anomalous defects. Initial work will introduce coprecipitation as a simple method of adding sintering aids on a fine and reactive scale. Control of agglomeration, solvent, mixing speeds, time, and other dispersion characteristics will be investigated as methods of improving the coprecipitation procedure versus previous research. The thickness and uniformity of the coating will be evaluated through zeta potential measurement, TEM, and BET.
After an effective coprecipitation procedure has been established, a series of samples will be processed through coprecipitation, conventional wet ball milling, and aqueous propeller mixing with the goal of demonstrating differences in the mixedness, microstructure, and selected mechanical properties. Compositions containing 5 weight percent aluminum oxide and yttrium oxide sintering aids with a molar ratio of 4 Al:Y will be processed using aqueous ball milling and propeller mixing. The same final composition will also be achieved through coprecipitation starting with nitrate precursors.
A fluorescence lifetime measurement method will be employed to determine the mixedness of these samples versus standard processing methods. EDS analysis in the
TEM will be used to probe the mixedness of the sintering aids.
Objective I and its tasks are discussed in Chapter 3 of this thesis.
37
Objective II: Understand the Influence Coprecipitation has on Phase and
Microstructure
The second objective of this thesis is to better understand the influence the coprecipitation procedure has on the control and development of the microstructure of silicon carbide. The chemistry of the grain boundary phase will be controlled to explore the effects of compositional changes on boundary thickness, polytype, crystallinity, and grain shape, as well as other features. Coprecipitated samples will be processed with a number of different rare earth elements of varying cationic radius, namely yttrium, gadolinium, samarium, and lanthanum. These elements were chosen based upon their cationic radii differences and ability to modify the thickness and properties of glassy grain boundaries in SiC systems. Initial samples will further explore the sinterability of the aluminum and yttrium systems through modified hot pressing cycles. Additional hot pressed samples will be attempted for the gadolinium, samarium, and lanthanum systems.
Spark plasma sintering will be used to improve the sintered densities, the densification rate, and reduce the grain size of coprecipitated systems. Coprecipitated samples in the aluminum / yttrium system at both low (1 wt. % sintering aids) and high
(10 wt. % sintering aids) will be processed to investigate the sinterability of these systems relative to additive contents. The 10 wt. % samples will enable easier detection of minority crystalline phases and additive distribution, and a second comparison to standard processed samples will be made. A larger matrix of samples consisting of 2 and
5 wt. % additives for all four additives will be processed to high densities through spark plasma sintering. These samples will be characterized through standard scanning electron microscopy and x-ray diffraction. Chapter 3 will investigate the differences in
38 microstructure, phase, and grain boundaries between the coprecipitated and standard silicon carbides and indicate the important additive, additive content, heating method, schedule, and other choices that impact the processing and properties of LPS-SiC.
Objective II and its tasks are discussed in Chapter 4 of this thesis.
Objective III: Evaluate the Mechanical Properties and Plasticity behavior
The final objective of this thesis is to evaluate the elastic modulus, hardness, and plastic deformation response of the silicon carbide samples prepared throughout this thesis. The elastic modulus values for the samples will be determined from ultrasonic testing, and compared to the expected values. Knoop hardness will be measured at low loads of 100 grams, 300 grams, 500 grams, 1000 grams, and 2000 grams for all samples.
For a selected number of samples, additional higher load indents of 5000 and 10000 grams will be measured. The data from the high load indentation experiments will be evaluated in order to determine plasticity rankings for these materials using a method developed by McCauley et al. [97]. The indentation behavior will be examined as a method of interrogating the cracking and grain boundary changes developed through coprecipitation and other process changes.
Objective III and its tasks will be discussed in Chapter 5.
39
Chapter 3
Improved Coprecipitation Processing as a Method of Reducing Defects in SiC
3.1 Introduction
Chapter 1 introduced the basic structure and properties of silicon carbide and outlined the typical processing steps involved in developing high density SiC structural ceramics. Because of the limited sinterability of pure SiC, emphasis was placed on discussing the key powder processing steps that introduce and distribute the sintering aids. A general overview of the basic sintering process as well as innovative sintering methods and the connections between processing and sintering were included.
In this chapter, the impact of defects on the microstructure and mechanical properties of structural ceramics as well as the connection between processing and defects is discussed. Coprecipitation is introduced as an innovative and versatile processing method for reducing the population of defects in high density SiC ceramics while concurrently offering the ability to develop interesting grain boundaries and microstructures. A simple coprecipitation method was developed from literature, but results on defect removal and densification were poorer than expected. Therefore, a number of process improvements were instituted to modify the coprecipitation process.
Comparisons of microstructure relative to defects and grain size were made between conventionally processed samples, i.e. ball milling and wet mixing methods, versus samples yielded from the improved coprecipitation process. Finally, fluorescence intensity and lifetime measurements are offered as a more quantifiable measure of mixing improvements and defect reduction.
40
3.2 Processing – Structure – Properties Relationship
3.2.1 Impact of Defects on Mechanical Properties
As for all brittle ceramics; silicon carbide initiates failure at flaws present in the material. The importance of flaws in brittle failure was explained by Griffith in the
1920’s [4, 5, 98]. Griffith’s calculations outline the conditions for brittle fracture under quasi-static or low strain rate conditions and conclude that a ceramic will initiate fracture by cracking from the largest defect oriented favorably relative to the applied stress [5].
While this result indicates the importance of reducing flaws in engineered ceramics, the dynamics of certain applications of silicon carbide, like ballistic impact, may reveal a more critical situation than Griffith’s theory indicates. Higher strain rates and greater loadings can activate many or all flaws above a certain size leading to diffuse cracking
[99, 100]. The activation of many flaws has indeed been observed in brittle glass materials subjected to shock impact. Bourne observed a failure front rapidly proceeding behind the shock wave. The front appeared to nucleate at numerous locations in the glass and the front spread accordingly by cracking and coalescence. It was stated that the initiation sites were most likely flaws in the material although they were unable to establish whether the defects were microscopic or macroscopic. These results underscore the important role of improved processing in enhancing performance of SiC in structural or high performance applications [101].
Dutta tested the strength of four different silicon carbides and found a wide distribution in the Weibull data for strengths among the samples for each material [7].
Although different manufacturing techniques improved the strengths among the different
SiC materials, the high spread in the reliability of the data indicated a lack of
41 improvement in flaw distribution [7]. These results, shown in Figure 3.1, underscore the presence and importance of flaws in SiC, particularly in creating batch variability in large production lots.
Figure 3.1: Strength distribution and variability in commercial SiC materials showing high variability in fracture strengths. The wide distribution in strengths is indicative of a strong dependence on flaw size and distribution. [7]
3.2.2 Influence of Processing on the Population of Defects
The control of powder properties such as particle size and distribution, shape,
surface, and other factors all influence the processing and densification as well as the
number and distribution of defects in the final sample. Abnormal grain growth can be
limited through improved mixing of sintering aids, impurity control, phase transformation
control, and narrow particle size distributions [13, 102, 103]. Foreign inclusions can be
reduced by improving cleanliness of the processing steps, procurement of higher purity
starting materials, and powder washing and treatment. Many of these improvements can
and have been addressed to varying degrees by powder manufacturers and processors.
42
These steps have yielded improved materials over previous ones, but many sources of potential defects still exist.
Poor mixing and distribution of sintering aids is one of the most prevalent and continuing sources of defects in silicon carbide. Due to the highly covalent nature of bonding in SiC, sintering aids are necessary in order to activate effective diffusion mechanisms and produce a dense product. Choosing an effective sintering aid for the system, as well as the correct amount, is essential in achieving high densities. Excess sintering aids, agglomeration, and segregation of additives lead to density fluctuations, abnormal grain growth, and subsequent defects in the final sintered product [6, 39, 98].
Flaws are most commonly present as porosity or cracks introduced during the processing of ceramics [104]. This definition is incomplete however, as large grains, second phase inclusions, unreacted sintering aids, and agglomerates all function as stress concentrating flaws which lead to cracking and failure in engineered ceramics [17].
Bakas et al. performed scanning electron microscopy on ballistic rubble and found foreign inclusions, grains much larger than the average, poorly mixed sintering aids such as carbon, and other defects present on fracture surfaces. Figures 3.2 and 3.3 show carbon inclusions in the microstructures of Bakas and Dutta [6, 7].
43
Figure 3.2: Carbon inclusion (C) and surrounding interface grains (I) on fracture surface of SiC [6]
Figure 3.3: Replica micrograph showing carbon inclusions at triple points and grain junctions [7]
Baron et al. tested four similarly sized nano alumina powders from various
producers and observed that their green properties, namely density and strength, were
influenced by green processing changes and impacted the final sintering behavior and
properties. Three of the powders were similarly prepared while the fourth contained a
binder added during manufacturing. It was first observed that the presence of the binder
44 lead to the formation of hard agglomerates which could not be removed during typical sintering procedures, leading to lower densities than the other three powders. After a calcination step was added to remove the binder, the density of this powder returned to the same level as the other three. All of the powders were also subjected to attrition milling and freeze drying processes which further modified their behaviors. The reduction in particle sizes from attrition milling reduced the green density while the freeze drying process tended to create hard agglomerates which increased the green strength. These hard agglomerates impacted the sintering again however, leading to undesirable microstructural growth and lower densities [105].
Hamminger and others studied the fracture surface of SiC and found that carbon inclusions acted as flaws in the microstructure [7, 39, 57, 75, 98]. Other researchers have observed a correlation between the presence of defects like pores and lower fracture strength [31]. It is clear that improvements in processing are essential in reducing the population of flaws in SiC ceramics.
3.3 Improved Processing Methodology
The two sections above on the importance of defects in determining the mechanical properties of SiC and the linkage between defects and processing indicate that better processing methods may be necessary to achieve the properties desired of the next generation of SiC ceramics. Chapter 1 outlined the typical powder processing route for sintered SiC ceramics and also highlighted some of the previous research aimed to process and sinter interesting SiC materials. Here we are going to explore a possible
45 method for preparing silicon carbide, namely coprecipitation, which can reduce defects and allow for control for many features of the microstructure.
3.3.1 Coprecipitation
Chemical precipitation, where an insoluble solid phase is nucleated from a supersaturated liquid through any number of methods, has been well established in the production of many chemicals and materials. A single insoluble salt or combination of salts can be nucleated out of solution and grown into a range of particle sizes, shapes, and characteristics, making this type of processing highly controllable. Precipitation can occur at very rapid rates which can lead to very fine nanoscale particles that often create materials with improved mixedness, reactivity, and differing crystallinity [106-109].
Although there are many specific methods by which to perform a coprecipitation, a general description of a common procedure, without considering the development of a coating, is described here. The initial step in coprecipitation is the addition of a precursor salt into a suitable solvent where the precursor will dissolve into its ionic components.
The ions will then attach to the polar solvent molecules and surround them, a process known as solvation. An isolated species in solution will develop an ionic charge due to dipole effects, if not already charged. Oppositely charged ions will be initially attracted to the colloid forming a well defined layer of counter charged ions on the surface. With this initial layer present other oppositely charged ions are now somewhat repelled at a certain distance, but still feel the charge of the colloid strongly enough to diffusely surround it. As the accumulation of oppositely charged ions surrounds the colloid, they serve to shield the charge of the colloid from remaining ions in solution, allowing ions of
46 the same charge as the colloid to begin surrounding the outer layer. Accordingly, the number of ions of the opposite charge of the colloid will decrease with distance from the center while ions of the same charge will accumulate more towards the outside. This is known as the formation of a double layer shown below in Figure 3.4. This solvation process and formation of a double layer can result in ions repulsing each other, leading to good separation and mixing of the components throughout the solvent [109-111].
Figure 3.2: Formation of electrical double layer due to charge behavior of ions in solution. Curve indicative of changes in zeta potential with distance and ion adsorption near charged surface. Compiled from [110, 111].
If a second mixture containing another solvated ion pair is added to the first there is a possibility of precipitating a solid phase through a double replacement reaction, if one of the pairs of cations and anions combine to form an insoluble product. An example equation for a double replacement reaction is shown below in simplified form:
Al(NO3)3 (aq) + 3 Na(OH) (aq) ÆAl(OH)3 (s) + 3 Na(NO3) (aq) (Equation 5)
47
In more complicated systems such as during coprecipitation, the two possible reactions are happening concurrently over the pH ranges where they are favored. The formation of complexes and other side reactions can complicate the chemistry.
The insoluble solid phase will nucleate and grow rapidly from solution under the correct conditions. There are many essential variables to consider when controlling the coprecipitation process including pH, temperature, saturation, rate of addition and many others. Through proper control of these variables, coprecipitation can be scaled or modified to influence many important features of the final product including particle size, purity, agglomeration, and many other features [106, 107, 109, 111].
The behaviors of particles in solution are described by colloid chemistry. It is well known that careful control of the surface characteristics of the powders such as the surface area, charge, and species present, as well as the conditions of the solvent, can be used to disperse, agglomerate, or otherwise manipulate the distribution of colloids.
Similar controls can enable coprecipitation to be utilized in more complicated ways such as in coating a powder from colloidal solution [107, 109].
The behavior of colloids in solution can be described in terms of the simplified
Derjaguin, Landau, Verwey and Overbeek (DLVO) theory and extended to describe the procedure of coprecipitation and coating. According to DLVO theory, the tendency of particles in solution to either adhere to or repel each other is governed by the competition between the attractive Van der Waals forces and the electrostatic repulsion from the double layer. As two particles approach each other in solution the interaction of their double layer provides an electrostatic repulsion, keeping them from associating.
Depending upon the factors that govern each force, such as bond strengths, charge, etc, a
48 curve can be drawn showing the total energy of the system and the contribution by each of the two separate forces. If the repulsive energy is sufficiently high, the ions will prefer to stay well dispersed. If the attractive energy is more favorable, the system will prefer to agglomerate if containing a single species or to have opposite charges associate with each other, e.g. form a coating. A graph indicating the competing forces is shown below in
Figure 3.5. Depending upon the shape and height of the total energy curve the normal collisions between particles may provide enough energy to encourage ions to stick to each other either temporarily or permanently. In many cases however, the use of surface modifying polymers, ionic strength, or other effects may be necessary to introduce the desired behavior[110-112].
Figure 3.5: Net interaction curve showing competition between repulsion and attraction that governs agglomeration or dispersion of particles in solution [111]
Coprecipitation has been shown to be a versatile and effective method of coating a base powder with fine nanoscale additives. Research on Si3N4 has revealed improved
homogeneity and mechanical properties with coprecipitation of sintering additives [113,
49
114]. Although much research has been done on coprecipitation for Si3N4, little has been
done on utilizing it as a method for introducing sintering aids for SiC.
Bellosi et al. found improved mechanical properties and sintering kinetics in SiC
with coprecipitated additives of Y and Al [115, 116]. The study compared
coprecipitation with ultrasonication and concluded that coprecipitation increased the
densification rate and lowered the final grain size. The amount of amorphous phase was
also observed to decrease from 6 to 3 wt. % through the use of coprecipitation. This was
likely due to the reduction of silica content and increased densification rate from the fine,
reactive sintering aids. While the mechanical properties did show some improvement,
the work implied that the strength and properties of these coprecipitated samples were not
as high as expected due to the observation of large porous regions, shown below in
Figure 3.6. These issues were likely due to poor processing procedures relative to
agglomerate breakdown and dispersion of the particles during mixing. This thesis will
attempt to remedy some of these issues.
Figure 3.6: Highly porous region in liquid phase sintered SiC prepared through coprecipitation. Region caused from lack of sintering aid penetration due to agglomeration of starting powder [116]
50
It is believed the fine, homogenous distribution of sintering aids obtained from coprecipitation will reduce the number of defects in SiC and lead to enhanced performance.
3.4 Procedures:
3.4.1 Initial Coprecipitation Procedure
3.4.1.1 Raw Materials
3.4.1.1.1 Silicon Carbide
As one of the goals of this thesis was to develop homogenous microstructures with a reduced number of defects, an α silicon carbide powder was chosen. This was aimed at limiting the formation of the large, acicular grains that develop from the β to α transformation that occurs at higher sintering temperatures when using β starting powders. The α powder used in this study was H.C. Starck UF-25 (H.C. Starck GmbH,
Goslar, Germany). This powder is a fine grained powder produced by the Acheson process. The manufacturer reported it consists primarily of the 6H polytype and has an average grain size of 0.45 microns and a BET surface area of 23-26 m2/g. The chemical
analysis provided by H.C. Starck indicated a free carbon content of 1.27 wt. %, 1.86 wt.
% oxygen, 0.03 wt. % aluminum, 0.03 wt. % iron, and less than wt. % 0.01 calcium. An
effort was made to more carefully characterize the polytype distribution and particle size
of the powder. This is discussed later in the Results section.
51
3.4.1.1.2 Sintering Aids
The sintering aids for this study were chosen from alumina and rare earth oxide systems with the intermediate products added as a fine coating through coprecipitation.
The discussion here covers the selection and initial properties of the precursor powders.
Four rare earths, shown in Table 3.1 below, were selected on the basis of covering a wide range of cationic radii in order to modify the grain boundary properties. The concept of modifying grain boundary behavior through sintering aid and cationic radius control was mentioned briefly in Chapter 1; it is discussed at length in Chapter 4. Aluminum and rare earth nitrates were chosen as the precursor powders for this study due to their high solubility in water and alcohol and the relative ease of burnout for their residual organics.
Table 3.1: Ionic radii of selected rare earth elements, from Langhorn [117] Rare Earth Ionic Radius(Å) Y 1.019 Gd 1.053 Sm 1.079 La 1.16
The starting powders used in this study were yttrium nitrate hexahydrate
(Y(NO3)3·6H2O, Alfa Aesar, Ward Hill, MA), aluminium nitrate (Al(NO3)3·9H2O, Acros
Organics, NJ), lanthanum nitrate hexahydrate (La(NO3)3·6H2O, Alfa Aesar, Ward Hill,
MA), and gadolinium nitrate hexahydrate (Gd(NO3)3·6H2O, Alfa Aesar, Ward Hill, MA).
A small number of comparison samples utilized alumina and yttria added directly as
oxide powders. The alumina was A16 SG (Alcoa). The yttria was from Pacific
Industrial Development Corporation (PIDC).
52
3.4.2 Composition Selection
Compositions were chosen on the basis of sintering aid systems based upon Al2O3 and a single RE2O3 chosen from the group of yttrium, lanthanum, gadolinium, and
samarium. The weight percentage of final oxides was established as well as the molar
ratio of alumina to the rare earth oxide. The procedure for introducing the additives was
coprecipitation of the aluminum hydroxide and rare earth hydroxide components out of
solution onto the SiC grains as a fine coating. Before utilizing coprecipitation as a coating method, evaluation of the behavior of the SiC powder and individual components was required.
3.4.2.1 Titration Curves
Before any coprecipitations were performed for samples containing SiC, simple titrations of each of the single components were performed in order to identify the pH ranges at which each of the components was expected to precipitate. This was important to insure that each of the systems would be possible and to establish the necessary pH conditions to run the reactions. Analysis was performed in water initially as to compare to literature and a second analysis was performed in isopropanol for a number of the systems to ensure its success in that solvent as well. The nitrates were dissolved in 50 mL of the solvent and stirred continuously in a beaker. A second solution of NH4OH was prepared at a high concentration and was added to the beaker at the same approximate volume per drop. A pH meter was used to continuously track the pH.
These titrations are shown below in Figure 3.7. The curves established that the aluminum hydroxide or complexes came out of solutions at a pH of between 6 and 7,
53 while all of the rare earth oxides had slightly different but similar ranges of between 8 and 9. This indicated that a pH of above 9.0 would be most effective for simultaneously coprecipitating alumina and rare earth hydroxides controllably for good coating.
Figure 3.7: Titration curves for selected rare earth systems
3.4.3 Coprecipitation Procedure
The versatility of the coprecipitation process allows for easy scalability of the
batch size and throughout this thesis, sample batch sizes ranging from 10 to 1000 grams
of powder were prepared. Figure 3.8 is a simple flow chart indicating the common steps
in the coprecipitation procedure used here; some of the process improvements
implemented throughout the work are not listed for simplicity. The silicon carbide was
measured out and added to a beaker containing water or isopropanol as a solvent.
Enough solvent was added to make a dilute and fluid slurry. One weight percent of
Triton-X 1000 (Sigma Chemical, St. Louis, MO) was added to the slurry and it was
mixed by propeller (Cole-Palmer Instrument Co. Stir-pak Laboratory stirrer; model#
54
4554-00) for 30 minutes before the coprecipitation was began. An electronic pH meter,
(Fisher Scientific Accumet Model 25 pH/ion meter) was carefully situated in the slurry to monitor the pH at this time and throughout the coprecipitation process. The necessary amount of aluminum nitrate required to yield the calculated alumina content of the sample was measured out and dissolved in the solvent. The necessary amount of the rare earth nitrate was then added to the solution and allowed to dissolve and mix until the solution was clear. The required amount of ammonium hydroxide (Fisher Scientific,
Pittsburgh, PA) needed to precipitate all of the added nitrate components was calculated.
For some earlier samples, a solution of tetramethylammonium hydroxide pentahydrate
(TMAH) (C4H12NOH·5 H2O, Fisher Scientific, Fair Lawn NJ) was used because of its effectiveness as a dispersant for SiC [61]. Either hydroxide was then dissolved in solvent and mixed, separate to the first mixture. The nitrate solution and the hydroxide solution were each poured into separate burettes. The hydroxide was added slowly to bring the
system up to a pH of 9.5; then both solutions were added dropwise to the beaker
containing the SiC as it rapidly stirred while carefully controlling the rates of addition to
keep the pH as close to a constant 9.5 as possible. As the coprecipitation neared
completion, the rate of addition was slowed and the pH level was raised to 10.0. After all
of each solution was added to the beaker, the contents were aged for 1 hour at a pH of
10.0 at room temperature of 20°C.
55
Figure 3.8: Example flow chart outlining steps in coprecipitation process
After the aging process, the powders were washed in order to aid in the removal of the ammonia and nitrates. The slurry was allowed to settle and the supernatant was poured off. Fresh solvent was then added to the beaker and the contents were mixed by propeller for 30 minutes. This process was repeated three times and the powder was subsequently dried in an oven or pan dried.
For some earlier batches, samples were dry pressed in a stainless steel die at this point. The densification results of these samples were poor however, and it was suggested in literature that the samples be fired to low temperature in an inert atmosphere to further remove the residual organics [115, 116, 118]. The powders were loaded into alumina crucibles and placed in a tube furnace and sealed. The samples were then fired under an Ar atmosphere at a heating rate of 2 degrees per minute to 500°C and held for
56 one hour. After cooling, the powders were crushed by mortar and pestle and screened through No. 200 (74 micron) and No. 325 (44 micron) sieves. For spark plasma sintering, the powders were stored in a Nalgene bottle and left in an oven overnight before sintering. When hot pressing, the fired powders were dry pressed into 2.25 inch or
1.00 inch discs in a stainless die at 35 MPa.
3.4.4 Coprecipitation Process Improvements
After the initial coprecipitation provided some low and fluctuating density samples, a series of improvements were made in the process to ensure better homogeneity and control. This section describes the procedure for performing some of them. Their development and effectiveness are discussed in the Results sections to follow.
3.4.4.1 Milling Treatment
The majority of coprecipitated samples in this study were subjected to a milling treatment to reduce agglomeration and particle size before coprecipitation. This milling hoped to develop a more uniform dispersion of SiC before the coprecipitation. Varying sample sizes were explored during this thesis so the sizes and amounts listed below are a single example. 150 g propylene glycol (PG) (Fisher Scientific, Pittsburgh, PA) were added to a 1L milling jar (Nalgene, Rochester, NY) followed by addition of 300 g of 0.19 inch diameter SiC milling media (Union Process, Akron, OH). Pristine SiC powder, 30 g, was added to the jar which was sealed and tumbled for 2 hours. The media was then separated from the suspension by sieving and the suspension was allowed to settle. The
PG was then removed by siphoning, and fresh isopropanol (Fisher Scientific, Pittsburgh,
57
PA) was added, followed by 10 minutes of mixing by propeller. This washing procedure was performed 4 consecutive times. After the final settling, the isopropanol was removed by siphoning and the milled and washed slurry was ready for use in coprecipitation.
3.4.5 Comparison Samples
In order to establish if the coprecipitation process did indeed lead to a reduction in defects, samples were prepared from simple mixing of powders as a comparison. SiC,
Al2O3, and Y2O3 were weighed out to match the final oxide composition of a
coprecipitated sample. Some batches were then dispersed in isopropanol to a moderate
solids loading in order to ensure good mixing. The powders were ball milled for 3 hours
with 3/16” SiC media, sieved, dried, and ground by mortar and pestle. Aqueous mixing
of powders using the same propeller as when coprecipitating was also explored. The
powders were added to a beaker with isopropanol and Triton-X 100 in the same way as
the coprecipitation, and stirred at the same speed setting as the coprecipitation solutions.
3.4.6 Note on Sample Naming and Nomenclature
As the nature of this thesis focused on the development of an effective powder
processing procedure and included a large variety of additives and sintering methods, a
considerable number of batches of powder were required. Batches of different sizes,
modified methods, and varied additives were processed during the stages of this work. In
the pages that follow, the exact composition of the sample is often listed with a secondary
designation of the letter R followed by a number, i.e. 54. The letter R indicates this is
from the specific series of samples interested in rare earth doped SiC. The number is
58 present in order to identify this specific batch of powder. In a number of cases, multiple powder batches were prepared for the same composition. The numbers have no other specific meaning besides indicating the batch and do not identify the percentage or the additive. Many of the prepared batches were made larger than required for a single sample firing. In these cases, a second number was added to differentiate the batch being split, i.e. R43-1, R43-2. Not all of the numbered batches contributed to the work discussed, so a gap in the number sequence does not indicate any specific problems.
Care was taken throughout the process to document the conditions and perform the procedure consistently in order to limit batch to batch and sample to sample variability.
3.4.7 Sintering Procedures
3.4.7.1 Hot Pressing
Samples for hot pressing were stored in a drying oven and checked for dimensions and weight before loading. Sample sizes varied throughout the study in order to conserve powder and address sample loading issues. Larger samples, measuring approximately 2.25 inches in diameter, were slid into a 2.25 inch ID graphite sleeve
(Weaver Industries, Denver PA) after surrounding the inner section of the sleeve with either graphite foil or painting the tube with a boron nitride (BN) slurry or spray. The samples were also surrounded with graphite on the top and bottom and sandwiched between two graphite discs that acted as spacers. Smaller samples, 1 inch diameter, were prepared using the same sleeve because of alignment issues, but the samples were instead surrounded with BN powder. The sleeve was secured in a 6 inch OD by 5 inch ID by 4 inch high graphite die. This assembly served as the workpiece for the induction heating.
59
The graphite die and sleeve assembly were then placed on the stage of the induction hot press. The stainless steel rams of the hot press were controlled by hydraulics allowing their movement and positioning. The bottom ram was locked in position through the use of an isolation valve. A graphite rod was then placed onto the stainless steel ram while the other end of the rod is slid into the bottom of the sleeve and into contact with the spacer. The sample and spacers were adjusted to make sure they were properly located in the hot zone of the furnace. The top graphite rod was then lowered into the sleeve. An insulation covering made of woven layers of graphite felt was slid down the rode and onto the top of the induction coils and die set. The upper stainless steel ram was then lowered into position manually, and the alignment of the entire die and graphite rods was adjusted into the best alignment relative to the rams.
Hot pressing was performed in a 30 ton uniaxial induction hot press
(Centorr/Vacuum Industries, Nashua, NH). The induction coil set was further insulated with a layer of mica and woven graphite felt sheets. After the die and plungers were loaded onto the hot press stage and aligned to the rams, the system was ready for evacuation. A rough vacuum was pulled on the chamber until an approximate vacuum of
10-2 Torr was reached. The system was then backfilled with argon, and evacuated again
in order to better reduce the residual oxygen levels. The system was then switched over
to diffusion pump and the system was pumped down to a vacuum of 10-3 to 10-4 Torr.
Although a better vacuum should be possible with a diffusion pump, the age and condition of the system prevented better ultimate vacuums from being reached.
Heating in the hot press was controlled in two separate regimes because the temperature was pyrometer controlled. As the pyrometer only measured temperatures
60 above 900°C, the power output of the high frequency induction power supply was manually controlled to heat the system to a readable temperature. When the digital readout of the pyrometer registered temperatures above 1000°C, the control of the induction power output was switched from manual control to the automatic furnace control mode. The system was continuously under vacuum during the manual control range in order to remove organics and water below 1000°C, but was quickly backfilled at the change to automatic control in order to prevent the loss of sintering aids due to possible volatilization during the remaining heating.
Because of differences in the sinterability of the various rare earth additions and compositions, a number of different sintering cycles were utilized. The cycle presented here is a representative cycle from which the changes in later cycles can be identified.
The heating rate was 30°C/ minute up to 1000°C, 10°C/min from 1000°C to 1400°C, 5°
C/min from 1400°C to 1700°C, and 3°C/min from 1700°C to 1900°C with a final hold of
1 hour at 1900°C. After the final hold, the power was shut off and the furnace was allowed to cool at its own pace. The pressure cycle was applied at 1400°C by slowly and manually ramping the applied load. The final pressure was set at 35 MPa and was reached in approximately 2 minutes. Modifications to the cycle above are noted for any relevant samples in the later sections.
3.4.7.2 Spark Plasma Sintering
Spark plasma sintering runs were performed using a 10 ton SPS unit (Model 10-4,
Thermal Technologies, Santa Rosa, CA). Loading samples for spark plasma sintering consisted of wrapping the inside of a 20 mm graphite die with a small sheet of graphite
61 foil and sliding the bottom punch into position. Two small circles of graphite foil were then placed on top of the punch in order to prevent reaction and bonding with the bottom punch. Between 4 and 5 grams of powder were then poured into the die, and another two circles of graphite foil placed on top. The top ram was then inserted and the die assembly cleaned was cleaned with isopropanol. The die was wrapped with an insulation layer of graphite felt with a small hole cut into it. This hole is made to allow for the pyrometer to be sighted on a small hole machined in the die to a depth of 0.75 inches of the distance through the width. This allows for more accurate temperature determination.
The SPS configuration used here consisted of two hydraulic rams, a moveable top ram and a stationary bottom ram. A cold ram spacer was placed on the bottom ram, followed by a secondary spacer. The die was then placed on the secondary spacer. The pyrometer is then carefully aligned to the hole in the die/insulation by rotating and adjusting the die to aiming crosshairs visible on the screen connected to the pyrometer.
This alignment is essential to ensure consistent and accurate temperature reporting.
After the samples were properly loaded, the door was sealed and the chamber evacuated via rough pumping until a vacuum of 10-2 Torr was reached. The system was
then backfilled with argon, and evacuated again in order to better reduce the residual
oxygen levels. Once the appropriate rough vacuum was reached, the turbo pump was
activated and the system was pumped down to a high vacuum of 10-4 to 10-5 Torr. Once
the system passed the high vacuum checks, the interlocks allowed the power system to be
activated.
The heating in the SPS is controlled by a programmable Eurotherm controller.
The cycle described below is again a general cycle that was modified for specific sample
62 runs. The temperature is monitored through a pyrometer that begins reading at 600°C, so the initial stage in the program has a ramp of 200 °C/min with a hold requiring the pyrometer to register at 600 °C for 30 seconds before proceeding. The program then continues as designed, with a ramp rate of 200 °C /min to the final sintering temperature of 1900 °C. The hold time at the final temperature varied between 5 and 15 minutes. For a number of samples, a modified cycle with a 200 °C /min ramp to 1800°C with a 5 minute hold, followed by a 200°C/min ramp to 1900°C and a final hold of 5 to 10 minutes. The pressure was applied at 5 MPa from the start of the run, and ramped up to
50 MPa during the initial heating steps. The power was quickly shutoff at the ultimate temperature allowing for rapid cooling, while the pressure was ramped down to 5 MPa.
3.5 Characterization Techniques
3.5.1 BET
BET (Brunauer, Emmett, Teller) analysis was utilized to determine the specific surface area of some of the powders during various stages of processing to determine if the coating was indeed present. The analysis was performed using a Micromeritics
Gemini 2375 (Micromeritics Instruments, Norcross, GA) through the absorption of nitrogen. Samples were degassed overnight at 110oC prior to surface area analysis.
3.5.2 FTIR
Fourier-transformed infrared spectroscopy (FT-IR) analysis was performed in
transmission using a FTIR spectrometer (Galaxy Mattson 5000) via KBr pellets
63 technique. Pellets were prepared by mixing 0.2mg SiC-based powder in 100 mg KBr
(FTIR grade, Alfa Aesar).
3.5.3 Electrokinetic analysis / zeta potential measurement
The electrokinetic measurements were carried out on a zeta potential analyzer
(Brookhaven Zeta PALS, Holtsville, NY) using suspensions of 0.01g/mL SiC-based
-3 powders prepared in KCl 10 M. The zeta potential was measured over a pH range of 2-
10. From the data, the isoelectric point was determined for all samples.
3.5.4 Density
Density was measured by the Archimedes displacement method, following the
ASTM standard (B962). After sandblasting to remove the graphite foil that adhered to
the sample during sintering, all samples were carefully cleaned with acetone. Samples
were then allowed to dry. Five to ten weights were taken for each sample for dry, wet,
and suspended measurements. Samples were then boiled in distilled, deionized water
containing a few drops of Photoflo (to prevent air bubbles) for 1 hour. Suspended
weights were then taken by weighing the samples while they were suspended in distilled, deionized water. Care was taken to keep the samples in boiling water between measurements. A damp paper towel was then used to remove the water from the surface, and saturated weights were taken.
The theoretical densities of these samples were calculated using the weight percents of the silicon carbide, alumina, and rare earth oxides based upon the initial
calculations. This calculation therefore assumed complete addition and conversion of the
64 nitrate precursors, as well as ignoring the effect of volatilization during sintering. The amount of silica present on the SiC powders was also not considered. These considerations introduced some percentage of error in the theoretical density, and the actual density of the samples was always observed to be higher by microscopy than the calculation indicated.
3.5.5 Grinding, Cutting, and Polishing
Samples of a sufficient size were mounted on alumina tiles using Crystal Bond
590 (SPI Supplies, West Chester, PA) and were surface ground with a 325 grit resin bonded diamond wheel (North Jersey Diamond Wheel, Cedar Grove, NJ) on a Kent automatic surface grinder/cutter, Model KCF-1020 AHD, (Kent Industrial, Tustin, CA).
After grinding, smaller pieces were cut from the main sample using a Leco VC50 (St.
Joseph, MI) slow speed saw. Pieces were then cleaned with acetone and mounted in epoxy (Epomet F, Buehler, Lake Bluff, IL) using an Simplimet 1000 (Buehler, Lake
Bluff, IL) mounting press. Upon cooling the samples were loaded in a Buehler
ECOMET 3000 (Buehler, Lake Bluff, IL) autopolisher and polished on 125, 45, and 15 micron diamond embedded discs. Fine polishing then continued using 6 and 1 micron diamond slurries. For samples intend for microscopy, a final polish was performed using a 0.25 micron diamond slurry.
3.5.6 Etching
Various etching methods were employed in order to better observe different aspects of the microstructure. Chemical etching using Murakami’s reagent was
65 performed in order to remove the grain boundary phase and examine the grain size or other features of the SiC. The Murakami’s reagent was prepared by boiling a mixture of
NaOH and potassium ferricyanide.
Plasma etching was utilized on samples in order to preserve the glassy phase for observation. A plasma etching system (SPI Plasma Prep II, West Chester, PA) containing a CF4:O2 mix at a ratio of 92% CF4 and 8% O2 by weight (Air Liquide
America Specialty Gases, Houston, TX) was used. The etching time was variable
depending upon composition, but best results were typically observed with long etching
times of 60 to 80 minutes.
3.5.7 SEM
After grinding, polishing, and etching, where applicable, care was taken to
prepare the samples for scanning electron microscopy. Samples were removed from the
epoxy mounts through careful cutting on the diamond saw. Excess epoxy was ground off
on silicon carbide paper. Samples were then cleaned by ultrasonication in isopropanol
for 5 minutes followed by a second ultrasonic treatment in hydrogen pyroxide for 5
additional minutes. Samples were then mounted onto aluminum studs using carbon tape
or silver paint in situations where high voltage energy dispersive x-ray spectroscopy
(EDS) was intended. Samples were desiccated for at least 1 hour before loading.
Scanning electron microscopy was utilized to examine polished surfaces.
Microscopy was performed using a Leo/Zeiss DSM Gemini 982 FESEM and a Zeiss
Sigma FESEM (Carl Zeiss AG, Germany). Procedures varied slightly between the two
scopes; the description below focuses on the Sigma as the majority of the images were
66 taken on that microscope. Images were taken around the approximate zero point of charge of SiC, 2.7 to 3.0 eV, depending upon where the image appeared best relative to the different compositions. Various imaging methods were used throughout this thesis, ranging from secondary electron imaging using the Everhart-Thornley detector to in-lens images to backscattered electron imaging. The working distance was kept around 8.5 mm as this is the analytical working distance for the Sigma. Samples were examined to determine grain size, porosity, microstructural homogeneity, presence of defects, and the distribution of the grain boundary phase.
3.5.8 Grain Size and Porosity
The images of polished and etched microstructures were analyzed using a software package Lince V2.42B (TU Darmstadt,FB Materialwissenschaft, FG NAW) in order to determine the grain size and amount of porosity. The Lince software package is a convenient method that takes a calibrated image and allows for the semi-automated calculation of grain size through the linear intercept method.
3.5.9 Hardness
Microhardness testing was performed on a Leco M-400-G3 Hardness Tester
(Leco Corp., St. Joseph, MI) using a Knoop diamond indenter. Hardness values were obtained by measuring the length of the longer diagonal in microns. This value is then entered into the following equation for calculation of the Knoop hardness:
P HK = (Equation 6) C L2 p
67
2 where HK is the Knoop hardness in kg/mm , P is the load in kg, Cp is a correction factor
related to the indenter shape, 0.070279 for standard Knoop indenters, and L is the load in
mm. Hardness indents were made at loads of 0.100 kg, 0.300 kg, 0.500 kg, 1 kg, and 2 kg. Either 5 or 10 indents were taken at each load at a well polished region.
3.5.10 Ultrasonic Nondestructive Evaluation
Elastic, shear, and bulk moduli values, as well as Poisson’s ratio, were determined by ultrasonic nondestructive evaluation. The sound waves of a known frequency range can be passed through a material where they interact with the features present in the microstructure. The effect of the interactions on the amplitude of the signal upon exiting the sample can be recorded, as can the time of flight (TOF) through the material. The data can be taken at a single point for a quick representative measurement on a small sample, or the transducer can be rastered across the sample in order to develop maps of this information. From this data, many mechanical properties can be determined using simple relationships and equations, and this data can also be mapped. Some of the relevant equations from this study are shown below:
2 2 υ = [1-2(cs/cl)] / [2-2(cs/cl)] (Equation 7) 2 E = [(cl) (ρ)(1-2υ)(1+ υ)] / [(1- υ)] (Equation 8) 2 G = (cs) (ρ) (Equation 9) K = E / [3(1-2υ)] (Equation 10)
where υ is Poisson’s ratio, cs is shear velocity (m/s), cl is longitudinal velocity (m/s), E is
Young’s modulus (GPa), G is shear modulus (GPa), K is bulk modulus (GPa), and ρ is
the density (g/cm3)
Although it was originally desired to make full C-scan maps for the majority of the samples in this thesis, the small size of the samples prepared from the current SPS
68 set-up made C-scans impossible at this time; therefore A-scan, or point measurements, were then used to determine average moduli values.
3.5.11 XRD
Powder x-ray diffraction analysis was performed by filling a small plastic sample holder with the necessary amount of powder. Sintered materials were analyzed with samples that were surface ground or polished to remove any reaction layers and obtain a smooth surface finish. The sintered samples were then mounted into a bulk sample holder using sticky tack. Care was taken to keep the surfaces of the sample flat and parallel with the holder to prevent avoidable displacement errors.
A number of different XRD machines were used throughout this work because of varying equipment and resolution problems. Initial scans used a Siemens Diffractometer
D-500 with Cu Kα radiation and a nickel filter. Scans with an improved signal to noise
ratio were desired later in the work, and so scans were repeated or supplemented using a
Phillips PANalytical X’Pert system. The highest resolution scans were later taken on an
additional Phillips PANalytical X’Pert MPD Pro system.
3.5.12 Transmission electron microscopy (TEM)
TEM imaging of the coating quality was carried out on a Topcon 002B TEM at an accelerating voltage of 200 kV. The samples were prepared by dispersing the powders
in isopropanol and depositing a few drops on lacey carbon grids. TEM / EDS analysis of
bulk samples was performed on a JEOL 2010F Field Emission TEM operating at 197 kV.
69
Samples for this analysis were prepared from the H-bar technique using a focused ion beam system.
3.5.13 Characterizing Mixedness Changes
The reduction of defects in a microstructure through advanced processing is intended to improve the measurable mechanical properties and. Judging the extent of improvement however can be difficult because of how many microstructural features have influence over properties and the effect of subtle changes. Measurements such as strength and fracture toughness can provide valuable information on improvements in bulk sample behavior, but the appearance of the sample may not be particularly different with an improvement in additive mixing. Simple observation of differences in the microstructure through electron microscopy can be employed as a rough judgment of the presence of defects and the level of the mixedness of additives. The use of energy dispersive x-ray spectroscopy to quantify impurity elements, inclusions, or unwanted phases in images can give some indication of process improvements, but these techniques can be complicated in nanoscale systems or in low contrast regions. These subjective methods do not allow for evaluation of sample mixedness beyond a coarse level of the microstructure, and may not distinguish one method from another.
Researchers have designed methods for analyzing microstructure and mixedness improvements in more quantifiable ways. Computational modeling efforts have been employed as methods of evaluating and comparing differences in mixing; a summary of the relevant research work in the area can be found in Czerepinski et al [119].
70
Czerepinski and Riman also developed an innovative evaluation method using fluorescence spectroscopy specific for rare earth doped ceramics. Fluorescence in rare earth containing systems is well understood and can be used to evaluate the mixedness of ceramics because of concentration quenching. When rare earth ions are in too close proximity, they can exchange energy between each other instead of contributing to the radiative loss of energy as fluorescence. By measuring the fluorescence intensity and decay time, it can be inferred that a well mixed sample will display greater intensity and longer decay times because of lessened concentration quenching while a poorly mixed sample should have clusters of rare earths that are quenching. In this way, measurements of fluorescence intensities and times can be used as a probe of structure and mixedness in rare earth containing systems [119].
3.6 Results and Discussion
3.6.1 Processing
3.6.1.1 Coprecipitation
The initial goal of this section was to establish the effectiveness of coprecipitation as a method of providing a silicon carbide powder with a fine, reactive coating of sintering aids. It was expected that the improvements in processing, homogeneity, and scale of the sintering aids obtained through coating would reduce the population of anomalous defects in SiC. The removal of defects could enable the emergence of deformation mechanisms and behaviors governed by the grain boundary phase that are otherwise overshadowed by the presence of these larger features. The manipulation,
71 control, and scale of the additives during coprecipitation could also allow for greatly modified structure and properties of the grain boundary phase.
Early samples focused on identifying the essential variables required for successful coprecipitation of a coating onto SiC. The earliest samples were prepared from a very basic procedure; later work focused on developing an improved procedure.
All early samples were prepared from compositions in the aluminum-yttrium system because of the abundance of literature for comparison. The subsequent characterization work on the resultant powders was to establish the success and identify the changes introduced by the coprecipitation of a coating of sintering aids onto SiC. Examination of some fired samples was performed to establish further proof of success.
Small amounts of powder from before and after many of the coprecipitation steps were saved and examined in order to evaluate the method.
3.6.2 Initial Characterization
3.6.2.1 Zeta Potential
Zeta potential measurement is an effective method of determining the electrokinetic properties of the surface of a powder in suspension. The zeta potential is related to the overall charge a particle acquires in a specific medium. Oxides and hydroxide species have a common feature called the isoelectric point which is the pH where the zeta potential equals zero. The value of the isoelectric point however, is more or less specific, within certain limitations, to a given compound. Therefore, identification of the isoelectric point can serve as a useful reference in identifying the surface character of a species. A coating on the surface of a powder is expected to shift the isoelectric
72 point from a value representative of the uncoated surface to a different valuable more similar to that of the coating species.
The surface character of these samples before and after coprecipitation in an aqueous KCl 10-3M suspension was analyzed by measuring the zeta-potential as a function of pH. Initial measurements were performed on the SiC powder directly from the container without any additional treatment. The results in Figure 3.9 are for the raw SiC powders which had an isoelectric point of 3.7. This value is consistent with literature data for SiC powders and is similar to those measured for silica powders [63, 116]. This is indicative of the presence of a thin layer of SiO2 on the surface of the SiC powders.
Figure 3.9: Zeta potential graphs for raw SiC powder and Al/Gd, Al/La, and Al/Y systems at a molar ratio of 4:1 Al to RE. Isoelectric point of raw SiC powders is similar to that of SiO2, as expected. After coprecipitation, isoelectric point values observed to shift to values representative of aluminum and rare earth oxides.
The three remaining curves in Figure 3.9 show the zeta potential versus pH
curves for samples containing SiC with 2 wt. % of various aluminum and rare earth (RE)
sintering aid combinations with a fixed molar ratio of aluminum to the rare earth element
of 4 to 1. The zeta potential results display the expected behavior of metal hydroxide
precipitates coating the SiC powders as shown by other researchers [115, 116]. The
73 presence of a coprecipitated coating significantly changes the isoelectric point of the SiC powder in aqueous suspensions. Whereas pure SiC shows an isoelectric point at a pH of
3.7, the addition of the coprecipitates shifts the isoelectric points into the more basic pH range of 7.4 to 8.6, representing the presence of aluminum and rare earth hydroxides on the surface of the SiC. Literature supports this conclusion as aluminum and rare earth hydroxides alone in solution have isoelectric points in the range pH = 8 - 9 and 9 - 9.4 respectively [115, 116, 120].
Samples were also coprecipitated using a different molar ratio, 3:2 Al to Re, to acquire further confirmation of successful coating. The curves illustrating the zeta potential versus pH, Figure 3.10, show IEP values trending towards even higher pH values than in the 4:1 molar ratio samples. The isoelectric point values are indeed higher as would be expected due to the higher isoelectric points of the rare earth hydroxides as opposed to the aluminum hydroxide. The observed shifts offer proof that a coating has been achieved successfully.
Figure 3.10: Zeta potential curves showing shifts for Al:RE systems at higher RE content of 3:2 Al to RE. Increase in shift compared to Figure 5.1 at 4:1 Al:RE is due to higher isoelectric point of RE hydroxides compared to aluminum hydroxide
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3.6.2.2 FTIR
FTIR results for SiC, Al/La 3:2, and coated SiC Al/La 3:2 before and after sintering confirm the presence of a coating on the SiC powder and the formation of a glassy phase upon firing at 1900°C. The spectrum from the uncoated SiC powder, shown in Figure 3.11, shows both the longitudinal optical (LO) and transverse optical (TO) phonon bands at 929 cm-1 and 861 cm-1 as expected for a lattice constructed by polar
chemical bonds, bands assigned to stretching vibrations of Si-O-C and Si-O-Si,
respectively [121]. The band at 1023 cm-1 can be attributed to the Si-O asymmetric
stretching vibration since it is similar in shape to the band observed in amorphous silica,
indicating the presence of a silicon oxide layer on the SiC powder [121]. The band at 761
cm-1 corresponds to the Si-C stretching vibration mode [122].
Figure 3.11: FTIR Spectra of hydroxide precipitates, uncoated SiC, and SiC - Al/La 3/2 before and after sintering at 1800°C and 1900°C
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The spectrum of only the coprecipitated metal hydroxides confirms the coprecipitation formed hydroxides as expected due to the presence of the bending vibration band of H-O-H at 1670 cm-1. A shoulder observed at 554 cm-1 (not shown) was
-1 -1 assigned to AlO6 octahedra. The band at 803 cm along with that at 1030 cm can be assigned to La-oxycarbonate species according to some authors [123]. On the other hand, the band at 1030 cm-1 can also be assigned to N-based organic residual groups on the
-1 - hydroxides’ surface [124]. The band at 1385 cm is attributed to adsorbed NO3 from the starting nitrates on the surface [124].
FTIR of the system with the hydroxides coprecipitated in the presence of SiC shows the Si-C stretching vibration mode band at 761 cm-1. This band shows no change
from the uncoated SiC as expected. The band at 855 cm-1 appears to be slightly shifted
towards lower wavenumbers as compared to the TO band at 861 cm-1 from the uncoated
SiC, but actually the band may be an overlap between that at 803 cm-1 present in the
coating hydroxides and the unshifted TO band.
Analyzing the samples after annealing at 1800°C for 1 hour, the position of
stretching vibrations of Si-C-O and Si-O-Si bands were shifted towards higher
wavenumbers after annealing (886 cm-1 with a shoulder at 942 cm-1) [125]. The band of
the Si-C stretching vibration is shifted at 780 cm-1 after annealing as well, as reported by
Oliveira et al. [122]. These changes may represent a change in the polytype distribution
of the SiC during annealing. The developing band at 830 cm-1 can be attributed to
condensed AlO4 tetrahedra identified in most aluminates and is indicative of the start of formation of an aluminosilicate glassy phase [126]. The absorption bands due to the presence of La(III) in the glassy phase are usually situated in the spectral range of 400 –
76
500 cm-1 and could not be identified in our spectra because of the frequency range of the
Mattson spectrometer starts at 500 cm-1 [123].
The spectrum recorded after annealing of the sample at 1900oC is different than
that recorded for the sample annealed at 1800oC. An infrared spectrum of silica-rare earth-alumina glassy phase is expected to show bands at ~ 1100 cm-1 and 943 cm-1 respectively. These bands are related to the number of non-binding oxygen atoms and presence of the rare earth compounds [127]. The shoulder in the spectra at 933 cm-1 is actually an overlap between the LO band of SiC and the latter above-mentioned vibration. The shoulder at 670 cm-1 is assigned to the stretching vibration of isolated
tetrahedra found in the silica-rare earth-alumina glass structure [127]. The band due to the
stretching vibration of Si-C becomes sharper and more intense after annealing at 1900oC
as it was previously found by Alekseev et al. The band also shifted towards lower
wavenumbers [121]. This may be an indication of changes in the polytypes of SiC in the
presence of a glassy phase.
As observed in Figure 3.11, in all of the spectra, the strong LO and TO absorption
bands of the SiC masked the other absorption bands, making complete identification through FTIR difficult.
3.6.2.3 Density
As outlined in the Procedures, Section 3.4.3, the earliest samples produced by the standard coprecipitation were washed three times, dried, and ground by mortar and pestle. Samples were then dry pressed and placed in an oven for 24 hours to dry them completely before being loaded into the hot press.
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The density values returned by these samples were consistently lower than expected regardless of the choice or amount of additive. A low density area obtained from a coprecipitated sample is shown below in Figure 3.12. Observation of the coprecipitation process revealed a stratification of the powder into differently colored layers during the washing and settling procedure. It is theorized that the wide particle size distribution, agglomerated nature of the starting SiC powder, and poor dispersion of the slurry resulted in the heavier particles settling fast during the mixing leading to poorer coating quality. The agglomeration of the powder could also result in many particles failing to be coated which could cause defects such as those low density areas observed in the work of Bellosi, Figure 3.6 [115, 116]. Other researchers indicated the importance of the dispersion and slurry properties during coprecipitation [128, 129]. Other work suggested the importance of a prefire or calcination step in removing residual organic phases before sintering [113, 115, 130]. Steps were therefore taken to further improve the coprecipitation process.
Figure 3.12: Micrograph of lower density coprecipitated sample from early powder batch
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3.6.3 Improved Coprecipitation Processing
This set of samples was intended to demonstrate the effectiveness and viability of coprecipitation as a method of producing SiC with high densities, low defect distributions, and improved microstructures. Although coprecipitation was initially effective in developing a successful coating, the overall effectiveness and level of control were lower than desired. It was decided that the results were limited because of shortcomings in the initial procedure. A number of process modifications and improvements were therefore initiated.
3.6.3.1 Milling
A milling step was instituted before coprecipitation in order to reduce the particle size, break up of agglomerates, and narrow the particle size distribution. These steps should serve to better the dispersion of SiC for further processing. It was also hoped that this step could eliminate some of the variability in coating quality that might be responsible for the glassy phase deficient regions observed in the work of Bellosi, as shown in Figure 3.6. These regions are very undesirable as they may be critical defects in structural applications. Literature indicated improved coprecipitation of calcium zirconate through the use of polymeric stabilization and ball milling which reduced the size of agglomerates [131]. Although the authors of the study debated the mechanisms and roles of nonionic polymers in promoting deagglomeration of dispersed particles, the results suggested that the procedure greatly enhanced the reactivity and kinetics of the
79 coprecipitation. Similar work by Yao was equally as supportive of the concept of milling in a surfactant to reduce agglomeration and improve the coprecipitation of ceria based nanoparticles. Polypropylene glycol (PG) was described as an effective dispersant for a ceramic powder, yielding finer precursor grains and little impurity in the final products
[132]. Based on the success of that work and similar studies, polypropylene glycol (PG) was chosen as the solvent in which to disperse the raw silicon carbide for milling [132].
The particle size distribution of the silicon carbide was evaluated for raw SiC powder. These values were compared to previously coprecipitated powder batches and powder that was given the milling treatment in PG. The milling treatment in PG is described in more detail in section 3.4.4.1. Following the milling, the majority of the PG was drained off. The SiC was left in a small amount of the PG to keep it from drying and agglomerating. The particle size measurements were then made after dispersing some of the powder in H2O. The results of the particle size measurements on the milled and non-
milled powders are shown in Figure 3.13. The results indicated a shift in the average
particle size from 452 nm to 269 nm. A considerable reduction in the amount of 1 to 1.5
micron particles is also apparent.
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Figure 3.13: Particle size graph for raw SiC powder versus post milling treatment powder with measurement in H2O. Decrease in average particle size observed, as well as reduction in 1 to 1.5 micron particles.
3.6.3.2 Coprecipitation in Isopropanol
A review of literature and previous experiments with non-coprecipitated systems had indicated that the dispersability of silicon carbide is greatly improved in alcohol versus aqueous systems [50, 64, 133]. A similar particle size analysis to the previous was
performed except the silicon carbide was now dispersed in isopropanol instead of water
before the analysis. A second test was performed with the SiC milled in PG, but then
also dispersing this milled SiC in isopropanol before particle sizing. This was intended to
evaluate if coprecipitation in isopropanol would also augment the dispersion of SiC. The results are shown below in Figure 3.14 and indicate further improvement in the dispersability and reduction of large particles.
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Figure 3.14: Particle size graph for raw SiC powder versus post milling treatment powder with dispersion in isopropanol. Isopropanol when combined with milling treatment reduces particle size and eliminates particles larger than 750 nm.
A review of the literature referenced in this thesis indicated little work on the
viability of coprecipitation in alcohol based systems, but research was initiated to see if
this improved the features of coprecipitated products further.
The composite graph below, Figure 3.15, shows the total improvement in the
deagglomeration and dispersion of SiC gained through the milling and use of isopropanol
as a coprecipitation medium. The particle size of 73 nanometers appears misleading and
the average particle size is not realistically that low. The changes in the shape of the
distribution, reduction of large particles, and general trend of the data is reliable in
representing the benefits of the milling and isopropanol mixing. The data is also listed in
number percent which is appropriate for showing values over such ranges, but may become harder to interpret for very fine powders or non-spherical powders without accompanying data in volume. Because of the better dispersion after milling and mixing in isopropanol, the particle sizing by light scattering is able to register a sufficient amount
82 of well dispersed, fine SiC in this size range. The remainder of the curve does show a longer tail indicating a reasonable amount of SiC over a wider range of submicron sizes.
The results indicate these process changes do enhance the dispersion of SiC.
Figure 3.15: Comparison of all particle size graphs, indicating improved dispersability of SiC with milling treatment and mixing in isopropanol
3.6.3.3 Calcination
A review of literature also pointed out the importance of a calcination step in burning off any remaining organic material after washing. It was also believed that the calcination transformed or decomposed the hydroxides present after coprecipitation into their respective oxides [113, 115, 130]. This reaction is favored versus an in-situ reaction that would occur while heating during the sintering cycle.
Densities were observed to increase after the calcination step at 500 °C was added to the procedure even though the powders were previously washed. Although the washing was expected to remove excess PG and nitrate ions, it was unlikely to remove organics that are bonded to the surfaces. It is theorized that the conversion of the
83 hydroxides into oxides during calcination, as well as the removal of other complexes, enhanced the sintering versus samples that were not calcined. Table 3.2 compares the densification results from some compositions relative to different prefiring conditions.
The percent theoretical densities listed are calculated from the rule of mixtures considering the oxide phases, and as will be shown later, are not absolutely accurate.
Table 3.2: Density improvement with powder calcination and atmosphere control Calcined / Density Theoretical Percent Composition atmosphere (g/cm3) Density (g/cm3) Dense
90 wt.% SiC, 6 % Al2O3, 4 % Y2O3 N 2.73 3.32 82.2
90 wt.% SiC, 6 % Al2O3, 4 % Y2O3 Y / air 3.15 3.32 94.9
90 wt.% SiC, 6 % Al2O3, 4 % Y2O3 Y / Ar 3.20 3.32 96.4
X-ray diffraction analysis was performed on coprecipitated powders from both before and after the calcination step in order to ensure the success of the coprecipitation and confirm the condition of the sintering aids. The scan from the sample before calcination, Figure 3.16, is consistent with an α silicon carbide powder, displaying a majority 6H polytype and a small percentage of 4H polytype as well. No indication of any crystalline phases from the coprecipitation process or sintering aids is observed.
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Figure 3.16: X-ray diffraction pattern and phase matches for coprecipitated powder before calcination treatment
Figure 3.17 below shows the x-ray diffraction scan from the sample after the calcination in argon. This scan is very similar to the previous sample before calcination indicating little noticeable change in the structure of the SiC from the treatment. A small peak from an impurity phase is now visible at 28 degrees. Attempts to identify this minor phase were unsuccessful. Results indicate that coprecipitated coatings are amorphous in nature even after calcining the samples at 500oC. To assess the presence of the coating on the SiC surface, zeta potential analysis was carried out.
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Figure 3.17: X-ray diffraction pattern from coprecipitated sample after calcination in argon at 500°C cycle
3.6.4 Characterization of Coprecipitated Samples After Process Improvements
3.6.4.1 Zeta Potential
Figures 3.18 and 3.19 show the zeta potential versus pH curves for samples coprecipitated in isopropanol at 2 (Fig. 3.18) and 5 (Fig. 3.19) wt. % additives with a fixed molar ratio of aluminum to yttrium of 4 to 1. Pre-calcined and calcined coated samples along with pristine SiC were analyzed. As discussed in the previous zeta potential section, the presence of a coprecipitated Al-Y based coating, either in hydroxide or oxide form, significantly changes the electrokinetic properties of the SiC powder in suspension. While pure SiC shows an isoelectric point at a pH of 3.99, the presence of the coprecipitates shifts the isoelectric points into the more basic pH range of 7.5 to 8.5.
86
This again implies the presence of aluminum and yttrium hydroxides on the surface of the
SiC. The isoelectric point increases slightly after the calcination step but still suggests a viable coating. The increase is likely due to the removal of organics from the powders.
Figure 3.18: Zeta potential curves for 2 wt. % additive coprecipitated samples after milling treatment and with coprecipitation in isopropanol
Figure 3.19: Zeta potential curves for 5 wt. % additive coprecipitated samples after milling treatment and with coprecipitation in isopropanol
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For all coated powders, the curves illustrating the zeta potential values versus pH show IEP values shifted towards higher pH values. The zeta results display the expected behavior of metal hydroxide precipitates coating the SiC powders as shown by other researchers, and appear unaffected by any of the process changes. The observed shifts and values are proof that a coating has been successfully achieved via coprecipitation.
3.6.4.2 TEM of Coated SiC
Transmission electron microscopy was utilized in order to evaluate the surface of the SiC and the homogeneity of the coating. Images were taken on the raw SiC powder, shown below in Figure 3.20. The image shows a grain of silicon carbide, 6H polytype, as determined by the periodicity and spacing of the atomic layers. There is a small rim of amorphous material surrounding the edge of the sample. This is believed to be the thin silica layer that is present on native SiC powders. It varies in thickness from approximately 1 to 6 nanometers.
Figure 3.21 shows TEM images from the silicon carbide powder after coating with yttrium and aluminum sintering aids during the coprecipitation process. A fuzzier, amorphous layer, difference in appearance and thickness than the layer on the raw SiC, was now visible as a coating all over the SiC grains. A relatively uniform thickness of 6
– 10 nanometers was observed for this coating. The appearance of this second coating layer is believed to be representative of the coprecipitated additives and is similar to images of coatings on SiC as shown in the work of Sciti and Bellosi [115, 116].
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Figures 3.20 and 3.21: TEM images of uncoated SiC powder and coated powder from after coprecipitation process
3.6.4.3 BET
BET surface area measurements were also employed as a simple method of determining the presence of a coating. By evaluating the surface area changes at varying points in the processing, it can be established if a coating is indeed present on the SiC.
As shown in Figure 3.22, the surface area increases considerably from 22 m2/g for the
native SiC to 29 m2/g after coprecipitation for the 2 wt. % additive system. For 5 wt. % additives, a further increase to 34 m2/g is observed. Compared to other literature, it
appears our coprecipitation is as successful in coating the SiC [116].
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Figure 3.22: BET results on coprecipitated powders at 2 and 5 wt. % sintering aids
3.6.5 Comparison of Coprecipitated versus Conventional Mixing Processes
3.6.5.1 Hot Pressed Samples
After the initial studies indicated that the improved coprecipitation procedure was capable of coating a SiC powder with sintering aids, a designed set of samples was created for the purpose of directly comparing conventional processing techniques with coprecipitation. Samples of the same composition were prepared through coprecipitation and also processed using conventional processing methods: wet ball milling with powder additives and simple aqueous mixing of the powders using the propeller system that is also used for coprecipitated samples. Compositions using aluminum and yttrium as sintering aids with 5 weight percent additives at a molar ratio of 4 Al:Y were processed.
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3.6.5.1.1 Density
The densities of all three samples prepared for the comparison study were high; around 97% of theoretical based on a rule of mixtures calculation using the expected oxide additives. The densities were quite similar regardless of the method of processing as shown in Table 3.3 below.
Table 3.3: Density values for coprecipitated (C1), ball milled (B1), and propeller mixed (P1) samples in comparison study Theoretical Processing Hot Press Density Percent Sample Composition Density Method Cycle (g/cm3) Dense (g/cm3)
95 wt. % SiC, 1950°C, C1 3.22 % Al2O3, Coprecipitated 30 min, 3.18 3.27 97.2 1.78 % Y2O3 35 MPa
95 wt. % SiC, 1950°C, B1 3.22 % Al2O3, Ball Milled 30 min, 3.18 3.27 97.2 1.78 % Y2O3 35 MPa
95 wt. % SiC, 1950°C, P1 3.22 % Al2O3, Propeller Mixed 30 min, 3.18 3.27 97.2 1.78 % Y2O3 35 MPa
Although the fine, nanoscale additives from coprecipitation were expected to
enhance the densification kinetics during sintering when compared to standard
processing, initial density values indicated no clear advantage. This comparison utilized
the same sintering cycle for all three samples however, and as the hot press system did
not have a dilatometer, there is no ability to consider if the coprecipitated sample
achieved final density at a faster rate and the additional time only contributed to grain
coarsening. Firing the coprecipitated sample for a longer time or at too high of a
temperature also may have promoted the volatilization of the fine sintering aids and this
could have reduced the density slightly. Additional batches of samples in later sections
91 of this thesis will attempt to address these issues. This initial comparison of density also assumes the same grain boundary phase forms in all three materials, yielding the same theoretical density. As will be shown in detail later, this may not be the case when considering materials of different processing method and scale.
3.6.5.1.2 SEM
Examination of a number of regions from each sample revealed greater uniformity in grain size and porosity in the coprecipitated sample when compared to the ball milled or propeller mixed samples. Images 3.23 and 3.24 below show representative secondary electron and in-lens images of the microstructures from the coprecipitated sample, C1. Images 3.25 and 3.26 show the same image types from the propeller mixed sample, P1. The samples were plasma etched together to obtain contrast and preserve the grain boundary phase for examination.
Figure 3.23: Secondary electron SEM image of coprecipitated sample C1
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Figure 3.24: In-lens SEM image of coprecipitated sample C1
Figure 3.25: Secondary electron SEM image of propeller mixed sample P1
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Figure 3.26: In-lens SEM image of propeller mixed sample P1
The coprecipitated sample displayed much finer, uniformly distributed pores and
very few clusters of large porosity. The triples points and glassy pockets are small and
well distributed throughout the microstructure. The size and amount of pullout from
polishing is lessened in these samples likely due to the smaller grain and triple point size.
Grain size measurements were taken from linear intercept analysis using the computer
software package Linc, revealing a grain size for the coprecipitated sample of 0.96
microns. There are a number of very small circular pores present in many of the triple points. It is believed that this may be from trapped gasses due to volatilization or segregation effects because of the non-optimized heating cycle. These trapped pores in the glassy phase are also present in mixed samples fired in the hot press.
While there are some non-porous, uniform regions in the mixed samples, there were also regions of clustered large pores and non-uniform glassy phases such as those
seen in Figure 3.26. The frequency of large pullouts from polishing is greater in the
94 mixed sample. The triple points appear much less uniform in size and grain boundary phase distribution versus the coprecipitated sample. The grain size is somewhat larger in the milled sample at 1.45 microns versus the 0.96 microns of the coprecipitated sample.
There are very interesting differences in the etching behavior, contrast, and grain boundary phase coloring between the samples. The samples were plasma etched together so any differences in appearance are more likely from material differences than sample preparation. The small size and distribution of grain boundary phase in the coprecipitated sample makes the observation of features in the secondary electron image very difficult because of low contrast. The topography of the image is very flat as well. When examining the inlens image where more compositional contrast is obtained, the coprecipitated sample has consistently small and uniform triple points that appear white in color. Some of the larger grains display the distinct “core-rim” or “core-shell” pattern expected from liquid phase sintered silicon carbides that have undergone solution- reprecipitation.
On the other hand, the secondary and in-lens electron images of the milled samples show a much different topography and contrast. The grain boundary phases and triple points appear raised in the secondary image which is expected when plasma etching
SiC. This difference in etching behavior could be indicative of a difference in grain boundary phase composition between the samples. The milled sample often contains larger triple points with bright white inner regions surrounded by a grey rim while some of the smaller triple points are more similar to those seen in the coprecipitated sample.
This could imply poorer mixing or segregation into other phases. As also seen in the coprecipitated samples, there are some trapped circular pores present in many of the
95 larger glass pockets. There is also a more frequent occurrence of core-rim grains in milled sample; this could be due to differences in the liquid phase formation temperature and behavior in these samples as well as the larger grain size. Although the samples are similar in density, there does appear to be subtle and noticeable differences in microstructure between these samples.
The similar density values and percentages for all three processing methods were puzzling initially when compared to the micrographs from these samples. The coprecipitated samples appeared to have a lower average amount of pores than the other samples. The differences in the appearance and coloration of the grain boundaries and triple points raised belief that there might be a compositional or crystallinity difference between coprecipitation and other processing methods. This warranted investigation into the cause of the density discrepancies between the observed and theoretical densities in these samples and many other samples as the work proceeded.
3.6.5.1.3 XRD
X-ray diffraction work was initiated on these samples as an attempt to ascertain the cause of the density differences. Scans on these samples revealed the expected peaks of silicon carbide consisting of 6H, 4H, and 15R polytypes and the presence of some crystalline phases. The higher noise level and quality of the scans taken on the Phillips
Diffractometer units made clear and complete phase identification difficult however. An example of a low intensity scan of coprecipitated sample C1 is shown below in Figure
3.27. Previous researchers have typically found some amount of crystalline content as well as amorphous phases, at least when the additive content was high enough. With only
96
5 wt. % additives total in these samples, a considerable amount of the grain boundary would need to form a crystalline phase in order to be detectable and easily identified.
Figure 3.27: XRD pattern and matches to SiC polytypes for coprecipitated sample C1
Later XRD scans were undertaken using an XRD system with a better signal to noise ratio and improved intensity. Figure 3.28 compares samples C1 and B1. The scans appear reasonably similar and it is now possible to identify an aluminosilicate phase, likely mullite, present in both samples. There are some subtle differences, however. The strong, intense peak at 13° in the coprecipitated sample may be indicative of a silica phase, but a strong match was not found. In the milled sample, the small peak around
20° could not be matched. Samples with greater additive contents will be prepared and higher resolution scans undertaken in order to confirm and further analyze these results.
These results will be discussed in Chapter 4.
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Figure 3.28: Improved resolution XRD pattern and phase matches for ball milled (B1) and coprecipitated (C1) samples in comparison study
3.6.5.1.4 Mechanical Properties
While a more comprehensive study of mechanical properties was undertaken and
discussed in Chapter 5, a brief examination of the elastic modulus and hardness curves
for the early comparison samples is included below. This was intended to uncover and
highlight the need for deeper mechanical property studies and to establish if the
differences in grain size, porosity, microstructure, and phase could lead to any
measurable differences in properties.
Knoop hardness values for the three samples in this study are plotted below in
Figure 3.29 at the test values of 100 g, 300 g, 500 g, 1000 g, and 2000 g. The values are
reasonably similar among all three materials. Considering the strong influence of density
on hardness in silicon carbide however, it is interesting to observe that the coprecipitated
98 sample is slightly harder at most values in the curve. Whether or not this is related to the grain size or phase differences is difficult to ascertain at this time.
Figure 3.29: Knoop hardness curve for three comparison samples
The mechanical property data for samples C1 and B1 obtained from ultrasound
analysis is shown below in Table 3.4. Sample P1 was unable to be evaluated accurately
because of the very small size of the available samples due to breakage during removal
from the hot press. Because the other samples also broke during removal from the hot
press die, there may be some additional error in these calculations. The size and shape of
the pieces made the preparation of flat, parallel surfaces on a piece of sufficient size
difficult. Further comparisons of mechanical properties will be made on other samples in
Chapters 4 and 5.
Table 3.4: Mechanical property data returned from ultrasonic non-destructive evaluation of coprecipitated and ball milled samples Sample Density (g/cm3) υ E (GPa) G (GPa) K (GPa) C1 3.180 0.198 419 176 227 B1 3.184 0.222 392 160 241
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3.6.5.2 Comparison of Coprecipitation and Ball Milled samples through SPS
While the early results above were promising, it was theorized that greater benefits of having fine sintering aids as a coating could be yielded if more rapid heating was employed during sintering. A second batch of comparison samples was prepared using alumina and yttrium sintering aids, but instead of hot pressing, spark plasma sintering was used as the densification method. After the earlier success of samples containing 5 weight percent sintering additives, it was decided to also evaluate samples containing only 2 weight percent additives.
3.6.5.2.1 Density
Densities greater than 95 percent theoretical were found for all 4 samples using the same SPS cycle consisting of a 5 minute hold at 1800°C and a second 5 minute hold at 1900°C with 50 MPa of pressure applied from 1000°C to the conclusion of the cycle.
These results are in Table 3.5. The two-step sintering cycle was chosen because of advantages in grain size that will be discussed later. In both the 2 and 5 wt. % additive cases, the coprecipitated samples showed statistically significant higher densities; 3.20 g/cm3 versus 3.15 g/cm3 for 2 wt. % additives and 3.16 g/cm3 versus 3.12 g/cm3 for 5 wt.
% additives. It is important to note however, that these theoretical density values are based off of the rule of mixture calculations for the percentage of additives and that previous analysis has suggested that the theoretical numbers are inaccurate. The true percentage of porosity is lower than these results indicate. Regardless, it appears that the the rapid heating of SPS is more favorable for densification of coprecipitated samples than the slow kinetics of the hot press.
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These samples are revisited for further analyses as part of a larger matrix of samples from this thesis. SEM images of the coprecipitated samples can be found in
Section 4.3.2.2. Some hardness data for these samples is shown below with remaining analysis in Chapters 4 and 5.
Table 3.5: Density values for spark plasma sintered samples at 2 and 5 wt. % sintering aids, prepared through coprecipitated and ball milling as a comparison Processing Percent Sample Composition SPS Cycle Density Method Theoretical 1800°C, 5 min 98 wt. % SiC, 1.29 % CS2 Coprecipitated 1900°C, 5 min, 3.20 98.97 Al O , 0.71 % Y O 2 3 2 3 50 MPa 1800°C, 5 min 98 wt. % SiC, 1.29 % BS2 Ball Mill 1900°C, 5 min, 3.15 97.51 Al O , 0.71 % Y O 2 3 2 3 50 MPa 1800°C, 5 min 95 wt. % SiC, 3.22 % CS5 Coprecipitated 1900°C, 5 min, 3.16 96.87 Al O , 1.78 % Y O 2 3 2 3 50 MPa 1800°C, 5 min 95 wt. % SiC, 3.22 % BS5 Ball Mill 1900°C, 5 min, 3.12 95.43 Al O , 1.78 % Y O 2 3 2 3 50 MPa
3.6.5.2.2 Mechanical Properties
Figure 3.30 shows the Knoop hardness curves for the four samples prepared in the
comparison study. Not surprisingly, the values are clustered into two groups with the
samples containing only 2 wt. % sintering additives having a higher hardness than the 5 wt. % samples. Otherwise, the coprecipitated samples appear to generally have a slightly
higher hardness values at most of the indentation loads, but the differences are slight and
inconsistent.
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Figure 3.30: Knoop hardness curves comparing SPS samples prepared through coprecipitation (CS) and ball milling (BS) of samples at 2 and 5 wt% sintering aids
Comparison of the elastic moduli values for the differently processed samples reveals some interesting observations. Although the modulus values, Table 3.6, should strongly trend with density, the values for these samples are less sensitive than expected.
The values would also expect to show some trend with the amount of crystalline and amorphous or glassy phases present as these would have an impact on both the density and ultrasonic speed values used in the calculations. The coprecipitated samples which are more dense do show higher moduli values than ball milled samples, but the two are nearly equal, regardless of the CS5 sample having more than twice as much additive as
CS2. Interestingly, the two ball milled samples also show similar moduli values, 381
GPa for 2 wt. % additives, and 377 GPa for 5 wt. %. These results are counterintuitative at first glance, as the density effect is less pronounced as expected, and the amount of additives also appears to have less impact than would be believed. Yet, it is apparent that the different processing routes do have impact on the mechanical properties.
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Table 3.6: Comparison of density and elastic modulus values for coprecipitated and ball milled samples in SPS comparison section Sample Density (g/cm3) Percent Theoretical E (GPa) CS2 3.20 98.97 397 BS2 3.16 97.51 381 CS5 3.19 96.87 401 BS5 3.12 95.43 377
3.6.6 The Use of Fluorescence Analysis to Compare Mixedness
One of the long term issues encountered in this research was how to best indicate
the true improvement in mixing offered by coprecipitation. Logically, the more intimate
mixing and reactivity of the additives should reduce the appearance of certain defects in
the microstructure such as large agglomerates of sintering aids and/or regions deficient in
sintering aids, but the evaluation of such features is typically based on observations of the
microstructure and is not easily quantifiable, provable, or time efficient. Examining the
microstructure also only evaluates mixing of additives on a very large scale and ignores
differences in mixing on a finer level that could result in different phases or properties.
To better analyze the mixing of additives measurements of the fluorescence intensity and decay times of rare earth doped compounds were performed, based on the method outlined by Czerepinski [119]. When rare earth compounds become too close together, as in a poorly mixed region, they have a tendency to undergo a process known as concentration quenching where the two rare earth species will transfer energy resulting in non-radiative energy loss from the system. Properly mixed samples will have better spaced rare earth ions and therefore less energy loss between species, resulting in higher emission intensities and longer lifetimes. If a sample is displaying a shorter decay time and/or a lower intensity, this would be caused by concentration quenching that is indicative of rare earths being close together, additive clustering, and overall poor
103 mixedness. The intensity and lifetimes of a bulk sample can be easily measured in standard spectroscopes and the comparison of different samples can reveal information about the mixedness of samples.
Initial trials indicated that of the rare earths already being used in this study, samarium had the best, although not optimal transmission window for fluorescence in
SiC samples. Using a handheld 365 nm UV light, samarium doped SiC samples were observed to emit a clearly visible orange glow when excited. This emission was also found to be detectable using an Edinburgh Instruments FLS 920 Spectrofluorimeter.
Based on this, a series of samples were prepared in order to determine the mixedness of rare earth doped materials in this study. The samples are described in Table 3.7 below.
Notes on the nomenclature used for identifying samples throughout this thesis can be found in the section 3.4.6.
Table 3.7: Samples prepared for fluorescence analysis Sample Composition Mixing Method
R-54 95 wt% SiC, 2.69% Al2O3 , 2.30% Sm2O3 Coprecipitation
SM5 95 wt% SiC, 2.69% Al2O3 , 2.30% Sm2O3 Ball Milling
R-62 98 wt% SiC, 1.08% Al2O3 , 0.92% Sm2O3 Coprecipitation
SM2 98 wt% SiC, 1.08% Al2O3 , 0.92% Sm2O3 Ball Milling
Comparing the emission intensities and lifetime measurements from the
comparison samples reveal some interesting results. Figure 3.31 below compares the
emission spectra from the two coprecipitated samples at varying additive contents, 2 (R-
62) and 5 (R-54) weight percent samples. The reason for comparing two samples
prepared in the same manner, but with varied additive levels is to indicate the validity of
the method. As can be seen in Figure 3.31, the emission from the sample with 2 wt. %
additives is about four orders of magnitude higher than the 5 wt. % additive sample. This
104 would be expected, even among two well-mixed samples as a higher dopant concentration should result in some rare earth ions being closer to themselves and this should contribute to increased concentration quenching and lower intensities. Lifetime measurements between the two samples, Figure 3.32, show a much longer lifetime for the lower dopant sample which matches the emission results and indicates less quenching.
Figure 3.31: Emission spectra from coprecipitated samples at 5 (R54) and 2 (R62) weight percent sintering additives
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Figure 3.32: Fluorescence decay curves from coprecipitated samples at 5 (R54) and 2 (R62) weight percent sintering additives
A comparison of the coprecipitated and ball milled samples prepared with 2 wt. %
sintering aids indicates some possible differences between the two processing techniques.
Figure 3.33 is the emission spectra showing a two orders of magnitude greater emission for the coprecipitated sample versus the ball milled sample. This result suggests that there is a mixedness advantage in using more advanced colloidal processing techniques like coprecipitation in order to ensure more intimate mixing of the sintering aids.
Perhaps even more meaningful, the shapes of the two spectra are somewhat different.
Differences in peak heights, shifts, and splits are believed to be indicative of local structure modifications.
Figure 3.34 below is comparing the spectra from two locations in the sample and also shows noticeable differences. This gradient effect is common in liquid phase sintered materials as the edge of the sample often reacts with the local environment, i.e.
the die, and this can change the viscosity, composition, and properties of the liquid in this
106 region of the sample. These results indicate that processing can also be highly influential on the local structure and crystallinity of materials.
Figure 3.33: Emission spectra comparing coprecipitated and ball milled samples at 2 weight percent additives total
Figure 3.34: Fluorescence decay spectra showing possible structural gradients in ball milled samples at 2 weight percent additives total
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The fluorescence lifetime measurement comparing coprecipitation to ball milling,
Figure 3.35, reveals further evidence that coprecipitation develops samples with better mixedness. The lifetime of the coprecipitated sample is almost 2 times higher than the ball milled sample. This could be indicative of a more uniform grain boundary composition throughout the sample as well as less regions deficient in sintering aids.
This is sound initial evidence that coprecipitation could have important ramifications and benefits in the development of SiC.
Figure 3.35: Fluorescence decay spectra showing improved fluorescence decay in coprecipitated versus ball milled samples at 2 weight percent additives total
Optical and scanning electron microscopy examination of coprecipitated
microstructures indicates these conclusions based on fluorescence are reasonable, as
these samples show little porosity or large regions deficient in sintering aids, such as
those seen in samples prepared Sciti and Balbo [115, 116].
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3.6.7 TEM/EDS of Grain Boundaries
As another confirmation of the fluorescence results the chemical composition analysis was performed using EDS. The goal was to compare the grain boundaries in samples prepared from coprecipitation versus those prepared from ball milling. Because of the fine, nanoscale grain boundaries any careful compositional analysis required high resolution to show discrete or well mixed components. Attempts to perform this analysis in the SEM were shown to be unsuccessful because of a lack of resolution so EDS analysis in TEM was pursued instead. This was originally intended to be a substantial study, but difficulties with sample preparation necessitated a reduction in the number of samples. The later discussion of future work will strongly endorse continuing this methodology, but even the brief analysis here gives further credence to the improved mixing of coprecipitation.
For this TEM work the samples were chosen from the 5 wt. % additive samples previously investigated in the comparison study in section 3.6.5.2. As a review, samples were prepared with the same composition, 95 wt.% SiC, 3.22 wt.% Al2O3, 1.78 wt.%
Y2O3, but one was prepared through the improved coprecipitation process while the other
was prepared from simple ball milling of the SiC and oxides. The samples were
densified using the same spark plasma sintering cycle.
Figure 3.36 is an EDS map on the ball milled sample. Because the majority of the
sample is SiC the signals from Si and C strongly dominate the area of the sample, and
better color maps are obtained when ignoring them. The map shows the presence of
aluminum, yttrium, and oxygen in the grain boundary pockets. Each is represented by a
specific color: yellow for aluminum, purple for yttrium, and blue for oxygen. The maps
109 indicate there are regions of discrete alumina present in the ball milled sample. There is no detectable yttrium content in these areas as the maps shown. This is indicative of poorer mixing and the possibility of large phase variability throughout the sample.
Figure 3.36: Elemental analysis and map of aluminum, yttrium, and oxygen at grain boundaries in sample prepared from ball milling. Note aluminum rich and yttrium deficient area in upper left of images.
The results from the coprecipitated sample, Figure 3.37, indicate a different
situation. With the improved mixing and reduced scale of the sintering aid additions
introduced by coprecipitation, discrete regions of alumina or yttrium were not observed
110 in this sample. The colors mix uniformly throughout the image. This can be interpreted as improved mixing and an overall different microstructure relative to phase. This change could lead to modified mechanical performances and vastly different materials.
Figure 3.37: Elemental analysis and map of aluminum, yttrium, and oxygen at grain boundaries in sample prepared from coprecipitation showing improved mixing and distribution of elements
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3.7 Summary
An innovative and improved coprecipitation procedure was developed for preparing liquid phase sintered silicon carbide materials with reduced large defects due to poor mixing and scale of sintering aids. By precipitating the sintering aids out of solution as opposed to adding them as discretely sized powder, the ability to process fine, well mixed silicon carbide was gained. Further control of the dispersion of the SiC as well as the electrical stabilization of the system was utilized to ensure the best quality coating of sintering aids around SiC versus previous researchers. Comparison samples to standard ball and aqueous milled samples that were hot pressed revealed some differences in density and slight differences in mechanical properties, but not as significant as expected.
More significant differences were observed when the samples were densified via spark plasma sintering. Clear and strong differences in the level of mixing and phase formation were implied from fluorescence measuments and later confirmed by nanoscale TEM of the grain boundary phases. The next chapter, Chapter 4, focused on refining the remaining process and procedures to better control the microstructures and interrogate the phase and structure differences.
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Chapter 4:
Coprecipitation as a Method of Microstructure Control
4.1 Introduction
The advanced coprecipitation procedure developed in Chapter 3 was shown to be a simple and versatile processing method that was successful in improving the mixedness and scale of sintering aids. This also suggested possibilities for reducing the size and frequency of certain microstructural defects common to SiC, namely a reduction in large pores and sintering aid inclusions. These results were supported through innovative fluorescence measurements as well as direct microstructure comparisons of coprecipitated versus ball milled samples. While it was clear from comparing the microstructure of coprecipitated and ball milled samples that they were different relative to grain and pore sizes, it was surmised that more significant microstructural differences may have been gained through the modified processing route. This chapter therefore focuses on the understanding of the differences in microstructure between coprecipitated and ball milled samples, as well as the establishment of what processing variables are relevant in controlling the microstructural development of advanced SiC ceramics.
4.2 Microstructure Development in SiC
The green processing of polycrystalline ceramics presents many opportunities for innovation and control. The high percentage of covalent bonding necessitates sintering aids in order to achieve high densities and the ability to solid state or liquid phase sinter
SiC gives much flexibility. By directing the processing and sintering, features such as
113 grain boundaries can be controlled and emphasized in order to develop new fracture behaviors and mechanical properties. On the other hand, if processing is neglected in the development of grain boundaries and microstructure, defects can be created which degrade performance and make the benefits of grain boundary changes negligible.
4.2.1 Grain Boundary Engineering
By changing the features of the grain boundary phases in SiC, very different mechanical responses can be observed. The reaction of the native SiO2 layer with any
additives introduced into the system can develop liquid phases and resultant glasses with
widely varying properties such as refractive index, dielectric contrast, and mechanical
properties like elastic modulus and hardness [12]. Zhou et al. developed LPS-SiC with varying mechanical properties by adding a number of rare earth elements. One of the key factors in differentiating the effect of one rare earth from another is in the cationic radius effect. The properties of a typical glass are defined by the bond strength of the SiO2 bonds; the addition of cationic species as network modifiers changes the repulsion of the neighboring atoms and the corresponding bond strengths. This change is modeled similarly to the effects of the gravitational field of two objects on each other and this effect can likewise be defined by a modified gravitation equation:
2 2 Fc-a = ZaZce /r . (Equation 11)
where Fc-a is the field strength of the bond, Za and Zc are the charge of the anion and
cation respectively, e is the elementary electron charge, and r is the distance between the
cation and ion, which is equivalent to the addition of the ionic radii. The end result is that
the bond strengths and resultant thickness of the glass is sensitive to the cationic radius of
114 the additives. These additions can therefore be used to strongly influence the properties of the intergranular phases [12, 14, 16].
Liquid phase sintering with the correct combination of additives also allows for the development of grain boundary zones with gradient compositions that may offer beneficial behaviors. The formation of a liquid phase can encourage the dissolution of impurity phases or SiC grains into the liquid during the sintering process. Upon cooling, the solubility of the liquid decreases with temperature and the impurities and SiC can be slowly reprecipitated out of the liquid as a fine gradient region at the grain boundary.
This process is called solution-reprecipitation and leads to the development of an easily identified microstructure consisting of a SiC “core” region and a boundary zone known as the “rim” or “shell”, consisting of the expelled phases [134-136]. An example of this phenomenon as observed in a sample from this thesis is shown below in Figure 4.1.
Figure 4.1: Example of core/rim microstructure observed in LPS-SiC. Solid, red lines and circles indicate core regions. Blue, dotted lines outline rim regions.[137]
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The development of a core-rim microstructure is a highly versatile method of processing SiC that has the potential to influence many desirable properties. Hu et al. hot pressed SiC with varying amounts of AlN and small additions of Y2O3. They observed
clear evidence of diffusion of some AlN into the core of SiC grains during the solution formation process while a greater amount of AlN was present in the reprecipitated shells surrounding the grains. With higher amounts of AlN additions into the sample, AlN based grains precipitated out in competition with the AlN-rich shell regions. Amorphous grain boundary phases and crystallized triple points were also observed with varying compositions. Careful compositional controls over the amount of Y2O3 added allowed
the researchers to develop microstructures with little intergranular glass phase and small
triple points, resulting in microstructures with good room temperature properties such as strength, 1130 MPa, and fracture toughness, 6.2 MPa m1/2. The careful choice of
composition and design of the microstructure resulted in a high temperature strength only
one-third lower than the room temperature, but this loss was accompanied by a one-third
gain in fracture toughness at high temperature [134]. A number of other researchers have
done studies on core-rim microstructures and the range of properties and modifications
possible through additive controls and processing conditions [13, 14, 55, 96, 135, 138].
The importance of engineering the grain boundaries in SiC is also seen in the
work of Gu et al. A reduction in grain size and modification of the grain boundary
composition were shown to promote grain boundary sliding and other plasticity
mechanisms in SiC tested at high temperatures in tension and compression. Stress-strain
curves indicated deformations of 44 % in tension and 56 % in compression when testing
at 1700°C to 1800°C in argon. The materials failed through interdiffusion effects leading
116 to the segregation of crystalline phases at the grain boundary. These phases prohibited further grain boundary sliding and lead to eventual fracture. Although these tests represent behaviors promoted at high temperatures, information gathered on microstructure changes and phase development could be useful for low temperature applications. Silicon carbide with engineered grain boundaries could potentially show improved plasticity through limited grain boundary sliding or microcracking behaviors and lower temperatures [74].
4.2.2 Grain Size
Hall and Petch indicated that reduced grain size may indeed lead to higher strengths, but those assumptions do not consider the uniformity, interfaces, grain boundary characters, and polytypes of the grains involved [139, 140]. Although many researchers have postulated that nanograined polycrystalline ceramics would display exceptional combinations of hardness, ductility, and yield strength, the proof and physical realization of these effects has been more elusive. Some researchers have in fact observed an inverse Hall-Petch relationship where some nanograined polycrystalline materials soften with a reduction of the grain size below a certain level. It is still an open question as to which effect is more likely [139, 140]. If a Hall-Petch relationship does exist, accessing such behavior may not be as simple as reducing grain size. Krell has discussed the challenge of producing dense nanograined ceramics without introducing defects which would overshadow the improvement in properties [141].
Some researchers have been successful in controlling grain sizes to develop microstructures with both good hardness and improved ductility. Zhan was able to
117 produce ultrafine SiC microstructures using 90 nm β-SiC starting powders. Through the seeding of some samples with α-SiC, Zhan could produce microstructures with fine, equiaxed grain structures or more anisotropic grains in a finer matrix, or combinations.
Although the finer grained equiaxed structures were able to display superplastic deformation capabilities, it was found that the mechanical properties could actually be improved with careful annealing and growth to final grain sizes of 1 micron [103]. Jiang and Wang were able to model the plasticity effects of different grain sizes and boundary characteristics, as well as study the model in comparison to nano-sized TiO2. They found
that there is indeed an optimal grain size range for compressive strength which is very sensitive to porosity, grain size, and secondary phase characteristics [142]. Other researchers working with nano-sized SiC have found an optimal range in which strength and fracture toughness or other effects can be optimized [143, 144].
Processing brittle materials such as silicon carbide with the correct distribution of grain size and grain boundary width might allow for the discovery of a sweet spot where the material can have both high strength and hardness as well as improved ductility. By modifying the processing compositions to lower temperatures, it is hoped that the grain sizes can be reduced sufficiently to influence the behaviors. Innovative processing methods such as spark plasma sintering (SPS) or microwave sintering may also be explored as methods of reducing grain size and improving microstructure. Being able to understand and control the features of the microstructure though processing is an essential pathway to better performing structural SiC.
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4.3 Results
4.3.1 Microstructure Control
Microstructure development in ceramic systems is a complicated process that is strongly impacted by every processing step, change, and variable. It can be difficult to independently isolate a single processing change and observe its influence on the development of the microstructure. In a more general sense, changing the amount of sintering aid and studying the change in the microstructure may not be a simple study; the sintering cycle may require longer heating or higher temperatures to achieve the same density for fair comparison. Likewise, modifying a variable like grain boundary thickness can be achieved through both simple increases in the amount of sintering aids, but also through compositional changes, i.e. modifying the rare earth species. The results are not expected to be the same in resultant properties, and as such, both methods are explored separately in the sections to follow. Keeping some changes and comparisons independent from other variables can be difficult and result in even more daunting and larger sample sets. Therefore, many of these sections will contain smaller studies of variables chosen for a specific sample set, but with some overlap and redundancy.
4.3.2 Microstructure
4.3.2.1 Influence of Sintering Aid Mixing – Coprecipitation versus Ball Milling
As this thesis is aimed at preparing high strength structural SiC, the majority of samples processed were of low additive content. Attempts to quantify the presence of any crystalline or amorphous phases at the grain boundaries with samples consisting of 2 or 5 wt. % additives proved to be difficult however, both when using bulk XRD
119 techniques to detect the presence of any crystalline phases and through EDS in the
FESEM for mapping or phase identification purposes. The small quantity of grain boundary phase, the thinness of the boundaries, and the small size of the triple points in these samples had made these analyses difficult. A set of samples containing a larger amount of sintering additives was therefore processed in order to be used as a diagnostic for determining some of the phase properties. At 10 wt. % additives, it was not expected that this sample would show desirable mechanical properties even with the improved processing of coprecipitation, but would instead allow for easier characterization and comparison. A second 10 wt. % sample was prepared through standard ball milling techniques in order to enable a comparison of the grain boundary phases. An examination of the densification behavior of the 10 wt. % samples was also carried out.
4.3.2.1.1 Densification Results
A number of powder batches with 10 wt. % sintering aids, compositions of 90 wt% SiC, 6.4 wt% Al2O3, and 3.6 wt% Y2O3, were prepared from coprecipitation.
Samples of this composition were fired over a wide temperature range to assist in
determining the densification properties. Micrographs of samples from this batch (R61)
are shown below in Figures 4.2 and 4.3.
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Figure 4.2: SEM micrograph of coprecipitated sample R61, 90 wt. % SiC, 6.4 wt. % Al2O3, 3.6 wt. % Y2O3. The density appears to be much higher than theoretical density calculations would indicate
Figure 4.3: SEM micrograph of coprecipitated sample R61, 90 wt. % SiC, 6.4 wt. % Al2O3, 3.6 wt. % Y2O3.
In some of the earlier work in Chapter 3 it was observed that the densities calculated through Archimedes density measurements made the samples appear less dense than observation of microstructures through scanning electron microscopy indicated. The samples with higher additive content again appear to be much more dense
121 than the theoretical calculations would imply. With the higher amount of additives present in these samples, the theoretical density as calculated would be 3.32 g/cm3 while densities returned from initial samples were in the 3.20 g/cm3 range. Observations of the micrographs reveal little porosity however.
It is possible to infer differences in phase formation and reactivity through better
analyzing the densification process through dilatometry, as this can reveal if sintering is
proceeding through a lower temperature or different route or if outgassing or phase
changes are occurring. Unfortunately, a lack of working dilatometers on the sintering
systems made true measurements impossible, so a densification curve was instead
prepared from the individual density values of samples sintered under different cycles.
As Figure 4.4 shows, over the temperature range of 1825°C to 1870°C the density
is relatively unchanged and seems to plateau around 3.20 g/cm3. This would imply a density percentage of only 96%, while the images reveal less porosity than that. The samples seem to be 98% dense and above. The densification curve and microscopy imply that coprecipitation is an effective densification method, even at relatively low
sintering temperatures and times, but the assumption of theoretical densities is incorrect.
Weight loss from additive volatilization and the coprecipitation process modifying the amount of silica present were investigated as likely causes for some of the differences, but neither seemed to fully justify or explain the discrepancy.
It was therefore theorized that the different processing methodologies and sintering methods employed in this thesis have a strong influence on the phase formation
at the grain boundaries, both in content and crystallinity. The fact that the elastic
modulus values in Figure 4.4 drop off for the lowest sintering temperature even though
122 the densities are very similar may indicate a shift in phase crystallinity during sintering.
Differences in grain boundary phase and crystallinity are often easily implied from x-ray diffraction results showing the presence or lack of certain crystalline phases. For these reasons, XRD analysis of the higher additive systems was employed as the primary method for determining the extent of phase differences introduced through processing.
Figure 4.4: Variations in density and elastic modulus values with different SPS sintering cycles for 10 wt. % additive samples
4.3.2.1.2 XRD Analysis
X-ray diffraction scans were performed on a sample prepared with 10 wt. %
additives. The initial scans were performed on a Phillips Diffractometer and revealed a
few small peaks that may be indicative of some detectable crystalline phases, but the high
noise level and quality of the pattern prevented the identification of any of the obvious
possible or expected phases such as yttrium aluminum garnet (YAG) phase, excess
alumina or yttria, amorphous phases, or any crystalline silicate phases. The size of the
peaks relative to the background made their determination versus noise peaks very
123 difficult. The quality of the sample surface and possible contamination of carbon also influenced the desire to redo this composition and obtain better scans.
In order to better observe and quantify the crystallinity differences, a newer x-ray diffraction system (Phillips X’pert) with a much better signal to background ratio was utilized. Results from the improved scans are shown below in Figure 4.5. The improved noise level revealed some small peaks, but identification of the phases of interest was still difficult.
Figure 4.5: XRD pattern of 10 wt. % additive sample. Only 6H, 4H, and 15R polytypes of SiC identifiable
An even higher resolution XRD (Panalytical X’Pert MPD Pro) was therefore used for some samples of interest. Scans from sample R61 are shown below in Figures 4.6 and 4.7. The very high count rate and low background intensity from this machine enabled a much better characterization of the 10 wt. % composition. As seen in all other
124 samples, the expected diffraction peaks indicative of the 6H, 4H, and 15R polytypes were observed. Some card files indicated the presence of other less common SiC polytypes, but differentiating them was difficult. The other peaks apparent in the sample indicate the formation of a crystalline 2:1 mullite phase, 2 Al2O3-SiO2. The stoichiometry
indicates the composition is oxygen deficient, which from literature is common for 2:1
mullite.
Figure 4.6: High resolution XRD scan on coprecipitated sample R61
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Figure 4.7: XRD pattern of R61, zoomed in on area of interest. 2:1 mullite phase now identifiable in improved scan
Estimating the amounts of the individual phases present in this sample and many others in this thesis from either RIR calculations using the Easy Quantitative function in
Jade (Data on Jade programs) proved to be challenging because of the considerable overlap of the most intense peaks, texturing from the bulk sample, and from the possible presence of a considerable amount of glassy phase. By selecting peaks carefully and refining the pattern fit, quantitative results were obtained; these are shown below in Table
4.1. The results are expected to be less accurate for determination of exact polytype amounts, but are more revealing for the amounts of minority phases. It is interested to note that there is no visible crystalline phase containing yttrium which could indicate the improved mixing resulted in the formation of a yttrium containing glass.
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Table 4.1: Phase amounts in coprecipitated sample R61 as estimated by Easy Quantitative function from peak fit in Jade program Phase RIR Wt.% SiC (6H) 1.40 87.8 (28.2) SiC (4H) 1.26 6.4 (2.1) SiC (15R) 0.85 1.5 (0.5)
Al4.64Si1.36O9.68 0.77 4.2 (1.4)
Figure 4.8 below is an XRD scan taken under the exact same conditions for a 90
wt. % SiC, 10 wt. % additive sample prepared through ball milling and spark plasma
sintered at the same conditions as the coprecipitated sample. It is clear that the different
processing route results in a sample with a very different distribution of crystalline
minority phases present in the XRD scan. For the coprecipitated sample the only visible
crystalline phase was mullite, but in the ball milled sample a number of smaller peaks are
observed. Identification of all of the peaks was difficult, but yttrium silicate, Y2Si2O7, and Al2O3 are both observed in the pattern. The selected XRD regions shown below in
Figures 4.9, 4.10, and 4.11 highlight the important peaks for the phases of interest.
Figures 4.10 and 4.11 emphasize the appearance of the alumina in the ball milled sample. The presence of the alumina and a separate yttrium containing phase may be indicative of the poor mixing and reduced reactivity expected of a ball milled system.
Since the initial composition is alumina rich and the powder is coarse, excess alumina is expected and observed in previous research in conventionally prepared SiC systems [88].
Figure 4.12 shows the two graphs on the same image in to allow for easy comparison.
Other than the differences in phases mentioned above, there are no other obvious major
changes in polytype or texturing within the expected error calculations in the
quantification. Although the amount of crystalline phases is small it is apparent that the
127 composition along the grain boundaries in each material is subtly different due to the reactivity and mixedness differences introduced through the varied processing methods.
Figure 4.8: High resolution XRD pattern from ball milled 10 wt. % additive sample
Figure 4.9: XRD region showing appearance of Y2Si2O7 phase in ball milled sample
128
Al2O3
Figure 4.10: XRD pattern of 10 wt. % additive sample prepared through ball milling. Highlighting the peaks attributed to the excess alumina
Figure 4.11: XRD patterns from 10 wt. % sintering aids for both coprecipitated and ball milled samples. Note excess Al2O3 in ball milled sample
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Figure 4.12: XRD pattern comparison of 10 wt% additive samples from coprecipitation and ball milling
Table 4.2 is the quantitative analysis results from the ball milled sample showing the presence of appreciable amounts of the phases indicated. The amount of the alumina suggested is unrealistic since the total amount for this composition is 6.4 wt. %. This is likely due to the strong overlap of some of the alumina peaks. The fit gave higher error for this calculation itself. Even accounting for this, the final composition is detectably different than samples prepared through coprecipitation.
Table 4.2: Phase amounts in 10 wt. % additives ball milled sample as estimated by Easy Quantitative function from peak fit in Jade program Phase RIR Wt.% SiC (6H) 1.40 75.9 (39.6) SiC (4H) 1.26 12.6 (6.6) SiC (15R) 0.85 3.0 (1.6)
Al2O3 1.04 6.4 (3.4)
Y2Si2O7 1.34 2.0 (1.1)
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These results offer further confirmation of the improved mixing offered by coprecipitation and are complimentary to the observations of the TEM/EDS study of
Chapter 3. The ball milled sample examined in the TEM showed discrete areas of alumina present in the microstructure while the XRD indicated the presence of alumina as a discrete phase. The coprecipitated sample color maps displayed well mixed colors and no discrete alumina; the XRD also did not reveal any alumina present for coprecipitated samples. This is a significant result demonstrating the mixing improvements and possible changes in phase that can occur with coprecipitation.
Comparing these results to other researchers who processed silicon carbide with similar additives revealed these results as intriguing. Zhou et al. was able to find YAG present in 5 vol. % doped samples prepared by mixed oxide powders [14]. Mrotek observed the formation of YAG in silicon carbide samples processed with starting oxide powders as well as sol gel additives [20]. Balbo et al. coprecipitated SiC with 10 wt. % additives and were able to detect 5 wt. % YAG in some samples with the presence of trace amounts of crystalline alumina as well. They explained the presence of the alumina as being a result of the excess alumina in the starting compositions versus the stoichiometry of YAG. Their samples were sintered through a pressureless sintering method with a powder bed which may account for some of their phase differences [115,
116]. Interesting work performed by Can et al. compared theoretical thermodynamic calculations to gas pressure sintered and hot pressed samples with 10 wt. % additives.
The compositions similar to the one in this thesis indicated the likely formation of YAG, mullite, and Y2Si2O7, but showed different products depending upon the sintering method
and atmosphere: Y2Si2O7 and Al2O3 for hot pressing, while gas pressure sintering did
131
reveal YAG and Al2O3. Although the equilibrium calculations from thermodynamics
indicated mullite formation as a likely product, neither method of experimental sample
preparation in their work showed that mullite formed. The key difference between which
phases were yielded is relative to the loss or retention of SiO2 during sintering; the added
confinement of the die and pressure utilized during hot pressing retained the SiO2 in the sample leading to the formation of silicate based boundary phases. In the gas pressure sintered sample where the SiO2 is lost the final phases are aluminates [88].
The initial scans for the coprecipitated samples prepared in this thesis showed
neither measurable amounts of YAG nor excess Al2O3, which as discussed above, most
other researchers observed for similar compositions. The calculated densities from this
thesis were also lower than any of the above researchers’ indications. It was surmised
that the improved coprecipitation and/or rapid heating of the SPS system was leading to
different phase formation and crystallinity versus other methods depending upon the
temperatures and cycles used for densification.
4.3.2.1.3 EDS Mapping in SEM
The high amount of sintering aids in sample R61 enabled EDS maps to be created
for selected areas of the microstructure. In samples with less additives, the smaller size
of the triple points and residual phase areas were problematic for mapping because of the
large interaction volume of x-rays. A region of sufficient grain size was found in sample
R61 and generalized compositional information was generated using the Inca x-ray
system on the Sigma FESEM at 30 keV. The area of the map is highlighted in purple in
Figure 4.13. Figure 4.14 is the compilation of the individual element maps obtained
132 during the scan while Figure 4.15 is a reconstructed color map using the mixing of colors to indicate where certain combinations of the elements are. The scans indicate the clear presence of oxygen and aluminum along the grain boundaries and larger pools. Although it appears some regions are aluminum and oxygen rich, implying the presence of discrete alumina areas, the XRD data did not detect any excess alumina in the coprecipitated sample. While it does appear these regions are lacking in silicon, the fact that silicon is present in the bulk phase causes the intensity throughout the scan to be heavy in silicon from overlapping so the regions that appear dark are not necessarily absent of silicon, but just appreciably lower. It is possible that these regions are the crystalline mullite phase which would also be alumina rich. Although the yttrium was not detected in any specific phase in XRD, these results do confirm the presence of it in the microstructure, and it does appear to be generally located near the grain boundaries. These results, especially the strong oxygen and aluminum signals framing around the grain boundaries, loosely support the presence of a well distributed glassy phase.
Figure 4.13: Area of interest for EDS mapping in FESEM
133
Figure 4.14: Individual elemental maps from EDS scan of sample R61
134
Figure 4.15: Mixed color map from EDS data showing association of specific elements
4.3.2.1.4 Presence and Effect of Glass at Grain Boundaries
The results from the XRD study of the samples with 10 wt. % additives indicated the presence of a limited amount of crystalline phases in the microstructure. The coprecipitated sample specifically did not show the presence of any yttrium containing phase although the EDS results did confirm the coprecipitation was successful in adding it. Although the open structure of the mullite is capable of incorporating reasonable amounts of some cations such as the smaller and midsized transitions metals, the larger size of yttrium and other rare earth cations make their incorporation in large amounts unlikely [145]. Micrographs of the samples shown above and all other samples examined in this thesis reveal the presence of fair amounts of material at the grain boundaries and in the triple points. The conclusion from this, as is observed by many other researchers, is that a considerable amount of the liquid phase formed during sintering remains behind in the microstructure as a glassy phase [55, 88, 115, 116].
The XRD results however, failed to show any obvious amorphous hump as expected with the presence of an amorphous phase. A series of XRD scans were then
135 performed in order to establish the detection limit for the amount of glassy phase that can be detected in silicon carbide. Figure 4.16 shows the XRD scans from a P281 glass frit.
Although the composition of the glass frit is likely different than the unknown glass composition in the sintered samples, it was decided that it would allow for an approximation of the detectability of glass in SiC. Figure 4.16 is a scan of 100 percent of the powdered glass and reveals a clear amorphous hump at 23 degrees and some limited crystalline peaks, likely silica in this case. A matching scan of the silicon carbide powder only was also performed. The same glass frit was then carefully mixed with the silicon carbide powder at 5 and 10 weight percentages through the use of a mortar and pestle.
Scans using the same conditions were then run on these samples and are shown in Figure
4.17. Examining the scans reveals that the glassy phase is difficult to detect in SiC at the percentages of 5 and 10 wt. %. While there is no clear amorphous hump shown, there are differences in the slope and intensity of the three curves in the region where the amorphous hump would be. This indicates the presence of the glass although not in an easily detectable amount. The work of Kuchinski indicates a possible method of calculation for the amount of amorphous phase present, although it was not attempted here. It requires more careful spiking studies and a known glass composition to achieve accurate estimations [146]. Since the maximum additive percent used in these samples is 10 weight percent, and not all of the additives would be expected to form a glass, this study demonstrates why no amorphous hump was identified in any of the XRD scans throughout this thesis. It is therefore still reasonable to assume the glass content of these samples can be reasonably high, as opposed to assuming an unidentified phase.
136
Figure 4.16: XRD pattern obtained for 100 percent glass frit
Figure 4.17: XRD patterns from raw SiC powder, overlayed with scans of the powder mixed with 5 and 10 wt. % of the glass frit
137
From the XRD and EDS analyses it can be concluded that the coprecipitation and ball milling processes develop SiC materials with glassy grain boundaries, but with clear differences in residual and secondary phases crystallinity. It could be concluded from these observations that the glass properties and structure could also be fundamentally different. Coprecipitation may not just make a better mixed material versus ball mill processing, but a material that can be characterized as different altogether.
4.3.2.2 Influence of Sintering Aid Percentage
While samples prepared with high additive contents make characterization easier, they may not be of the greatest interest to structural ceramic systems because of the loss of hardness and other property changes imparted from high glass contents. Samples with lower additive contents but the benefits of coprecipitation may provide better results and are of more interest here.
4.3.2.2.1 Densification
For the study of varying sintering aid amounts, aluminum - yttrium based samples were targeted with 2 and 5 wt. % of sintering aids. A number of powder batches at each composition were prepared from coprecipitation.
Compositions at both 2 and 5 wt. % in the yttrium systems showed good sinterability over a range of sintering cycles indicating both the success of coprecipitation in making a highly sinterable powder and the importance of sintering method and cycle design in microstructure control. These variables will be discussed at length in later
138 sections. A brief discussion of the effect of sintering cycle is warranted here however in order to explain why the sintering cycle as shown was chosen.
Because these compositions are liquid forming systems, it was reasonable to believe that many of the effects and benefits of liquid phase sintering would apply for both coprecipitation and spark plasma sintering even though literature in the sintering of such a system is sparse. Mrotek, Cordrey, and other researchers have indicated that the temperature of liquid formation and the amount of liquid both have strong influence on the grain size. The presence of greater amounts of liquid at lower temperature promotes more nucleation events during solution-reprecipitation because of the longer diffusion path in the greater liquid volume; this discourages growth after nucleation because of crowding [10, 20]. Since this thesis work was targeting fine grains as an important means of improving properties like grain boundary plasticity, the cycles were designed with this effect of liquid in mind.
Initial samples supported this theory and the importance of the cycle relative to grain size. Dense samples at 5 wt. % sintering aids were obtained with a simple cycle consisting of a fast ramp up to 1900°C with a 10 minute hold. A second sample with the same ramp but with an intermediate hold at 1800°C for 5 minutes, and a subsequent 10 minute hold at 1900°C resulted in nearly the same density. Microstructures of the two samples are shown in Figure 4.18 The grain sizes are different however as the sample sintered directly to 1900°C resulted in a grain size of 0.86 ± 0.40 μm while the grain size from two step sintering cycle sample was 0.61±0.10 μm. Observation of the microstructures and grain size measurements reveal that the two step cycle resulted in a microstructure with a slightly lower grain size and less deviation between sizes. It is
139 theorized that the lower temperature hold allowed for the development and distribution of the liquid phase at this stage. When the temperature was then ramped to the final hold where grain growth would be most encouraged, the presence of the liquid and its improved distribution helped constrain the grain growth during further solution- reprecipitation. While it would be expected that this effect would be more pronounced with the longer, slower heating in a hot press, it is apparent from these results that it also holds for the rapid heating in the SPS. For this reason, all remaining SPS samples in this thesis were prepared from cycles with a lower temperature hold. This does not mean that the same cycle was sufficient for densifying all samples however.
Figure 4.18: Grain size comparison for SPS cycle with no intermediate low temperature hold (1900°C for 10 min.) versus cycle with lower temperature intermediate hold (1800°C for 5 minutes, 1900°C for 10 minutes)
4.3.2.2.1.1 2 weight percent
Samples containing 2 wt. % alumina-yttria sintering aids were sintered to high density in the SPS with cycles consisting of ramping at 200° per minute to 1800°C with a
5 minute hold, followed by a ramp to 1900°C and a 10 minute hold. The pressure was slowly ramped to 50 MPa and applied from 1000°C until the end of the final hold. The micrographs can be seen at two different magnifications in Figures 4.19 and 4.20. These
140 samples returned an Archimedes density of 3.19 g/cm3 (98.6 %) and displayed very little
circular porosity. They do however display grain pull out at almost every triple point.
The actual size of the pull outs in some cases is on the scale of a few nanometers at small
isolated triple points; features of this size are not common to polishing pull out. This
could indicate the strength or thickness of the glassy phase at the triple points was
insufficient to give high to strengths. The grain boundaries seem relatively free of glassy
phase with the majority of the additives having migrated to the triple points. The contrast
and appearance of the images is interesting in comparison to other RE systems at the
same additive content. These images are shown in Section 4.3.2.4
Figure 4.19: SEM Micrographs of R48, 98 wt. % SiC, 1.29 wt. % Al2O3, 0.71 wt. % Y2O3. Note vacant boundaries and triple points.
141
Figure 4.20: In-lens image of R48, 98 wt. % SiC, 1.29 wt. % Al2O3, 0.71 wt. % Y2O3
4.3.2.2.1.2 5 weight percent
Samples were also prepared at 5 wt. % sintering aids. They were shown to be sinterable to a high density using the same cycle as the 2 wt. % samples. These
samples, Figures 4.21 and 4.22, showed roughly the same bulk densities as those at 2 wt.
% additives, 3.19 g/cm3, but because of the greater percentage of additives present represent a lower percentage (97.7%). It was expected that the greater percentages of liquid phase additives would enhance densification behavior and result in higher densities
and lower temperatures versus the samples prepared from lower additive amounts. This
is not the case however as these samples again proved difficult to optimize. The higher
amount of additives may have lead to increased volatilization and weight loss of the
additives in these cases or modified the formation of secondary phases. At 5 wt. %
additives the grain size shifts to finer sizes besides the appearance of occasional larger
grains. The grain boundary thickness increases to where plasma etching reveals clearly
142 distinguishable boundaries between the majority of the grains. The size of triple points increases and at many locations the additives have pooled.
Figure 4.21: SEM images of R52, 95 wt. % SiC, 3.22 wt. % Al2O3, 1.78 wt. % Y2O3
Figure 4.22: In-lens image of R52, 95 wt. % SiC, 3.22 wt. % Al2O3, 1.78 wt. % Y2O3
The coprecipitation process was shown to be effective in enhancing the densification rate of SPS prepared samples so it was surmised that a lower quantity of sintering aid could be added through coprecipitation than observed in literature of similar systems. As such, a small number of samples were prepared at 1 wt. % sintering aids.
143
4.3.2.2.1.3 1 weight percent
A sample was prepared with 1 wt. % additives in order to judge the effectiveness of coprecipitated additives as sintering aids, as well as to add another comparison for grain boundary properties. The measured density for this sample was 3.14 g/cm3 versus the theoretical density of 3.22 g/cm3 or approximately 97.4% dense. It is believed that
small adjustments in the amount and process could further increase the density. A review
of the published literature referenced in this thesis did not indicate samples successfully
sintered with less than 2 wt. % sintering aids when added through ball milling.
The physical appearance of this sample shows a gradient in coloration between
the edge of the sample and the center, with the center appearing a light greenish grey and
the outside edge of the sample appearing a darker more metallic silver-grey. This is an
interesting effect believed to be related to sintering silicon carbide with well mixed liquid phase forming additives as opposed to carbon. Lundquist observed silicon carbides ranging in color from yellow to green to black while performing a detailed chemical and
XRD study to determine the levels of aluminum, iron, carbon, and other impurity phases.
He determined that the lightest colored silicon carbides were most pure and did not have clusters of carbon inclusions [30]. Other researchers also were able to process silicon carbide with green coloration through liquid phase sintering without carbon [147]. In the case of this sample the low additive content, coupled with the lack of added carbon and calcination steps, allows for the center of the sample to remain carbon free and pure while the outer rim which has the darker grayish color of standard SiC, likely contains carbon that is diffusing into the sample rim from the Grafoil or graphite die.
144
The microstructure of the sample itself is interesting in that polishing and etching appear to contribute to high amounts of grain pullout for a sample of such density. The micrographs are shown below in Figure 4.23, and they appear to show that the majority of the glassy phase accumulates in triple points at this composition. This effect may be from the occlusion of the limited amount of liquid at the end of sintering. This could also imply the grain boundaries are of low strength and are easily removed through polishing or etching. Although the microstructure was of interest and it is believed the density could be improved, the remainder of this work focused on 2 and 5 wt. % systems as intermediate between too little and too much grain boundary phase.
Figures 4.23: SEM micrographs showing low contrast and vacant triple points in 1 wt. % sintering aid sample.
Samples sintered using aluminum and yttrium based sintering aids showed good sinterability in the compositional range from 1 wt. % additives to 10 wt. % additives, although slight modifications of the sintering cycles were required. Slight differences in holding time and temperature did not appear to have a drastic effect on density for many of these samples indicating either the spark plasma sintering process or coprecipitation
145 result in easily sinterable SiC with a wide range of processing variables to modify in order to change microstructures.
4.3.2.2.2 Discrepancy in density calculations
Many researchers have previously noted challenges and discrepancies when evaluating the actual and theoretical densities during the processing and sintering of silicon carbide systems [20]. Research by Ishevskyi and Bressiani indicated that it is difficult to liquid phase sinter SiC with aluminum and rare earth based additives to greater than 98 % density even with control of the sintering atmosphere and other methods [148]. Effects of weight loss, powder bed effects, wicking or skimming of liquid additives, segregation, additive losses on processing, sintering atmosphere, and the crystallinity and structure of the final grain boundary phase are all important considerations that influence the ability to determine the correct density for a given system [20, 88, 148-152].
The calculation of theoretical density is often performed through simple assumptions based on the rule of mixtures. Because the exact phase amounts or additive losses or changes during sintering are hard to account for, these calculations are by nature subject to error. An example calculation for the theoretical density of a 5 wt. % Al2O3-
Y2O3 system is shown below in equation 12:
(0.95*3.21g / cm3 )+ (0.0322*3.97g / cm3 )+ (0.0178*5.01g / cm3 ) = 3.27g / cm3 (Eq.12)
These very simplified calculations were used in determining the theoretical densities throughout this thesis. As the work progressed, it became apparent from both microscopy and phase analysis that the observed microscopic densities appeared higher
146 than those indicated from these calculations. As such, attempts to outline the causes and find corrections for this discrepancy were undertaken.
Because the calculations used for determining the theoretical density assumes the rule of mixtures, weight loss is not considered in the typical calculation. Weight loss is possible during sintering due to the reaction of the native SiO2 layer with carbon as
shown by the possible equation [77]:
SiO2+ 3C = SiC + 2 CO (Equation 13)
The volatilization of additives during the high temperature stages of sintering is another
contribution to the weight loss of the system [88]. Although coprecipitation does not add additional carbon directly to the SiC as in standard solid state sintering, any residual organics species that did not burn off during prefiring could leave behind carbon. Carbon is also present in the graphite dies, punches, insulation, and other areas of the furnace.
The reaction of the silica with the carbon can account for a portion of the weight loss and
if this was occurring, it would influence the composition of the liquid phases and their
viscosity during sintering. The maximum weight loss expected could be calculated from
the starting oxygen content of the powder, 1.8 wt. % oxygen, and would indicate a
complete loss of silica as accounting for approximately 3 weight percent. This
assumption would not be reasonable however because there is considerable error in the
calculation due to the processing involved in this work; that situation is discussed further
below. Even though applied pressure is beneficial for reducing some of these gas phase
effects, literature indicates some of the gas phase reactions and weight losses are still
relevant. Controlled and careful calculation of weight losses during processing, prefiring,
and sintering steps are necessary in order to improve the accuracy of density calculations.
147
Attempts to carefully calculate the amount of weight loss that was occurring in these samples was difficult in many cases due to the spillage and loss of powder during the loading of the smaller SPS die sets, difficulties in the accurate loading and unloading of hot pressed samples, and because of the tendency of some samples to chip around the edge during removal from the die or sandblasting to remove the graphite foil used during
SPS. On a few measurable samples however, losses of 3 to 4 percent were observed during sintering. This would account for some, but not all of the discrepancy.
A related factor that can impact that calculation of additive weight percents, liquid phase properties, and crystallinity, is changes in the native SiO2 layer on the SiC.
Although the starting percentage of SiO2 on the powder is available, many steps in the
processing such as the coprecipitation at a basic pH, the prefiring step in argon, and the
final hot pressing, can greatly change the composition and surface character and make it
difficult to track the oxygen and silica contents. Although oxygen content measurements
through the Leco method are feasible to monitor changes throughout the powder
preparation steps, accounting for changes in-situ during the sintering process is not
reasonable. Different researchers themselves approach this in a number of ways. Some
researchers assume the pressure and confinement of hot pressing keep all of the SiO2 in the sample as discussed in the paper by Can and Adler [88]. While this will slightly drop the theoretical density values since accounting for the silica (density - 2.19 g/cm3) lowers
the SiC contributions (density – 3.21 g/cm3), the result of this assumption is that some
compositions return densities over 100 percent theoretical. The percentage of silica is therefore not included at all in these initial density calculations.
148
The coprecipitation itself may be responsible for some of the increased density differences versus previously studied methods. By coprecipitating aluminum into the system as a sintering aid, Al2O3 is formed as a sintering aid as opposed to the alternative
AlN system. Alumina volatilizes more easily than AlN which can lead to higher porosity
and poorer densification under some conditions. Coprecipitation leads to very fine
sintering aids which themselves are likely more reactive and may increase volatilization.
Another likely source of the density differences connected to many of the above
issues is changes in crystallization of the grain boundary phase. The changes in SiO2 because of the non-aqueous coprecipitation procedure may not only influence the amount, but also the final crystallinity of the grain boundary phases. The fine size and amorphous nature of the coprecipitated sintering aids could lead to crystallization of the glass from the liquid phase at lower temperatures while the improved mixing and reactivity could also yield less residual crystalline phases or assist in retaining a more amorphous grain boundary. The coprecipitation process and spark plasma sintering methods used in this thesis both had a strong influence on the phase and composition results, and therefore the theoretical densities values did not seem correct for many of the studied samples.
4.3.2.2.2.1 Reassessment of Theoretical Density Values
The results of the secondary phase analysis through XRD indicated the most likely cause of the discrepancy between the theoretical density and the actual observation of pores in micrographs as being due to the assumption of certain phases forming in the initial calculations. Although this rule of mixtures approach considering the likely phase
149 formation is simple and common in research, it appears to introduce error in cases where a moderate amount of glass or mixed phases can form.
The approach of Bellosi in addressing the discrepancy does not rely on the measurement of silica but instead utilizes the XRD results to determine the amount of
YAG and extrapolates the amounts of amorphous phase from the difference between the theoretical density and the inclusion of the YAG [116]. That technique is applicable here, but requires consideration of phases other than YAG since it was not observed in most samples. The lower amount of additives and different processing methodology versus Bellosi’s work makes the quantifiable detection of phases difficult in many of these samples. From the samples where a successful detection of phases could be made it is clear that the phases and amounts are less than Bellosi observed even though they also employed coprecipitation processing. This is likely related to the differences between aqueous and non-aqueous coprecipitations as the loss of silica appears to be reduced when the reactions are performed in isopropanol, yielding more aluminosilicate phase and less YAG or other crystalline phases in these samples. By accounting for the presence of the crystalline phases, we should be able to perform a similar calculation.
Further study of the density and mechanical property values of these samples was necessary.
The earlier assumptions on the theoretical density calculations and the percentages of amorphous and crystalline phase content can be better justified through the modulus values. In the work of Bellosi et al., the rule of mixtures was utilized in conjunction with the data from the XRD to show that the lower mechanical properties of their systems were expected due to a change in the amount of amorphous phase. The
150 amount of YAG phase was determined from their XRD peak intensities. Through microscopy, Bellosi stated that the amount of porosity was low enough to establish it as the experimental. From here, they accounted for the YAG and SiC amounts from the
XRD and back calculated the amount of amorphous phase. They later used the properties of a simple silicate glass containing small amounts of certain modifiers to improve their estimates, and checked their calculation using the SiC, YAG, and amorphous phase totals.
When examining the samples prepared using yttrium systems with varying amounts of sintering aid, it was noted that the mechanical properties showed a similar decrease as that of Bellosi’s samples. Table 4.3 below shows the important characteristics of the samples from this section. The elastic moduli values are relatively insensitive with the density values of these samples. The values are also relatively insensitive to the amount of the rare earth additions. Later in Chapter 4 it will be shown this insensitivity also holds for different rare earth phases. This implies some grain boundary phase similarities between samples of similar processing.
Table 4.3: Density and moduli values for Al2O3-Y2O3 samples Density (g/cm3) E (GPa) G (GPa) K (GPa) 1 wt. % Al,Y 3.14 399 169 205 2 wt. % Al,Y 3.19 397 162 237 5 wt. % Al,Y 3.19 401 170 207 10 wt. % Al,Y 3.20 378 157 211
This also implies that it is reasonable to believe that there is a considerable
difference in phase content when compared to standard prepared silicon carbide of
similar high densities. What is interesting to note about the mechanical properties of all of these samples is that the elastic modulus values are consistently low when compared to
151 typical high density silicon carbides. High density hot pressed silicon carbides typically have moduli values in the range of 440 GPa while sintered silicon carbides are approximately 420 GPa. These samples, and as will be shown later in this chapter, other samples in this thesis, consistently revealed density values of around 400 GPa as revealed by ultrasound analysis. The accuracy of the ultrasound point measurements on samples of this size, thickness, flatness, and polish is likely to introduce some small percentage of error, but this still does not explain why samples with little porosity would have such low values. This implies a possible phase difference, and as such, Bellosi’s method may give insight into the character of these samples.
Although some of the assumptions of their method appear to be debatable, a modified method based on theirs is used here. The XRD results from this thesis indicated a very different detectability of YAG in these samples versus those in Bellosi’s work.
This complicated the procedure somewhat, but also can help account for many of the more noticeable differences in density and mechanical properties. The limit of detection study performed earlier for the newer XRD equipment indicated minority phases should be detectable to a limit of 1 wt. % for most scans. In the vast majority of the samples studied YAG was not detectable at all while in a few samples a maximum of 1 weight percent was observed based on the peak intensities.
A second improvement in the calculations of this thesis versus Bellosi’s is in the density and mechanical properties of the amorphous phase. Bellosi utilized the density
(2.5 g/cm3) and elastic properties (E = 73 GPa) of a very simple silicate glass with low
amounts of alumina and sodium. But, if the amount of detectable YAG is very low and
crystalline alumina is not detected, the phase present at the boundaries is likely an
152 aluminosilicate glass containing a moderate amount of yttrium. Although some of the yttrium is expected to form phases such as crystalline Y2SiO5 and Y2Si2O7, the XRD
results indicated these are detectable in low amounts and could be accounted for when
detected. In samples where mullite was also detected, the open structure of mullite could
accommodate some yttrium, but from literature it does not appear to be a large amount.
It is therefore more likely that most of the grain boundaries are aluminosilicate glass
containing yttrium as a glass modifier.
If the approach of Bellosi is used here, but with the phases we can detect and
identify, the lower mechanical properties and densities can be explained. Tables 4.4 and
4.5 below show the density calculations and elastic modulus values for the yttrium
systems studied while considering the phase contents from XRD.
Table 4.4: Reassessment of theoretical density values for yttrium containing samples Theoretical Density Calculations Composition Theoretical Density (g/cm3) Wt% Wt% Wt% If all If all Y- Accounting Accounting for Sample SiC Al2O3 Y2O3 YAG Al glass for mullite Al2O3+Y2Si2O7 R61 90 6.40 3.60 3.34 3.21 3.23 3.28 R52 95 3.22 1.78 3.28 3.21 3.22 3.25 R48 98 1.29 0.71 3.24 3.21 3.21 3.22 R56 99 0.64 0.36 3.22 3.21 3.21 3.20
Table 4.5: Measured and expected elastic modulus values if theoretical density, phase differences, and porosity are evaluated Density (g/cm3) Elastic Modulus Values (GPa) Percent All All Glass + Including Sample Density Theoretical Dense Measured Mullite Glass Mullite Porosity R61 3.20 3.23 99.0 378 388 384 386 382 R52 3.19 3.22 99.1 401 401 400 400 396 R48 3.19 3.21 99.3 403 408 408 408 403 R56 3.14 3.21 97.7 399 411 410 412 402
153
The chart and calculations outlined above explain why the density values and mechanical property numbers of the 2 and 5 wt. % yttrium samples are so similar. The compositions containing mullite which has an approximate density of 3.2 g/cm3, and a yttrium modified glass, density of 3.70 g/cm3 have a lower theoretical density than
expected and with the amounts present the theoretical density changes little between 2
and 5 weight percent. In samples containing YAG, density 4.55 g/cm3, or excess
alumina, 3.97 g/cm3, and/or yttrium disilicate, density 4.03 g/cm3, the value scales a bit
higher. Because of the circumstance of the numbers being reasonably close to that of
SiC, the theoretical density does not shift by a large amount even if the percent increases
from 2 to 5 weight percent. The small differences in density and porosity between the
samples therefore provide the remaining influence, leading to very similar elastic
moduli, even considering the lower amount of glassy phase in the 1 wt.% sample.
These results help to reinforce the previous observations and discussions about
the glassy grain boundary character of these materials and the importance of mixing in
phase development.
4.3.2.3 Effect of Sintering Method and Conditions
4.3.2.3.1 Microstructure Changes with Varying Hot Press Heating Schedule
Yttrium based systems are a common and effective sintering aid for SiC, and have
been researched over a relatively large range of compositions by a considerable number
of research groups. The coprecipitation process has been observed to be successful in
improving the mixing of these sintering aids and results in modified phase content. There
are other variables in microstructure control that need to be characterized to really
154 understand how best to influence, control, and design microstructures. It is therefore necessary to understand the differences in phase developed from densification rate effects, holding times, and heating methods.
A large batch of powder was coprecipitated with the composition of 98 wt. % SiC and 2 wt% additives in a molar ratio of 4 Al:Y, yielding a final theoretical composition of
98 wt. % SiC, 1.71 wt. % Al2O3, and 0.29 wt. % Y2O3. The batch was then split into
three smaller batches which were subjected to different firing schedules in the hot press.
The results of the different firings could be used to infer if the comparison samples
mentioned previously had such similar results due to the sintering temperature of the
coprecipitated sample being too high. A table showing the different heating schedules and results is shown below in Table 4.6. Micrographs of the samples sintered using different schedules are shown in Figures 4.24, 4.25, and 4.26.
Table 4.6: Density and grain size results from coprecipitated powder batch split into 3 smaller batches and densified using different hot pressing cycles Grain Hot Press Density Percent Sample Composition Size Cycle (g/cm3) Theoretical (μm) 98 wt. % SiC, 1.71 % Al O , 1950°C, 30 min, 1.33± R43-1 2 3 3.15 97.5% 0.29 % Y2O3 35 MPa 0.14 98 wt. % SiC, 1.71 % Al O , 1900°C, 1 hr., 0.99± R43-2 2 3 3.16 97.7% 0.29 % Y2O3 35 MPa 0.12 98 wt. % SiC, 1.71 % Al O , 1850°C, 1 hr., 0.96± R43-3 2 3 3.08 95.4% 0.29 % Y2O3 35 MPa 0.14
Even at low firing temperatures of 1850°C coprecipitated samples with 2 % additives, Figure 4.26, displayed good densification behavior resulting in 95 % with a 1 hour hold. This is indicative of the effectiveness of coprecipitation in sintering as
1850°C is a relatively low sintering temperature for similar LPS compositions especially
considering the low additive content. The research of She et al. demonstrated a decrease
in sintered density with reduction of additives to 5 wt. % at temperatures of 1850°C [86].
155
Samples fired to 1900°C for 1 hour and 1950°C for 30 minutes, Figures 4.24 and
4.25, showed improved densification up to 97.7 and 97.8 % respectively. This wide range of processing temperatures enabled an in depth look at some of the changes possible in microstructures through the influence of processing. Scanning electron microscopy revealed differences in the core-rim development between the three samples subjected to different temperatures and times. The samples at 1850°C and 1900°C for 1 hour both displayed a considerable amount of grains showing a clear and distinct core- rim microstructure, while the sample fired at 1950°C does not. This observation could be explained by a lack of appreciable core-rim development because of the higher temperature and shorter sintering time. On the opposite end of the spectrum, it could also indicate complete dissolution of the parent SiC grains leaving no discernable observation of core-rim behavior.
Another interesting feature of the sample sintered at 1950°C is in the appearance of the grain boundaries. The boundaries in the samples sintered at lower temperatures appear raised after plasma etching, likely indicating the presence of an appreciably thick glassy phase along the boundary between SiC grains. In the sample hot pressed at
1950°C, many of the SiC-SiC grain boundaries appear to lack a measurable grain boundary phase thickness, evidenced by the flatness and lack of contrast in the image.
The majority of the glassy phase appears to have concentrated into larger triple points.
The sample hot pressed at 1900°C appears to have the most uniform distribution of grain boundary phase and smaller well distributed triple points, while the sample processed at
1850°C appears intermediate to the other two, with some areas showing clear boundaries and others appearing to have grain boundary phase present.
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Figure 4.24: SEM images of coprecipitated sample R43-1, 98 wt% SiC, 1.71% Al2O3, 0.29% Y2O3. Sample was hot pressed at 1950°C for 30 minutes
Figure 4.25: SEM images of coprecipitated sample R43-2, 98 wt% SiC, 1.71% Al2O3, 0.29% Y2O3. Sample was hot pressed at 1900°C for 1 hour
Figure 4.26: SEM images of coprecipitated sample R43-3, 98 wt% SiC, 1.71% Al2O3, 0.29% Y2O3. Sample was hot pressed at 1850°C for 1 hour
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While it can be established that there were clear differences between the samples above, a question was raised about differences observed between these coprecipitated and hot pressed samples versus spark plasma sintered samples from earlier in the study. A set of samples was therefore designed to investigate what effect the sintering method has on the phase and microstructure development.
4.3.2.3.2 Differences Developed between HP and SPS
While the procedures sections in Chapter 3 depicted the methods of hot pressing versus spark plasma sintering samples, it is important to understand some of the key differences between the two. Table 4.7 below compares the two sintering methods.
Table 4.7 Comparison of HP and SPS Methods Hot Press Spark Plasma Heating Method Induction Joule or boundary (theory) Heating Rate <50°C / minute Up to 1000°C / minute Typical Cycle Time 4 to 6 hours 15 – 25 minutes Pressures Lower, limited by die Higher Pressures
4.3.2.3.2.1 Density and Grain Size
Experiments on the limit of detection of the newer x-ray diffraction system
indicated that it should be possible to determine differences in phase content in samples
with less additives. A series of scans were performed on samples prepared with 5 wt. %
sintering aids but processed and sintered under different conditions. Sample BM5 was
prepared through ball milling of the yttria and alumina sintering aids and was spark
plasma sintered. The same weight percentage of yttria and alumina was added to sample
R52 through the coprecipitation procedure and this sample was also spark plasma
sintered. Finally, a third sample R31, was coprecipitated similarly to R52 but was
158 densified by hot pressing. The density results have been covered in detail previously but are summarized below in Table 4.8. Microscopy of these samples was covered in previous sections.
Table 4.8: Density and grain size values for 5 wt. % samples Sample Density Grain Size BM5 95.4 0.76±0.08 R52 97.6 0.61±0.10 R31 97.4 0.95±0.16
4.3.2.3.2.2 XRD
The results of the XRD scans reveal very similar results for polytypes and indicate similarity in features for the majority SiC phase. There are subtle differences in the crystallinity and phase content of the minor phases however. Figure 4.27 below compares the XRD patterns for sample R52 and BM5. Both samples are of the same composition, 5 wt. % additives, but sample R52 was prepared from the coprecipitation of the sintering aids while BM5 was ball milled using the oxide forms of the aluminum and yttrium. The samples were both spark plasma sintered with the same cycles. As opposed to earlier scans with less sensitive XRD systems, the appearance of small amounts of a number of crystalline phases is now observed. In both samples, the easiest phase to detect is mullite, or a similar aluminosilicate phase, as strongly indicated by the broad peak at 26°. There are a number of other peaks that seem to be from the same phase in both samples, but can not be identified. There are also clear differences in the intensity of some peaks between the two samples, namely the appearance of a strong peak in the coprecipitated sample at 13°, while in the ball milled sample a strong peak is visible at about 18°.
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Figure 4.27: XRD scans comparing 5 wt. % additive samples from coprecipitation and ball milling
Figure 4.28 compares the scans from sample R52 to sample R31. In this case, both samples were coprecipitated, but sample R31 was hot pressed for 1 hour at 1900°C instead of spark plasma sintering. The minor phase peaks observed here are more similar than those between the ball milled and coprecipitated samples, but it is also apparent that the more rapid densification achieved in the spark plasma sintering process does influence the crystallinity of the grain boundaries in silicon carbide. The intensity of the peaks is stronger in the SPS processed sample R52, but the peaks are also broad, indicating the fine size of these phases including the SiC, in the coprecipitated and spark plasma sintered material.
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Figure 4.28: XRD scans comparing coprecipitated 5 wt. % additive samples relative to heating method, SPS versus hot pressing
Figure 4.29 compares all three of the samples allowing for an overall comparison of the important differences. It can be observed that coprecipitation introduces subtle changes in the grain boundary phase and crystallinity versus standard ball milling processes. It is just as important to note however, that the heating rate and densification method have a strong effect on the phase development as well. Samples prepared through ball milling and spark plasma sintered display peaks very similar to coprecipitated samples that were densified using a long heating cycle in a hot press.
When a sample is both coprecipitated and densified quickly in the SPS, a very different
XRD pattern is observed with sharper, broader peaks, and the emergence of some new phases. The rapid heating of the spark plasma sintering process also appears to result in lower losses of volatile species during sintering as evidenced by the strong broad mullite
161 peak present in both the ball milled and coprecipitated samples that were spark plasma sintered.
Figure 4.29: XRD scans comparing 5 wt. % additive samples relative to processing and densification method.
From the XRD results it was concluded that the non-aqueous coprecipitation procedure did indeed lead to different grain boundary phases than initially expected and observed in other authors work. It was also observed that these changes can be further influenced by the rapid, efficient heating of the spark plasma sintering system. The improved processing and densification methods allow for the development of fine grained silicon carbide containing a glassy grain boundary phase with varying species and degrees of crystallinity among the minor phases.
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4.3.2.4 Grain Boundary Modification through Varying Rare Dopants
As much of the relevant literature on mechanical properties and grain boundary engineering of LPS-SiC systems showed favorable results for introducing phases other than yttria, it was decided to explore the effects of other rare earth dopants in coprecipitated and/or spark plasma sintered systems. In addition to the previously discussed samples containing yttrium, samples with three other rare earth elements, lanthanum, gadolinium, and samarium, were prepared. These elements were chosen based on differences in cationic radius. As a number of researchers have shown in previous work, changes in the rare earth element have an appreciable effect on the mechanical properties of the system because of their ability to modify the thickness of the grain boundaries and influence the properties of the glass [12, 14, 16, 55]. A table of the different radii and information on the rare earths is shown in Table 3.1. Although the molar ratio of alumina to rare earth is a controllable variable in any of these systems, it was necessary to limit the amount of samples at this stage. Samples with a high aluminum to rare earth ratio were chosen at this juncture, as literature indicated this system should have good mechanical properties [88].
4.3.2.4.1 Hot Pressed Samples
4.3.2.4.1.1 Gadolinium based samples
A composition containing 95 wt. % SiC, 2.65 wt. % Al2O3, and 2.35 wt. % Gd2O3 was prepared to investigate the effectiveness of coprecipitation on other rare earth systems of varying cationic radius. Literature on the usage of gadolinium as a sintering aid in SiC systems is limited [135, 153].
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The sample, shown in Figure 4.30, displayed an interesting change in coloration from the yttrium samples. While the yttrium samples were typically light grey in color, similar to SiC obtained from standard processing, visually the gadolinium samples exhibited a darker green color. Upon polishing the sample surface appears more similar to conventional polished SiC but slightly lighter. This may be related to an absence of carbon in these samples, but also is influenced by the presence of gadolinium.
Densities as high as 96 percent of theoretical were achieved in this composition but results were inconsistent, indicating a somewhat poorer densification behavior for gadolinium samples. The grain size of these samples is relatively fine, 0.74 microns.
This may help explain the reasonable value for 1 kg. Knoop hardness of 19.0 GPa.
Figure 4.30: SEM micrograph of sample R32, 95 wt. % SiC, 2.65 wt. % Al2O3, and 2.35 wt. % Gd2O3
4.3.2.4.1.2 Lanthanum Samples
The results on hot pressing lanthanum based samples revealed poorer sintering
behavior than yttrium based compositions. Based upon considerations of the ionic radius,
it is expected that the size and refractory nature of lanthanum should result in more
difficult sintering [148]. Even at higher sintering temperatures of up to 2000°C,
164 lanthanum samples could only approach 94.8% density after similar cycles to those for the yttrium with many samples failing to densify above 88%. Because of these results and its high cost, samarium was not attempted in the hot press at this time.
Literature showing densification curves of lanthanum based samples indicate a very different cycle may be required to achieve full density for these samples [148].
Changing the atmosphere in the furnace from argon to nitrogen appears to be a potential solution, but it would be preferred to be able to sinter all of the compositions to high densities under similar conditions.
4.3.2.4.2 Spark Plasma Sintering
The previous work with hot pressing the coprecipitated samples revealed some interesting differences, but the results did not seem to necessarily indicate drastic changes in sintering behavior or mechanical properties. The poor densification of some compositions in the induction hot press could have been overcome with drastically different heating cycles or modified furnace atmospheres. The small differences observed in properties between coprecipitated samples and ball-milled or propeller mixed samples suggested some small benefits in grain boundary phase crystallinity and grain size, but the overall densification behavior was not as enhanced as expected. Other researchers have also observed this in previous work where coprecipitation has not drastically lowered the temperature required for high densities, but has enhanced the kinetics to allow for more rapid sintering behavior [115, 116]. With the focus on grain size and crystallinity behaviors as important issues in influencing the properties of the
165 grain boundaries, it was decided that spark plasma sintering would be more appropriate for the remaining work.
The short cycle times and rapid densification behavior of spark plasma sintered samples could enable the reduction of grain sizes to even finer values where the influence of the grain boundaries on performance would increase. The reduced time for the diffusion of the sintering aids and rearrangement of the liquid could also be exploited for more control of the microstructure development. The small, reactive sintering aids introduced during coprecipitation, when coupled with the innovative heating of SPS, could lead to quicker densification, further reduction in SPS temperature, and new and different grain boundary characteristics.
Samples were prepared through the coprecipitation procedure described previously and contain alumina and one of the following rare earths oxides: yttrium, lanthanum, gadolinium, and samarium. Compositions containing two weight percent additives and five weight percent additives were prepared for all four of the systems of interest. The yttrium samples are included in the tables and the discussion for purposes of comparison. Micrographs for the yttrium samples were shown previously in section
4.3.2.2. The matrix of samples was intended to characterize the differences in materials relative to coprecipitation processing, varied rare earth elements, and other changes introduced through spark plasma sintering.
4.3.2.4.2.1 Densification
Use of the SPS as a densification method resulted in high calculated densities, greater than 95% based from the theoretical calculations for all rare earth systems and
166 both additive amounts. Table 4.9 below shows the exact composition of the samples prepared from the various rare earths and at the two different percentages while Table
4.10 summaries the density results from these samples.
Table 4.9: Matrix of coprecipitated compositions studied in spark plasma sintering of rare earth modified samples
Composition 2 wt. % additives 5 wt. % additives
Y 98 % SiC, 1.29 % Al2O3, 0.71 % Y2O3 95 % SiC, 3.22 % Al2O3, 1.78 % Y2O3
La 98 % SiC, 1.11 % Al2O3, 0.89 % La2O3 95 % SiC, 2.78 % Al2O3, 2.22 % La2O3
Gd 98 % SiC, 1.06 % Al2O3, 0.94 % Gd2O3 95 % SiC, 2.65 % Al2O3, 2.35 % Gd2O3 Rare Earth Sm 98 % SiC, 1.08 % Al2O3, 0.92 % Sm2O3 95 % SiC, 2.69 % Al2O3, 2.30 % Sm2O3
Table 4.10: Density results for matrix of coprecipitated samples Composition 2 wt. % additives 5 wt. % additives Density Density Archimedes Theoretical Percent Archimedes Theoretical Percent (g/cm3) (g/cm3) Dense (g/cm3) (g/cm3) Dense Y 3.19 3.23 98.8 3.20 3.27 97.9 La 3.13 3.25 96.3 3.17 3.30 96.1 Gd 3.19 3.26 97.9 3.21 3.33 96.4 Sm 3.19 3.27 97.6 3.19 3.35 95.2 Rare Earth
Densities ranged from 3.13 g/cm3 for lanthanum containing samples to 3.19 g/cm3 for yttrium, gadolinium, and samarium based samples. The conventional theoretical density is shifted slightly depending upon the mass of the rare earth oxide species, from
3.233 g/cm3 for the yttrium based samples to 3.265 g/cm3 for the gadolinium samples,
making the percentages shift from 96.1% for the lanthanum sample to 98.9% for the
yttrium sample. These numbers are calculated assuming the rule of mixtures and are not
correct. Correcting them here is more difficult because the XRD on the more exotic rare
earths required better card files for careful phase matching. Observations of porosity
167 through microscopy in these samples indicate consistent and high densities similar to the yttrium samples.
Although it was initially believed that the SPS may allow for a considerable decrease in the sintering temperature versus that employed in the induction hot press, the results indicated this was not always the simple case for these systems of interest. The rapid heating rate and efficient heating did consistently allow for a shorter cycle time and yielded samples with finer average grain sizes. Initially, it was attempted to process all the different rare earth samples using the same cycle. It became apparent that the wide variation in liquid phase content, as well as the differently sized rare earth ions required slightly modified sintering schedules in order to achieve suitable densification. The required temperatures for densification as well as the location and duration of the holds was not straightforward to optimize as previously thought, nor were they consistent between different samples. Where as small shifts in temperature or holding time during hot pressing could result in very drastic shifts in the final density, the spark plasma sintering system proved to be much more forgiving overall; most cycles yielded high densities with a smaller deviation than observed in hot pressing. While this is promising, it also makes optimizing SPS densified samples of this type a more daunting task. The full optimization of the samples through SPS processing requires a detailed and specific study of the densification and heating behaviors native to spark plasma heating systems which was not the focus of this study. It is important to note the different cycles employed for densification between the different samples when comparing some of the features.
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Table 4.11 below shows some of the density values obtained from various batches of powder that were split and fired using different cycles in order to determine if the coprecipitation did result in samples densifying at lower temperatures or shorter times.
Table 4.11: Differences in densification and elastic properties with small changes in final SPS holding temperature
Density E G K Sample Composition (g/cm3) υ (GPa) (GPa) (GPa)
1850 95% SiC, 2.65% Al2O3, 2.35% Gd2O3 3.22 0.198 381 159 210
1875 95% SiC, 2.65% Al2O3, 2.35% Gd2O3 3.23 0.222 413 169 247
1850 95% SiC, 2.69% Al2O3, 2.30% Sm2O3 3.16 0.208 386 160 220
1875 95% SiC, 2.69% Al2O3, 2.30% Sm2O3 3.21 0.207 404 167 230
Samples prepared from rare earth elements at 2 wt. % total additive content,
Figures 4.31 – 4.33, appeared to have more distinguishable grain boundaries than the
yttrium based samples. Comparing these samples should expose any differences in grain
boundary properties develop among the varying sizes of rare earth elements.
Figure 4.31: SEM image of R62, 98 wt. % SiC, 1.08 wt. % Al2O3, 0.92 wt. % Sm2O3
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Figure 4.32: SEM image of R62, 98 wt. % SiC, 1.08 wt. % Al2O3, 0.92 wt. % Sm2O3
Figure 4.33: In-lens image of R62, 98 wt. % SiC, 1.08 wt. % Al2O3, 0.92 wt. % Sm2O3
As shown in Figures 4.34- 4.40 below, at 5 wt. % additives the grain size shifts to finer sizes besides the appearance of occasional larger grains. The grain boundary thickness increases to where plasma etching reveals clearly distinguishable boundaries
170 between the majority of the grains. The size of triple points increases and at many locations the additives have pooled.
Figure 4.34: SEM images of R55, 95 wt. % SiC, 2.65 wt. % Al2O3, 2.35 wt. % Gd2O3
Figure 4.35: In-lens image of R55, 95 wt. % SiC, 2.65 wt. % Al2O3, 2.35 wt. % Gd2O3
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Figure 4.36: In-lens image of R55, 95 wt. % SiC, 2.65 wt. % Al2O3, 2.35 wt. % Gd2O3. Infrequent regions of large grains such as this observed in some SPS samples.
Figure 4.37: SEM image of R54, 95 wt. % SiC, 2.69 wt. % Al2O3, 2.31 wt. % Sm2O3
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Figure 4.38: SEM image of R41, 95 wt. % SiC, 2.78 wt. % Al2O3, 2.22 wt. % La2O3
Figure 4.39: In-lens image of R41, 95 wt. % SiC, 2.78 wt. % Al2O3, 2.22 wt. % La2O3
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Figure 4.40: In-lens image of R41, 95 wt. % SiC, 2.78 wt. % Al2O3, 2.22 wt. % La2O3
4.3.2.4.2.2 Elastic, Bulk, and Shear Moduli
An earlier examination of the moduli values of samples prepared from varying
yttrium contents revealed very similar elastic modulus values at different additive levels.
It was decided to continue this analysis using the other rare earth elements in varied
amounts as well. This could help establish that the compositions are forming rare earth
modified glass as expected since XRD was more difficult in these cases.
Table 4.12 below summarizes the important values returned from these tests. The values trend well with the density values of the majority of these samples until they densities get sufficiently high; the values are then observed to plateau at 400 GPa as with
the previous yttrium samples. The values are also relatively insensitive to the different
rare earth additions. This again implies some grain boundary phase similarities between
samples of similar processing.
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Table 4.12: Density and moduli values for various samples throughout this thesis Density (g/cm3) E (GPa) G (GPa) K (GPa) 2 wt. % Al, La 3.13 375 157 205 5 wt. % Al, La 3.17 395 165 217 2 wt. % Al, Gd 2.88 323 137 167 5 wt. % Al, Gd 3.22 392 165 211 1 wt. % Al, Y 3.14 399 169 205 2 wt. % Al, Y 3.19 397 162 237 5 wt. % Al, Y 3.19 401 170 207 10 wt. % Al, Y 3.20 378 157 211 5 wt. % Al, Y Ball Milled 3.12 377 153 240 2 wt. % Al, Y Ball Milled 3.15 381 154 243
As highlighted previously, it is interesting that the elastic modulus values are
consistently low when compared to similar structural silicon carbides. While commercial
silicon carbides typically have moduli values in the range of 420 to 440 GPa, most
samples in this thesis, regardless of rare earth, had E values of around 400 GPa as
revealed by ultrasound analysis.
Earlier it was discussed at length that the coprecipitation process and rapid
heating of spark plasma sintering both had effects on the phase and crystallinity of the
grain boundaries. These differences are the most likely cause of the mechanical property
differences. As discussed previously, Can et al. processed silicon carbide samples with rare earth and aluminum sintering aids through hot pressing and gas pressure sintering and compared the densification, chemistries, and mechanical properties. They observed significant differences in the elastic moduli between the differently sintered samples which they attributed to the presence of silica. While gas pressure sintering, the open chamber and longer cycles cause silica to be volatilized resulting in the final grain boundary phase to be aluminate based. In the case of hot pressing, the external pressures
175 reduce the volatilization and develop silicate based boundaries instead. The mechanical properties, namely E in this case, are lower in the silicate materials and the overall modulus of the SiC containing a grain boundary phase is therefore lowered.
This reasoning can be applied to the samples in this study as well. The use of hot pressing and spark plasma sintering are both methods in which the sample is contained in a tightly fitting die and pressure is applied causing the retention of silica and reduction of weight loss in these materials. Silicate phases would therefore be expected and observed as the XRD results throughout this work showed. These silicates generally have lower mechanical properties and are therefore likely to lower the values compared to standard
SiC. The fact that this pattern held for samples with different rare earth elements and various doping levels further supports the presence and important of the glassy phases in these silicon carbide materials and their strong influence on the overall properties of these samples.
4.3.3 Chapter 4 Summary
While Chapter 3 outlined the development of an improved coprecipitation procedure capable of preparing silicon carbide samples with better mixing and fewer large defects, it only suggested the possibility of major microstructural differences that could exist because of the modified processes and mixing. The work of Chapter 4 aimed to investigate these microstructural differences developed in coprecipitated samples relative to the processing changes, additive additions, and changes in sintering cycle that are enabled because of the coprecipitation process. Since the samples in Chapter 3 revealed less phase information than desired because of their low additive content,
176 samples were prepared with 10 wt. % additives from conventional ball milling and coprecipitation. These samples showed clear and distinct differences in crystalline phases. The ball milled samples showed discrete alumina and yttrium silicate phases which implied a finite scale of mixing as the residual alumina particles were likely from larger powder particles from the initial sintering aid additions. The coprecipitated sample showed a non-equilibrium phase consisting of a 2:1 mullite. The lack of detectable yttrium phases in this sample, as well as the low yttrium silicate content in the ball milled sample, were determined to be due to the majority of the yttrium phase being contained by the grain boundary glassy phase present in both materials and confirmed from XRD and EDS analysis.
Samples were also prepared from coprecipitation at different additive levels in order to determine if a reduced additive content could return favorable properties in this system. At both 2 and 5 wt. % additives in the aluminum-yttium sintering aid system, coprecipitated samples yielded high densities and comparable properties. Samples with 1 wt. % additives were also attempted in this system with some interesting results.
Densities and appearance of these very low additive samples were variable likely due to the low additive contents and changes in carbon contents because of the gradients in liquid phase contents. Otherwise, the densities and properties of samples prepared from
2 to 10 wt. % additives were consistent with a successful LPS-SiC material.
Microscopy of these samples always revealed a higher bulk density than the theoretical calculations made from the rule of mixtures from equilibrium phases was indicating. This further reinforced the belief that the phase development in some of these samples was far from the equilibrium. While many other researchers had reached these
177 conclusions as well, it was interesting to note that the coprecipitation yielded phases not commonly observed in other work. A simple reevaluation of the phases comprising the sample, determined from the XRD and other analyses, allowed for a secondary recalculation of the expected densities and much better results relative to the observed values. By circumstance of the glassy and crystalline phases present, the densities for samples with increasing additive amounts still averaged to be near the theoretical density for SiC and this explained why the measured densities tended to plateau with additive content in these aluminum-yttrium samples although the theoretical densities expected an increase. It also confirmed why the observed densities were consistently high. In order to further confirm this observation, measurement of the elastic modulus were predicted from a rule of mixtures approach using the known phases and shown to be consistent with measurements from ultrasonic evaluation.
Because various sintering methods were used throughout this thesis, a characterization of the differences imparted in these materials from modifications in sintering method and cycle was undertaken. Spark plasma sintering was developed as a superior method of sintering coprecipitated samples because its rapid heating and short cycles was best able to take advantage of the improved additive mixing and finer scale of sintering aids. Although hot pressing cycles were also shown to be reasonably effective for densifying these samples, results on the grain sizes, densities, and grain boundary phases indicated that coprecipitated samples could develop more interesting samples when spark plasma sintering was employed.
After the success of preparing samples from the well understood aluminum- yttrium system, it was determined to further explore the potential of coprecipitation to
178 easily prepare LPS-SiC materials containing other rare earth dopants. Three other rare earth systems were chosen based upon cationic radii differences: gadolinium, samarium, and lanthanum. Results on these systems with ball milled samples displayed poor sinterability in hot pressing and SPS trials. Coprecipitation greatly improved the densification behaviors of these systems, but densities were still lower than desired when hot pressing. It was observed and confirmed in literature that varied rare earths often require drastically different sintering atmospheres and cycles. The spark plasma sintering of coprecipitated systems from these compositions at 2 and 5 wt. % additives indicated reasonable sinterability with generally similar SPS cycles.
The key observations of this chapter were that the coprecipitation process did not just make samples with better mixing than ball milling, but because of the improvements in mixing and additive scale, coprecipitation was able to develop samples with very different grain sizes and phase constituents. When coupled with the rapid and innovative heating of spark plasma sintering, coprecipitation was observed to be able to generate silicon carbide with glassy grain boundaries with limited crystalline contents. It was also seen that the coprecipitation could be easily scaled to other additive simple additive systems and possible enable the creation of samples with further reduction in sintering aid content. After making these observations it was hoped that strong differences in mechanical properties could be observed in Chapter 5.
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Chapter 5:
Effect of Improved Coprecipitation Processing on Hardness and Plasticity
5.1 Introduction
While the previous chapters outlined the coprecipitation process and interrogated some of the phase and microstructure features, very few results indicated any significant differences in bulk properties. Small differences in hardness were measured between a few samples, but much of the difference could have been due to reduced grain sizes. The variations in elastic modulus were relatively small and shown to be related to phase content, but did not seem to demonstrate a drastic change in behavior because of the differences in structure and phase in the microstructure. So, while it has been exhibited that coprecipitation can developed SiC materials with observable mixing and microstructure differences, the significance of these differences has yet to be established.
While the primary focus of this thesis was in the development of the method and establishing the primary differences, it was hoped that in this Chapter a limited study of mechanical and physical properties that could be relevant for structural applications would reveal some significant benefit to coprecipitation.
5.2 Importance of plasticity
Much work on improving the mechanical properties has focused on two of the properties with the best potential for yielding a large increase in SiC ceramics performance, namely fracture toughness / fracture path modification and effective plasticity. Conventional methods of improving the fracture toughness in ceramics rely on
180 increasing the crack path or thermal stresses through second phase inclusions or anisotropic grain growth. Much literature on these methods and their success in increasing fracture toughness in SiC can be found [35, 102, 103, 138].
The high strain rates and loads involved in the usage of SiC as a structural ceramic may dictate that conventional strengthening techniques could be detrimental.
The large, acicular grains typically seen in a toughened microstructure act as stress risers in the microstructure and can initiate failure under high loads or strain rates. Second phase inclusions and particles also develop thermal stresses in the microstructure upon cooling and also are believed to be deleterious to performance [39]. The development of fracture toughness through alternative methods without elongating grains is less developed.
The development of plasticity in ceramics in general has been difficult with only certain systems showing definitive plasticity and most of those rely upon higher temperatures to influence plasticity mechanisms. Pearson in the 1930’s observed elongations of hundreds of percents in fine grained metals at certain temperatures and strain rates. The ability of typically less ductile materials to display this type of elongation behavior was coined “superplasticity” [154]. These behaviors are less than common to many systems and have not been exploited for many applications.
Developing plasticity in SiC is a difficult task using conventional processing techniques. Many high performance SiC materials are processed with solid state boundaries showing little intergranular phase and display brittle, transgranular fracture without much slip plasticity upon compressive loading. Fine grained ceramic materials containing a liquid phase have been shown to display superplastic behavior [154].
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Ceramics that undergo martensitic transformations have also displayed surprising low temperature plasticity [155, 156]. Some SiC materials are processed in such a way to develop thin amorphous boundaries phases or annealed to crystallize the boundaries and triple points. These boundaries can change the fracture path and improve the plasticity behavior, but much of their development is centered on their ease of processing and high temperature creep properties, not for reducing brittleness and improved room temperature behaviors [32, 157-159].
The ability to process SiC with varying grain boundary characteristics such as composition and thickness opens up the possibility of engineering the grain boundaries of silicon carbide to exploit and develop favorable properties and behaviors. The addition of rare earth dopants and sintering aid control may also allow for the stacking fault and twinning behaviors to be influenced as a method to further influence plastic flow.
5.2.1 Measuring Plasticity
The mechanisms and methods for measuring plasticity in ductile metals under static or quasi-static loading conditions is well developed, but assessing the plasticity behaviors of ceramic materials is less established. Indentation hardness with sharp indenters, such as Knoop or Vickers, has long been accepted as being a simple method for determination of plastic properties of a material, but the results are not conclusive or simple to interpret with brittle materials. A hardness number alone does not reveal much about the plastic deformation capacity or mechanism of the material; only its inherent ability to resist penetration which contains both elastic and plastic components [160].
Conventional test methods for metals such as stress strain curves and bend tests are not
182 common or thought to be applicable for testing ceramics. Only isolated systems have lent themselves to easy evaluation through these methods. Shimazu developed Al2TiO5-
MgTi2O5 ceramics to mimic the microstructures observed in the highly plastic naturally
occurring rock Itacolumite. Three point bend tests were used to generate stress strain
curves, and the ductility of the ceramic was interpreted from these diagrams [161].
Simple comparison of mechanical properties such as hardness, elastic, shear, and
bulk moduli, and Poisson’s ratio can reveal information about differences in brittle versus
plastic behaviors. Bulk metallic glass (BMG) systems have shown promise in some
structural applications, but much like SiC, are limited because of their inherent
brittleness. A study by Yu and Bai revealed that improved plasticity behavior in BMG’s
can be easily compared through tracking the shift in Poisson’s ratio or the ratio of the
shear to bulk moduli. Modified compositions and processing lead to observable
differences in these ratios by ultrasonic analysis. Stress strain curves confirmed the
changes in plastic strain. Although this comparison could be introduced simply in SiC
systems through ultrasonic NDE analysis, the variability in processing and defects in
ceramics could result in a less clear predictive method [162].
The development of load-displacement curves is a much more relevant measure
of plasticity, but with conventional sharp indent hardness tests the high initial load at a
single point causes local plastic deformation even at low loads and severe sample
damage, historically making both measurement and evaluation difficult [163].
Researchers have progressed towards both qualitative and quantitative comparison
methods using hardness indentation as a measure of plastic behavior. Lawn and Marshall
devised a brittleness parameter by considering the magnitude of the loading required
183 during a hardness indent to transition a material from leaving a permanent hardness indent to the onset and growth of cracking. All materials can be classified under these conditions as under low loads, hardness behavior dominates and at high loads where cracking is often inevitable for ceramics, cracking behaviors become more essential. The ratio of the hardness determined from a Vickers indentation to the fracture toughness then becomes a loose metric for comparing the brittleness of different materials [164]. Quinn and Quinn performed a similar analysis using energy considerations and developed a similar brittleness measurement relating the hardness, H, Young’s modulus, E, and the fracture toughness, KIc [165]. ()HE B = 2 (Equation 14) K Ic
Recently, work by McCauley and Wilantewicz developed a simple load-hardness
curve analysis method that results in a comparable plasticity number. By taking a load hardness curve over a wider range of loads than is often considered, they were able to fit the shape of the curve to a power law, displaying a much better curve fit than the previously developed equations. By developing a log-log plot of the power law equation a straight line can be plotted and the inverse slope of this line was introduced as a plasticity parameter, allowing for simple comparison of various materials. In order to better demonstrate the validity of this method, the plasticity parameter was summed with the Knoop hardness at 1 kg and this value was correlated with transitional velocity from ballistic testing experiments on a number of SiC, Al2O3, and B4C materials. The
predicted values showed very strong agreement with the actual test values. This sort of
184 analysis appears to be a simple and strong method for determining plasticity and relating it to material microstructure differences [97].
Hertzian indentation testing is believed to be a much more pertinent testing method for determining plasticity relative to compressive loading. A blunt or spherical indenter develops a much more gradual rise in stress and keeps the damage to a minimum allowing for the observation of plastic behavior. The load-displacement curves developed from Hertzian indents can reveal information about plastic response. It is also possible to develop indentation stress-strain curves which can show the elastic, transition, and plastic responses of a material [166].
5.3 Procedures
5.3.1 Hardness
The procedure used for basic Knoop hardness measurements was previously described in Chapter 3, Section 3.5.9.
5.3.2 Indentation versus Load profiling
In order to evaluate if changing the grain boundary properties of silicon carbide impacted the deformation behavior, indentation hardness versus load profiles were developed for these materials. The analysis of the curve allows for the extrapolation of plasticity data as developed by McCauley and Wilantewicz [97]. Proper analysis requires hardness indents taken at higher values than standard hardness tests so a second machine
(Wilson-Tukon Model 300) was utilized for these procedure. Hardness indents were made at the loads of 0.100 kg, 0.300 kg, 0.500 kg, 0.700 kg, 1 kg, 2 kg, 5 kg, and 10 kg.
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The computerized system and microscope allows for careful monitoring and measurement of the indents over the wide range of sizes. After measuring a minimum of
7 indents for each load, curves of Knoop hardness versus load were made.
Using the equations outlined by McCauley below, a simple power law fit to the curve, followed by a log-log analysis yields a line, the slope of which can be used to determine a plasticity parameter.
c HK = kF (Equation 15)
Where HK is the Knoop hardness (N/m2), F is the load (N), k is a constant (N(1-c)/m2), and c is the scaling exponent which is dimensionless. Taking the log10 of both sides
gives the relationship:
log10 HK = log10 k + clog10 F (Equation 16)
A plot of the log10 HK versus log10 F will therefore return straight lines with slopes of c.
Taking the absolute value of the reciprocal of c then yields a value semi-quantitatively associated with plasticity. The addition of this value with the Knoop hardness at 1 N, as shown below, is a simple number for comparing the plasticity of different materials
plasticity = HK(1N) + [−(1/c)] (Equation 17)
5.4 Results
5.4.1 Plasticity Analysis
Hardness analysis was chosen as the primary method of comparing these
materials currently with belief that the modified grain boundaries would lead to enhanced
plastic behavior or interesting deformation mechanisms. Using the higher load Wilson-
Tukon hardness tester allowed for load versus hardness curves to be developed for these
186 samples. Following the analysis of McCauley and Wilantewicz, these curves were further analyzed to obtain plasticity data. The load-hardness curves are shown below in
Figure 5.1. The hardness curves are fit to a power law equation:
c HK = kF (Equation 18)
After taking the log of both the load and hardness values, they can be graphed, and the
slopes of the line as well as the y-intercept hardness values are recorded. The log-log plot is shown in Figure 5.2 along with a linear fit and the R squared value. The R squared of 0.97 indicates a good fit and valid relationship. The addition of the inverse slope and the y-intercept hardness value then yields a plasticity value.
Figure 5.1: Load-hardness curves for selected samples
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Figure 5.2: Log-log plot of load and hardness.
Somewhat surprisingly, all of the materials studied to date displayed very similar curves and plasticity numbers over a narrow range from between 35.0 and 36.3. It was
expected that the very different boundary characteristics and compositions would lead to
some changes in plasticity, but the results indicate fairly similar behavior from 1 to 5
weight percent of additives. The sample prepared with the least additive, 1 weight
percent, does show the highest hardness values in the majority of the curve, but the
values are fairly similar to other samples as well. Comparing SPS material to hot pressed
material of the same composition reveals that the SPS sample displays higher hardness at
loads above 1 kg. This does not appear to be indicative of a grain size reduction through
SPS processing at this time, as the average grain size for all samples studied is similar,
approximately 1.1 to 1.2 microns.
The full set of samples analyzed through the plasticity analysis method is shown
below in Figure 5.3. It is interesting to note that the values are similar between 2 and 5
wt. % additive systems and relatively insensitive to the rare earth. The plasticity number
is calculated from the addition of the y-intercept hardness and the inverse of the slope,
188 which is believed to be quantitatively related to plasticity [97]. The Knoop hardness values are generally higher for the 2 wt. % system because of the lower glassy phase content. Since the 5 wt. % systems appear to have very similar total plasticity values, it follows that they may have a greater contribution from plasticity. Increased amounts of liquid phase tend to decrease the grain size which could account for some of the difference although the hardness numbers at such low loads do not seem to show an increase from grain size. It is also possible that the contribution is legitimately from a difference in the strength, adherence, or thickness of the grain boundaries.
Among the 5 weight percent samples, there is an indication that the improved processing through coprecipitation, as well as sintering through SPS, is beneficial for plasticity. Figure 5.4 shows the 5 weight percent samples from the plasticity study and it can be observed that the coprecipitated and spark plasma sintered samples have plasticity values between 35 and 37. Ball milled samples that were spark plasma sintered had a plasticity value at 33.5. A coprecipitated sample that was hot pressed had a value of 33.2.
Again, it is possible that the grain size advantages developed from coprecipitation and spark plasma sintering are responsible for some of the contribution to these higher values.
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Figure 5.3: Plasticity values for a variety of samples
Figure 5.4: Plasticity values for 5 weight percent additive systems
5.4.2 Indent Analysis
After the plasticity results appeared so similar, it was decided that examining the individual indents in these samples may reveal information into how the materials are behaving under indentation and if there is some similar response that can explain the similar values. SEM images were taken from a number of the samples, focused on
190 looking at and into the indents to evaluate cracking and damage. A high density hot pressed SiC composition, SiC-N (BAE Systems, Arlington, VA), was chosen as a basis for comparing indents.
Figure 5.5 shows 2 kg. indents in SiC-N and in a 2 wt. % additive coprecipitated sample. There are considerable differences between the two indents. SiC-N, which has a grain size of 2 - 4 microns displays large regions of grain pullout and damage inside the indent. The finer grained, boundary modified sample has a clean, well defined indent with less grain pop out. SiC-N also shows considerable cracking emanating from the indent tips at this load while R62 shows no cracking at all in this case. It is possible that there is more subsurface damage in R62 which is preventing the cracking at the surface.
Figure 5.5: Comparing 2 kg indents in SiC-N (left) and coprecipitated SiC R62 (right)
191
Figure 5.6 shows a 10 kg. indent from R31, a 5 wt. % sample made using Al and
Y additives and hot pressed. The indent, even considering the very high load applied, appears to be fairly intact with intergranular fracture between many of the grains inside the indent, but with little grain pullout or spall. At this load, there is cracking observed from the indent tips and other regions, but the crack paths are long and intergranular with little damage otherwise observed. In contrast to this, Figure 5.7 below is the same composition as R31 but is instead densified by SPS. The density is very similar between the two samples, around 97.5 percent as calculated, but the indents look very different.
Where R31 had clean indents and little spalling, R52 showed large amounts of spall at nearly every indent. Since these samples are the same composition, it seems reasonable to state that the SPS has a very strong influence on the strength, cracking, or stress state present in the grain boundaries. Interestingly, the hardness and plasticity numbers are only slightly different however.
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Figure 5.6: 10 kg. indents in coprecipitated hot pressed sample R31 – 5 wt. % additives Al2O3-Y2O3
193
Figure 5.7: Indents in coprecipitated spark plasma sintered sample R52 – 5 wt. % additives Al2O3-Y2O3
Sample R41, Figure 5.8, with lanthanum instead of yttrium at 5 weight percent
additives displays some small areas of spalling, but also some clean indents. This
behavior appears intermediate to the extreme spalling displayed by the yttrium based
sample.
194
R56 was processed through the SPS, and was the lowest additive content sample made to date at 1 weight percent of Al,Y. As seen in Figure 5.9, the indents in this sample are very clean with limited intergranular fracture inside the indents and only small cracking at the indent tips. The appearance of pullout is very limited and very small, even compared to some of the other similar samples. This is a very interesting sample because of the behaviors shown at such low additive contents where microscopy on etched samples indicated little phase present at the grain boundaries.
Figure 5.8: 10 kg. indents in coprecipitated SPS R41 – 5 wt. % additives Al2O3-La2O3
195
Figure 5.9: Indents in coprecipitated SPS R56 – 1 wt. % additives, Al2O3-Y2O3
196
Sample R63, below as Figure 5.10, contains 2 weight percent additives based on aluminum and gadolinium. The indents here were generally clean and similar to those in
R56. There was more cracking inside the indents however and at the crack tips.
Interestingly, one of the indents showed considerable spalling which appeared to occur from a section of large grains observed right at the center of the indent. These regions of large grains, typically 5 microns or so, are observed in some SPS samples and may be due to local hot spots or another phenomenon caused by the heating in the SPS. These features are clearly severe enough to act as a defect as in this case, and would be an undesired result in an otherwise well-mixed material.
Figure 5.10: Indents in coprecipitated spark plasma sintered R63 – 2 weight percent additives, Al2O3-Gd2O3
Further analysis of some of the indents reveals some other interesting
observations. Figure 5.11 below is from R62, one of the samples from the fluorescence
197 study, and shows the surfaces inside the indent. Upon inspection it appears many of the grains are cracked intergranularly as expected for a grain boundary containing SiC material. A closer look reveals small cracks within many of the grains, roughly parallel to the central axis of the indent. This may imply some sort of plasticity mechanism or mixed fracture occurring in this sample. This could also be related to the grain boundary strength and imply that intergranular fracture only occurs along certain orientations of boundaries.
Figure 5.11: 2 kg. indent in coprecipitated SPS R62 – 2 wt. % additives, Al2O3-Sm2O3
5.5 Summary
Initial evaluations of Knoop hardness demonstrated little differences in values
between coprecipitated and ball milled samples or different rare earths or additive
amounts at the standard indentation load of 1 kg. A more comprehensive study of high
load hardness values and their extrapolation to determine plasticity values revealed some
trends. Coprecipitated samples that were densified through SPS showed higher plasticity
198 values versus either coprecipitated and hot pressed samples or ball milled samples that were spark plasma sintered. While this may be indicative of the importance of microstructure and grain boundary phase crystallinity, the values were generally over a very tight range, making any strong conclusions difficult with this type of measurement.
Actual microscopic observations of the high load hardness indents revealed considerable differences in the amount of spalling damage and cracking in certain samples. When compared to a commercial hot pressed SiC material used for structural applications, the behaviors of coprecipitated samples indicated greater influence of the grain boundaries in limiting spalling and providing a different mechanical response.
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Chapter 6 - Conclusions and Future Work
6.1 Conclusions:
Coprecipitation was developed as a processing route for making fine grained, liquid phase sintered silicon carbide. Non-aqueous coprecipitation processing was shown to be a successful method for coating sintering aids of alumina and yttrium, gadolinium, samarium, and lanthanum oxides onto silicon carbide over a wide range of total additive contents. Highly dense samples were prepared through both hot pressing and spark plasma sintering of coprecipitated powders.
The success of the coprecipitation process was shown to be highly dependent upon the powder preparation and dispersion steps. Milling treatments were shown to improve the particle size distribution and dispersability of the SiC powder before coprecipitation. Non-aqueous coprecipitation using isopropanol was developed as a preferable processing route because of the superior dispersability of SiC in isopropanol versus aqueous systems. The use of these methodologies was successful in yielding coprecipitated samples with a reduced population of defects from poor distribution of sintering aids and agglomerated powders.
Hot pressing was effective in densifying aluminum and yttrium containing systems at both 2 and 5 wt% of sintering aids, and developed microstructures with finer grain sizes than ball milled systems. Improved densification was observed for the other rare earth systems as well, but consistently high densities were not always achieved while hot pressing. Spark plasma sintering was shown to be effective for all compositions and rare earths studied, although slightly modified sintering cycles were required for the
200 various compositions. The short cycle times further reduced the grain size and modified the crystallinity of the grain boundary phase.
Comparisons between coprecipitated samples and standard ball mill and aqueous mixing processes indicated noticeable differences in grain size and crystallinity between the differently prepared samples. Additionally, the sintering method employed and temperature also seemed to exhibit a strong influence over the final properties of the systems. At the highest level of sintering aids studied, 10 wt% additives in the aluminum/yttrium system, x-ray diffraction of coprecipitated samples indicated the presence of a crystalline 2:1 mullite phase while examination of the ball milled samples revealed measurable amounts of a yttrium disilicate phase as well as excess alumina. The phase differences could indicate an improvement in the additive mixing in the coprecipitated systems, as well as possible changes in the silica levels between the different processing routes.
Fluorescence lifetime measurements were utilized to compare the level of mixedness between coprecipitated and standard processed samples revealing a quantifiable improvement in the mixing of additives from coprecipitation. Coprecipitated samples displayed higher fluorescence intensities and longer decay lifetimes than ball milled samples because of reduced concentration quenching. This again implies an improvement in mixedness when utilizing this coprecipitation procedure.
The mechanical properties of hardness, elastic modulus, and plasticity values were measured. Although the x-ray diffraction and fluorescence work indicated noticeable differences between coprecipitated and standard milled samples, as well as other differences between hot pressed and spark plasma sintered samples, analysis of the
201 hardness and elastic modulus data revealed only small and subtle changes. The majority of the samples displayed elastic moduli values of 380 to 400 GPa which appeared low considering the high densities of these samples. Upon further analysis, these values were determined to be due to the presence of silicate based glassy phases. The application of pressure during the hot pressing and spark plasma sintering processes as well as the coprecipitation process itself may have been influencing the silica content and resulted in the formation of silicate phases. The XRD phase identifications and quantifications were used as a basis for estimating what the actual densities of these systems, as well as what the expected elastic moduli values would be from the rule of mixtures. Estimations performed in this manner were shown to be reasonably accurate in determining the elastic modulus values, and were therefore offered as a justification of the density, porosity, and phase identifications for these systems.
The plasticity values obtained from instrumented load hardness measurements were also more uniform than expected, although the combination of coprecipitation and spark plasma sintering did yield some samples with small improvements in plasticity versus standard processing at the same additive level. Although the hardness values were typically higher with reduced additive content, the calculated plasticity values appeared to show some slight improvement with an increase from 2 to 5 wt. % additives.
Examination of the hardness indents through SEM revealed more limited surface damage and reduced spalling for some compositions, while others appeared to show interesting spall patterns. When compared to a high density structural SiC material with larger grains and less boundary phase, many of the SiC samples developed displayed less surface damage and spalling, and shorter cracks emanating from the indent tips.
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6.2 Future Work:
This thesis served to develop coprecipitation as a method for preparing liquid phase sintered silicon carbide with improved mixedness and better sinterability.
Although a large number of samples were developed throughout the process of developing the method, a number of variables in the coprecipitation procedure, as well as other compositional possibilities, were not explored as thoroughly as they could be. The effects of supersaturation, temperature control, and rate of the coprecipitation additions could all be exploited to gain further control over the size and quality of the sintering aid coating. In order to limit the number of samples for comparison, the molar ratio of aluminum to the rare earth element was fixed at 4 to 1 for all systems. This was chosen from literature and was assumed to provide good mechanical properties, but some preliminary work on other ratios indicated the method is applicable and effective at other ratios. Exact eutectic compositions or other phase fields are certainly worth exploring for their effect on the properties and phases in these systems. It would also be relatively simple to expand the compositions to include transition metals, mixes of rare earths, or other desirable elements.
The understanding on the role of grain boundary glassy and crystalline phases on the mechanical properties of these systems would be improved with better identification and quantification of the phases present. An expansion of the glass quantification study in section 3.3.2.1.4 where SiC was doped with glass frits in order to show the change in slope in the amorphous hump of the XRD pattern would be useful in the future.
Choosing a glass frit more exactly resembling the expected glass phase as well as varying
203 the amounts and using high resolution equipment might allow for a determination of the glassy content with a higher degree of confidence.
Difficulty in preparing suitable TEM samples because of the nature of the grain boundaries limited the TEM analysis of sintered samples to only a couple of samples, but further emphasis on electron energy loss spectroscopy or high resolution x-ray energy dispersive spectroscopy should be investigated as a means to probe the structure of the grain boundaries and triple points, as well as reveal valuable information on the location and state of the sintering aids. Better chemistry, additive distribution, and phase crystallinity information could be determined utilizing high resolution TEM, EDS, and
EELS.
Along these same lines, a more expansive TEM study of SiC surfaces and the adherence of sintering aid coatings to unfired powders would be useful in driving future work towards different and more effective coating systems. TEM analysis showing the measurement of coating thickness and evaluation of the uniformity could further help enhance the coating quality.
The use of fluorescence decay analysis in interrogating the mixing of sintering aids should be confirmed and expanded as a powerful technique. The design of sintering aid system to optimize their fluorescence for better analysis as well as confirmation of these measurements with confocal or other microscopy techniques would be an excellent next step in analyzing these samples.
The role of the silica surface on the SiC powders was mentioned throughout this thesis, and was used to an advantage in developing low temperature sinterable compositions. The amount of the silica clearly impacts the phase formation and
204 mechanical properties however, and should be explored as another controllable variable in these systems. It is also important to find a methodology for tracking the changes in silica content that occur during an aqueous or non-aqueous process.
Plastictity through instrumented load hardness analysis was chosen as a simple mechanical analysis method in this research because it was believed that it may have been more accurate and sensitive to grain boundary changes than standard toughness through indentation fracture toughness measurements. These plasticity measurements while simple to perform are empirical however, and their physical meaning is unclear.
Better measurements of plasticity through instrumented stress strain analysis or other controlled indentation and analysis methods would be beneficial. Sample size was also limited in much of this work which at the time discouraged strength measurements. It would be beneficial however, to further analyze the full range of mechanical properties for these systems, including toughness and strength, in order to better determine if there are any significant differences or advantages to coprecipitation. A definitive test relevant to the grain boundary strength would also be beneficial in future work.
The high price of rare earth elements limits the industrial relevance of these compositions somewhat, other than the lowest additive systems. The information and knowledge gained about grain boundaries and properties is valuable to study at this point, but it would be advantageous to expand this work into more common systems and cheaper additives.
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Curriculum Vita
Steve Mercurio
Education:
2011 PhD – Materials Science and Engineering, Rutgers University
2010 MS – Materials Science and Engineering, Rutgers University
2005 BS – Materials Science and Engineering, University of Pittsburgh
Work Experience:
2011 – pres Senior R&D Engineer, Materials, American Standard Brands, Piscataway, NJ
2009 – 2011 Materials Scientist II (Contractor), US Army Research Laboratory and Data Matrix Solutions, Inc., Aberdeen, MD
2005 – 2009 Graduate Research Assistant, Dept. of Materials Science and Engineering, Rutgers University
2004 – 2005 Materials Co-Op / Intern, Extrude Hone, Prometal Division, Irwin, PA
2003 Co-Op/ Intern, Process Technology and Product Development, AK Steel Butler Works, Butler, PA
Publications:
1) S.R. Mercurio, S. Miller, J. Michael, F. Cosandey, R. Haber. Improved Additive Mixing via Modified Coprecipitation Processing. (Work in Progress, 2011)
2) S.R. Mercurio, M. Jitianu, R.A. Haber. Grain Boundary Engineering of Silicon Carbide by Means of Coprecipitation. Advances in Ceramic Armor IV: Ceramic Engineering and Science Proceedings, Volume 29, Issue 6, 2009, Pages 141-152.
3) S.R. Mercurio, R.A. Haber. SiC Microstructure Improvements for Armor Applications. Advances in Ceramic Armor III: Ceramic Engineering and Science Proceedings, Volume 28, Issue 5