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Grain Boundary and Triple Junction Chemistry of Silicon Carbide Sintered with Minimum Additives for Armor Applications

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PAGE 1

GRAIN BOUNDARY AND TRIPLE JUNC TION CHEMISTRY OF SILICON CARBIDE SINTERED WITH MINIMUM ADDITIVES FOR ARMOR APPLICATIONS By EDGARDO L. PABIT A DISSERTATION PRESENTED TO THE GRADUATE SCHOOL OF THE UNIVERSITY OF FLOR IDA IN PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR THE DEGREE OF DOCTOR OF PHILOSOPHY UNIVERSITY OF FLORIDA 2005

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Copyright 2005 by Edgardo L. Pabit

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This document is dedicated to my loving wife.

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ACKNOWLEDGMENTS I would like to thank my adviser, Dr. Darryl Butt, for his guidance, patience and support during the course of my studies. I would also like to thank Dr. John Mecholsky Jr., Dr. Paul Holloway, Dr. Wolfgang Sigmund and Dr. Mark Orazem, my committee members, for taking the time to answer my questions and for the helpful comments on my dissertation. I would like to acknowledge Ceramatec, Inc. for the funding of this project and for providing the mechanical property data for this dissertation. I would also like to acknowledge Kerry Seibien of MAIC at the University of Florida and Dr. Helge Heinrich of MCF at the University of Central Florida. Their help on performing the HRTEM and EFTEM studies presented in this dissertation is greatly appreciated. Finally, I would like to thank Dr. Butts group members for their support and for providing a semblance of sanity during the long hours in the laboratory. iv

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TABLE OF CONTENTS page ACKNOWLEDGMENTS.................................................................................................iv LIST OF TABLES...........................................................................................................viii LIST OF FIGURES...........................................................................................................ix ABSTRACT.......................................................................................................................xv CHAPTER 1 INTRODUCTION........................................................................................................1 2 LIQUID PHASE SINTERING AND MECHANICAL PROPERTY MODIFICATIONS IN SILICON CARBIDE..............................................................6 Sintering Additives used in Densification of SiC.........................................................7 Effects of Processing Parameters in LPS-SiC............................................................10 Effect of Sintering Time and Temperature..........................................................10 Effect of Choice of Additives..............................................................................12 Effect of the Starting SiC Powders......................................................................15 Effect of Sintering Atmosphere...........................................................................17 Effect of Annealing Conditions...........................................................................19 Microstructure and Mechanical Property Modification in Sintered SiC....................20 Fracture Toughness Enhancement in SiC............................................................21 Crack Healing as a Tool for Mechanical Property Improvement in SiC............26 3 STATEMENT OF THE OBJECTIVE.......................................................................31 4 MATERIALS AND METHODS...............................................................................33 Materials.....................................................................................................................33 Characterizations........................................................................................................36 Mechanical Properties and SiC Polytypes...........................................................36 Microstructure.....................................................................................................37 Chemical Analysis...............................................................................................37 Transmission Electron Microscopy.....................................................................38 TEM sample preparation..............................................................................39 High resolution TEM (HRTEM)..................................................................40 v

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Energy dispersive spectroscopy (EDS)........................................................41 Energy filtered transmission electron microscopy (EFTEM)......................43 5 RESULTS OF MECHANICAL PROPERTY, GRAIN BOUNDARY, AND TRIPLE JUNCTION CHARACTERIZATION.........................................................46 -SiC Sintered with 0.6 wt.% B, 2 wt.% C and 0 4 wt.% Al (ABC-SiC)...............46 Processing and Polytypes....................................................................................46 Microstructure and Mechanical Properties..........................................................48 Grain Boundary, Triple Junction and Other Secondary Phases..........................51 Grain boundary width and crystallinity........................................................51 Grain boundary, triple junction and secondary phase composition.............53 -SiC 0.6B -2C...........................................................................................54 -SiC 0.5Al 0.6B 2C...........................................................................55 -SiC-1Al-0.6B-2C......................................................................................59 -SiC 1.5Al 0.6B 2C............................................................................61 -SiC 4Al 0.6B 2C...............................................................................65 -SiC Sintered with Al Additive................................................................................68 Processing and Polytypes....................................................................................68 Microstructure and Mechanical Properties..........................................................70 Grain Boundaries, Triple Junctions and Secondary Phases................................76 -SiC 1.65 wt % Al hot pressed at 1900C..............................................76 -SiC 1.65 wt % Al hot pressed at 2100C..............................................80 -SiC 1.65 wt % Al hot pressed at 2200C...............................................83 -SiC 3.3 wt % Al hot pressed at 2000C................................................85 -SiC 3.3 wt % Al hot pressed at 2200C................................................90 -SiC 1.65 wt % Al 0.5 wt % B4C 2 wt % C...................................93 -SiC Sintered with AlN Additive..............................................................................96 Processing and Polytypes....................................................................................96 Microstructure and Mechanical Properties..........................................................98 Grain Boundaries, Triple Junctions and Secondary Phases..............................101 -SiC 2.5 wt % AlN hot pressed at 2100C...........................................101 -SiC 2.5 wt % AlN hot pressed at 2200C...........................................104 -SiC 5 wt % AlN hot pressed at 2000C..............................................107 -SiC 5 wt % AlN hot pressed at 2200C..............................................109 -SiC 2.5 wt % AlN 0.5 wt %B4C.....................................................111 -SiC 2.5 wt % AlN 0.5 wt % B4C 2 wt % C................................115 SiC-N, -SiC-3.1Al2O3, -SiC-0.5B4C-2C..............................................................117 Processing and Polytypes..................................................................................118 Microstructure and Mechanical Properties........................................................119 Grain Boundaries, Triple Junctions and Secondary Phases..............................121 SiC-N..........................................................................................................121 -SiC 3.1 wt % Al2O3.............................................................................124 -SiC 0.5 wt % B4C 2 wt % C...........................................................126 vi

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6 SUMMARY AND DISCUSSION: MECHANICAL PROPERTY, GRAIN BOUNDARY, AND TRIPLE JUNCTION CHARACTERISTICS........................129 -SiC Sintered with 0.6 wt.% B, 2 wt.% C and 0 4 wt.% Al (ABC-SiC).............129 Influence of Residual Stress..............................................................................130 Absence of Intergranular Films in Materials with Al Additive in Excess of Solubility Limit..............................................................................................133 SiC Ceramics Fabricated from -SiC Starting Powders...........................................135 Comparison of Mechanical Properties and Grain Boundary and Triple Junction Characteristics of SiC Sintered with Equimolar Al Addition.........136 Comparison of SiC Ceramics with Equimolar Al Addition..............................136 Effect of B and C Additions..............................................................................140 Comparison of -SiC Ceramics Sintered with Al or AlN........................................144 Effect of Hot Pressing Temperature and Amount of Additives........................144 Achievement of SiC-N-like Properties..............................................................147 7 CRACK HEALING OF SELECT SPECIMEN.......................................................151 Materials and Methods.............................................................................................152 Results and Discussion.............................................................................................155 Conclusions...............................................................................................................160 8 CONCLUSION.........................................................................................................161 LIST OF REFERENCES.................................................................................................164 BIOGRAPHICAL SKETCH...........................................................................................173 vii

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LIST OF TABLES Table page 4-1 Compositions and processing conditions of the materials investigated...................34 4-2 Raw materials used in powder processing...............................................................35 5-1 Density and polytypes of SiC ceramics from -SiC starting powders.....................47 5-2 Grain size, aspect ratio, hardness, toughness and observed fracture mode in SiC ceramics with -SiC starting powders......................................................................49 5-3 Grain boundary width and crystallinity of -SiC with Al, B, and C additions........53 5-4 Density, chemical analysis and phase assemblage of -SiC sintered with Al additive.....................................................................................................................68 5-5 Mechanical properties of -SiC sintered with Al additive.......................................73 5-6 Density, chemical analysis and phase assemblage of -SiC sintered with AlN additive.....................................................................................................................97 5-7 Mechanical properties of -SiC sintered with AlN additive..................................100 5-8 Density, chemical analysis and phase assemblage of SiC-N, -SiC-Al2O3 and -SiC-0.5B4C-2C.......................................................................................................119 5-9 Hardness, toughness, and mode of fracture observed in SiC-N, -SiC-3.1Al2O3 and -SiC-0.5B4C-2C.............................................................................................120 6-1 Summary of observations for SiC ceramics sintered from -SiC starting powders..................................................................................................................130 6-2 Mechanical properties, grain boundary and triple junction characteristics of SiC materials with equimolar aluminum addition.........................................................136 6-3 Mechanical properties, grain boundary and triple junction characteristics of SiC sintered with Al or AlN addition............................................................................145 7-1 Composition, processing conditions and mechanical properties of materials used for crack healing.....................................................................................................152 viii

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LIST OF FIGURES Figure page 4-1 Signals generated from high-energy electron beam/thin specimen interaction.....38 4-2 Schematic overview of the three-window technique............................................44 5-1 Transmission electron micrographs of the microstructure observed in SiC ceramics fabricated from a -SiC starting powder................................................48 5-2 High resolution transmission electron micrographs of grain boundaries in -SiC with Al, B and C additives ...........................................................................52 5-3 STEM of a grain boundary and a secondary phase in -SiC-0.6B-2C..................55 5-4 STEM of a triple junction in -SiC-0.6B-2C ......................................................56 5-5 EFTEM of a triple junction and grain boundary in -SiC-0.6B-2C ...................56 5-6 STEM of a triple junction in -SiC-0.5Al-0.6B-2C ...........................................57 5-7 STEM of a grain boundary in -SiC-0.5Al-0.6B-2C ..........................................57 5-8 EFTEM of a grain boundary in -SiC-0.5Al-0.6B-2C..........................................58 5-9 STEM of a multiple grain junction in -SiC-0.5Al-0.6B-2C containing metallic impurities as shown by the corresponding EDS......................................58 5-10 STEM of a triple junction in -SiC-1Al-0.6B-2C.................................................60 5-11 EFTEM of a triple junction in -SiC-1Al-0.6B-2C ............................................60 5-12 EFTEM of a grain boundary in -SiC-1Al-0.6B-2C.............................................61 5-13 STEM of a triple junction in -SiC-1.5Al-0.6B-2C ...........................................62 5-14 EFTEM of a triple junction in -SiC-1.5Al-0.6B-2C ........................................63 5-15 STEM of a grain boundary in -SiC-1.5Al-0.6B-2C.............................................63 5-16 EFTEM of a grain boundary in -SiC-1.5Al-0.6B-2C..........................................64 ix

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5-17 STEM of a grain boundary and associated bulk grains in -SiC-1.5Al-0.6B-2C...........................................................................................................................65 5-18 Residual carbon in triple junction of -SiC-4Al-0.6B-2C.....................................66 5-19 EFTEM of a grain boundary in -SiC-4Al-0.6B-2C.............................................66 5-20 EFTEM of a multiple grain junction in -SiC-4Al-0.6B-2C.................................67 5-21 STEM and EDS of secondary phases found in -SiC-4Al-0.6B-2C.....................67 5-22 Transmission electron micrographs of the microstructures observed in -SiC sintered with Al additive........................................................................................71 5-23 Optical micrographs of Vickers indentation in -SiC sintered with aluminum additions.................................................................................................................73 5-24 SEM of fracture surfaces of -SiC sintered with aluminum additions..................74 5-25 High resolution transmission electron micrographs of several grain boundaries studied in -SiC-1.65Al hot pressed at 1900C.....................................................76 5-26 STEM and EDS of a grain boundary and associated grains in -SiC-1.65Al hot-pressed at 1900C............................................................................................77 5-27 STEM and EDS of a triple junction in -SiC-1.65Al hot-pressed at 1900C.......78 5-28 EFTEM of a triple junction and grain boundary in -SiC-1.65Al hot pressed at 1900C...................................................................................................................78 5-29 EELS taken from a triple junction in -SiC-1.65Al hot pressed at 1900C..........79 5-30 High resolution images and EDS of triple junctions in -SiC-1.65Al hot pressed at 1900C..................................................................................................79 5-31 HRTEM of several grain boundaries in -SiC-1.65Al hot pressed at 2100C......80 5-32 EFTEM of a grain boundary in -SiC-1.65Al hot pressed at 2100C...................81 5-33 STEM and EDS of triple junction in -SiC-1.65Al hot pressed at 2100C..........82 5-34 STEM of triple grain junction phase with lower dihedral angle in -SiC-1.65Al hot pressed at 2100C................................................................................82 5-35 EFTEM of a triple junction in -SiC-1.65Al hot pressed at 2100C.....................83 5-36 HRTEM of grain boundaries in -SiC-1.65Al hot pressed at 2200C..................84 x

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5-37 STEM of a triple junction in -SiC-1.65Al hot pressed at 2200C.......................84 5-38 EFTEM of a triple junction and grain boundary in -SiC-1.65Al hot pressed at 2200C...................................................................................................................85 5-39 HRTEM of grain boundaries in -SiC-3.3Al hot pressed at 2000C....................86 5-40 STEM of a grain boundary near a triple junction in -SiC-3.3Al hot pressed at 2000C...................................................................................................................86 5-41 STEM of a triple junction in -SiC-3.3Al hot pressed at 2000C.........................87 5-42 STEM of two grains with different composition found in -SiC-3.3Al hot pressed at 2000C..................................................................................................88 5-43 EFTEM of a triple junction in -SiC-3.3Al hot pressed at 2000C.......................89 5-44 HRTEM of grain boundaries in -SiC-3.3Al hot pressed at 2200C....................90 5-45 STEM of a grain boundary in -SiC-3.3Al hot pressed at 2200C.......................91 5-46 EFTEM of a grain boundary in -SiC-3.3Al hot pressed at 2200C.....................91 5-47 STEM and EDS of a triple junction containing primarily Al-rich phase and metallic impurities in -SiC-3.3Al hot pressed at 2200C....................................92 5-48 STEM and EDS of Al-rich grain found in -SiC-3.3Al hot pressed at 2200C....92 5-49 EFTEM of a triple junction in -SiC-3.3Al hot pressed at 2200C.......................93 5-50 HRTEM of grain boundaries observed in -SiC-1.65Al-0.5B4C-2C....................94 5-51 EFTEM of a grain boundary in -SiC-1.65Al-0.5B4C-2C....................................94 5-52 STEM of a triple junction in -SiC-1.65Al-0.5B4C-2C........................................95 5-53 STEM and EDS of other secondary phases in -SiC-1.65Al-0.5B4C-2C.............96 5-54 TEM of the microstructures observed in -SiC sintered with AlN additive.........99 5-55 Optical micrographs of Vickers indentation on a polished surface of -SiC sintered with AlN additive...................................................................................100 5-56 SEM of fracture surfaces of -SiC sintered with AlN additive...........................101 5-57 HRTEM of grain boundaries observed in -SiC-2.5AlN hot pressed at 2100C.................................................................................................................102 xi

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5-58 STEM and EDS taken from -SiC-2.5AlN hot pressed at 2100C.....................102 5-59 EFTEM of a triple junction in -SiC-2.5AlN hot pressed at 2100C..................103 5-60 Low and high resolution image of a triple junction in -SiC-2.5AlN hot pressed at 2100C................................................................................................104 5-61 HRTEM of grain boundaries in -SiC-2.5AlN hot pressed at 2200C...............105 5-62 STEM and EDS of grain boundary in -SiC-2.5AlN hot pressed at 2200C......105 5-63 STEM of a triple junction in -SiC-2.5AlN hot pressed at 2200C....................106 5-64 EFTEM of a triple junction in -SiC-2.5AlN hot pressed at 2200C..................106 5-65 HRTEM of grain boundaries in -SiC-5AlN hot pressed at 2000C..................107 5-66 STEM of triple grain junction in -SiC-5AlN hot pressed at 2000C.................108 5-67 EFTEM of triple junctions in -SiC-5AlN hot pressed at 2000C......................109 5-68 HRTEM of grain boundaries in -SiC-5AlN hot pressed at 2200C..................110 5-69 STEM of a triple junction in -SiC-5AlN hot pressed at 2200C.......................110 5-70 EFTEM of a multiple grain junction and triple grain junction in -SiC-5AlN hot pressed at 2200C..........................................................................................111 5-71 HRTEM of several grain boundaries observed in -SiC-2.5AlN-0.5B4C...........112 5-72 STEM of triple junction observed in -SiC-2.5AlN-0.5B4C...............................113 5-73 EFTEM of a grain boundary observed in -SiC-2.5AlN-0.5B4C........................113 5-74 EFTEM of a triple junction observed in -SiC-2.5AlN-0.5B4C..........................114 5-75 STEM of secondary phases observed in -SiC-2.5AlN-0.5B4C.........................114 5-76 HRTEM of several grain boundaries in -SiC-2.5AlN-0.5B4C-2C....................115 5-77 EFTEM of a triple junction and grain boundary in -SiC-2.5AlN-0.5B4C-2C...116 5-78 Another EFTEM taken from a different triple junction in this material..............116 5-79 STEM of several secondary phases observed in -SiC-2.5AlN-0.5B4C-2C.......117 5-80 TEM of microstructure observed in SiC-N, -SiC-0.5B4C-2C and -SiC-3.1Al2O3...............................................................................................................120 xii

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5-81 Optical micrographs of Vickers indent on a polished surface of SiC-N, -SiC-0.5B4C-2C and -SiC-3.1Al2O3...........................................................................120 5-82 SEM of fracture surfaces of SiC-N, -SiC-0.5B4C-2C and -SiC-3.1Al2O3......121 5-83 HRTEM of several grain boundaries observed in SiC-N....................................122 5-84 STEM of a grain boundary in SiC-N...................................................................122 5-85 EFTEM of a grain boundary in SiC-N.................................................................123 5-86 STEM of a triple junction in SiC-N.....................................................................123 5-87 EFTEM of a triple junction in SiC-N..................................................................124 5-88 HRTEM of several grain boundaries in -SiC-3.1Al2O3.....................................125 5-89 STEM of a triple junction in -SiC-3.1Al2O3......................................................125 5-90 EFTEM of a grain boundary and a triple junction in -SiC-3.1Al2O3................126 5-91 HRTEM of grain boundaries observed in -SiC-0.5B4C-2C..............................127 5-92 STEM of a triple junction in -SiC-0.5B4C-2C...................................................127 5-93 STEM of a triple junction containing residual carbon and metallic impurities in -SiC-0.5B4C-2C.............................................................................................128 5-94 EFTEM of triple junction and grain boundary in -SiC-0.5B4C-2C...................128 6-1 Calculated residual stresses due to the presence of secondary phase..................132 6-2 TEM of microstructure observed in -SiC sintered with equimolar Al addition................................................................................................................138 6-3 SEM of crack profiles in SiC materials with equimolar Al addition...................139 6-4 SEM of crack profiles in materials containing B4C and C..................................141 6-5 TEM of microstructure observed in -SiC sintered with further addition of B4C and C............................................................................................................142 6-6 SEM of crack profiles in -SiC sintered with Al additives.................................147 6-7 SEM micrographs of crack path in SiC-N and -SiC sintered with AlN addition................................................................................................................148 7-1 Optical micrographs of the Vickers indentation using a 98 N load.....................153 xiii

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7-2 Schematic diagram of the four-point bend testing fixture for the flexural strength measurement..........................................................................................155 7-3 Weight changes of the materials studied as a result of exposure in different atmosphere...........................................................................................................155 7-4 Fracture strengths of the SiC materials before and after crack healing ............156 7-5 SEM of fracture surfaces in materials exposed in air..........................................158 7-6 Crack radius calculated from the measured long axis diameter and short axis radius....................................................................................................................158 7-7 Fracture toughness calculated from the crack radius of the SiC materials..........159 xiv

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Abstract of Dissertation Presented to the Graduate School of the University of Florida in Partial Fulfillment of the Requirements for the Degree of Doctor of Philosophy GRAIN BOUNDARY AND TRIPLE JUNCTION CHEMISTRY OF SILICON CARBIDE SINTERED WITH MINIMUM ADDITIVES FOR ARMOR APPLICATIONS By Edgardo L. Pabit December 2005 Chair: Darryl Butt Major Department: Materials Science and Engineering Silicon carbide containing a minimum amount of additives for armor application was fabricated by hot pressing. Microstructural development, phase information, mechanical properties, and triple junction and grain boundary chemistry were described. Correlation between the characteristics of the secondary phases and mechanical properties was also presented. The grain boundary and triple junction phases were characterized using high resolution transmission electron microscopy, energy dispersive spectroscopy and energy filtered transmission electron microscopy. The grain boundary and triple junction phase characteristics in silicon carbide sintered with the addition of aluminum, boron and carbon varied with the amount of aluminum additive. Silicon carbide sintered with the addition of boron and carbon only did not form triple junction and grain boundary phase, while addition of 0.5 and 1 wt. % aluminum resulted in the formation of amorphous Al-O-rich triple junction phase. xv

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Crystalline Al2O3 formed at the triple junction pockets upon addition of 1.5 and 4 wt. % Al. Formation of amorphous intergranular films, with thickness of approximately 1nm, was also observed at these amounts of aluminum additive. An increase in toughness, from 2.7 to 6.1 MPa m1/2, accompanied the presence of this grain boundary phase, which was also coincidental with the transition from transgranular to intergranular mode of fracture. The increased toughness and the change in fracture mode from transgranular to intergranular fracture were attributed to the residual stresses in the intergranular films. The triple junction and grain boundary phase characteristics in silicon carbide sintered with Al, AlN, and Al2O3 addition did not vary with the source of Al additive. The triple junction phases were observed to be crystalline Al2O3 while the grain boundary phase composition remained primarily Al-O-rich. The intergranular film also remained amorphous with thicknesses of approximately 1 nm. Toughness value of 6.8 MPam1/2 was attained by hot pressing at 2200C and using 3.3 wt. % Al additive. The resulting microstructure in this material consisted of elongated grains in a matrix of fine equiaxed grains. The use of AlN additive resulted in a retarded the grain growth and equiaxed grain morphology. Addition of B4C and C resulted in increased grain growth. Heat treatment at 1300C for 3 hours of select SiC specimens showed that crack healing was possible. Exposure to argon, air and water vapor-containing environments resulted in strength recovery in all materials studied. Decreases in crack sizes were also observed upon crack healing. Toughening in solid state sintered SiC in water vapor-containing environment, due to the possible formation of secondary phases, was also observed. xvi

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CHAPTER 1 INTRODUCTION Liquid-phase-sintered silicon carbide (LPS-SiC) has attracted increasing interest for its ability to form an in-situ toughened material and potentially superior mechanical properties relative to solid-state-sintered SiC.1-4 Another incentive that makes LPS-SiC technologically attractive is the ability to slightly tailor the resulting mechanical properties through improved control of the microstructure of the sintered materials.3-10 Depending on the types and amount of additive used, types of starting powder, sintering atmosphere, and processing conditions, the resulting grain morphology and property of the materials varied. For instance, it is generally accepted that higher temperature and longer sintering time result in a pronounced grain growth.11-13 The use of nitrogen overpressure during sintering usually results in an equiaxed grain morphology regardless of the types of starting powder used,14 whereas use of argon results in equiaxed or elongated grain morphology depending on the starting powder.4,12 In the case of -SiC starting powder, the resulting grain morphology is generally equiaxed.6,7 For a -SiC or -SiC containing -SiC seeds, the resulting grains are generally elongated due to the -totransformation during sintering or annealing.5,7 It has been proposed that elongated microstructure usually increases the fracture toughness in SiC by crack bridging or crack deflection.15-17 The incremental increase in fracture toughness of LPS-SiC is generally promoted by the elongated grains with relatively weak intergranular grain boundary and triple junction phases.15-17 The intergranular phases and triple junction composition are 1

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2 generally controlled by the type of additives used. A fracture toughness of ~8 MPam1/2 has been reported in SiC with additions of yttrium aluminum garnet (YAG, Y3Al5O12).1 The amount of additives used in this system generally exceeds 10 wt. %. A slightly higher fracture toughness of ~9 MPam1/2 was also reported for SiC sintered with aluminum, boron and carbon (ABC-SiC) additives. In this case, higher toughness was achieved with additives of only about 5 wt. %.2 One of the applications envisioned for SiC with improved properties is in ceramic armor.18 Currently, the state of the art ceramic armor material against heavy threats is a silicon carbide densified with AlN, 19 produced by Cercom, Inc. under the trade name SiC-N. In armor applications, the structural ceramic component is desired to have both high hardness and toughness to perform its intended function. The armor is designed to defeat the projectile at the interface by increasing its dwell time.20 In ballistic tests, long dwell times are achieved for high hardness ceramics. And at the same level of hardness, higher toughness leads to an improved resistance to penetration of the projectile.18 Generally, high hardness results in low toughness and vice versa. The manufacturing challenge, therefore, is to find an optimal balance of both properties. Literature data generally show that higher fracture toughness is associated with intergranular mode of fracture in SiC ceramics. Intergranular mode of fracture is usually observed in the presence of intergranular films and triple junction phases in the sintered material. The presence of these secondary phases, however, diminishes the hardness in the final product. Investigation of the interdependence of these variables (grain boundary, triple junction chemistry, microstructure and processing conditions) would therefore be important in achieving a hard and toughened SiC.

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3 In this dissertation, several approaches were made in order to understand the interdependence of the grain boundary and triple junction chemistry, microstructure, processing conditions and the resulting mechanical properties of SiC ceramics. Several sets of materials were fabricated with minimum additives, and characterized with respect to the aforementioned variables. The ABC-SiC system has been of interest due to the high toughness available in this material. Five specimens were made and studied to assess the effect of aluminum concentration on the grain boundary and triple junction chemistry and the resulting mechanical properties. The materials were fabricated from a -SiC starting powder and hot pressed with the addition of 0.6 wt. % B, 2 wt. % C and varying amounts of Al (0 to 4 wt. %). Results obtained from these materials were then used as a basis for comparison to and development of succeeding test matrices. The effect of source of Aluminum on the resulting mechanical properties and grain boundary and triple junction characteristics was also investigated. The materials in this case were fabricated from -SiC starting powders and hot pressed with the addition of Al, AlN and Al2O3. The amount of additive used was made higher than the amount of Al in ABC-SiC which yielded higher toughness on the materials investigated. For the purpose of comparison, the amount of Al was also made to correspond (or be equimolar with) to the amount of Al added in the commercially available SiC-N. The effect of B and C on the densification and microstructure of SiC fabricated from -SiC starting powders was already known. In the case of SiC fabricated from -SiC starting powders, however, the effect of B and C addition has not been properly

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4 studied. To address this issue, incorporation of B4C and C on materials fabricated from -SiC powders sintered with Al or AlN addition was also investigated. Tailoring of the microstructure of SiC ceramics to achieve mechanical properties similar to that of state of the art ceramic armor material (SiC-N) was also attempted. -SiC with a different amounts of Al (1.65 and 3.3 wt. %) or AlN (2.5 and 5 wt.%) additions was hot-pressed at different temperatures (1900, 2100 and 2200C). Comparison of the triple junction and grain boundary phase characteristics and mechanical properties was performed and methods for possible enhancement were also presented. Improvement of fracture strength in SiC has been reported possible through crack healing. Crack healing can be accomplished via an additional heat treatment. In this dissertation, crack healing of select specimen was also presented. The heat treatments were carried out in air, argon and water-vapor containing environment. The work presented in this dissertation was performed in cooperation with Ceramatec, Inc. Ceramatec, Inc. fabricated all the specimens studied. The mechanical property data presented, including the fracture mode and grain size measurements, were based on the characterization performed by Ceramatec, Inc. Grain boundary and triple junction characterizations were performed at the University of Florida and the University of Central Florida. Crack healing experiments were performed at the University of Florida. Presentation of this dissertation will be as follows: Chapter 2 will cover the literature related to the liquid-phase sintering of SiC including discussions on the effect of starting powders, processing conditions, amount and type of additives and mechanical

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5 property modification; Chapter 3 will present the objective of this study; Chapter 4 will deal with the materials and methods used in this study; Chapter 5 will present the results of mechanical property, grain boundary and triple junction chemistry characterizations; Chapter 6 will show discussion of the results presented; Chapter 7 will deal with the results and discussion of the crack healing experiments; and finally, a general conclusion will be presented in Chapter 8.

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CHAPTER 2 LIQUID PHASE SINTERING AND MECHANICAL PROPERTY MODIFICATIONS IN SILICON CARBIDE Fabrication of SiC ceramics through liquid phase sintering made possible the use of SiC in demanding structural applications, such as armor component and high temperature-heat exchanger materials.19, 22-27 These two applications require different sets of material properties. In the case of armor component applications, researchers are in agreement that the ceramic material should possess relatively high hardness and toughness,18-20, 26-27 although these two properties in and of themselves are only part of the desirable attributes. High hardness and toughness are achieved in part through the judicious selection of fabrication parameters and use of liquid phase forming additives. For high temperature applications, however, the high-temperature strength retention of the ceramics should be increased. High temperature strength retention can be improved through the incorporation of minimal secondary phase.22-25 Thus, the variable that gave high toughness for ceramic armor becomes undesirable for SiC ceramics intended for high temperature applications. This phenomenon clearly shows the need for understanding the effect of processing variables on the microstructure that eventually controls the mechanical properties desired for a given application. To address this issue, discussion of the different additive system used for the densification of SiC ceramics will be presented in this section. In addition, the effects of several sintering parameters on the densification and resulting microstructure, and efforts 6

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7 reported in the literature for improvement of the mechanical properties of SiC ceramics will also be presented. Sintering Additives used in Densification of SiC The fabrication of dense SiC ceramics requires a sintering additive because of the extremely low self diffusivities in the strongly covalently bonded and -SiC structures.28 An effective additive promotes densification and creates an environment which inhibits the decomposition of the SiC powders at the high sintering temperature.13 In liquid phase sintering, the additive should be able to participate in liquid phase formation and the resulting liquid phase should act as a mass transport media during densification. To perform this function, the liquid phase generated should be of sufficient volume to allow complete wetting of the solid phase. Further, it can be desirable for the solid phase to have an appreciable solubility in the liquid to promote solution-reprecipitation.13,29 These characteristics are dependent on the particular additives, the relevant eutectic temperatures and the densification parameters such as sintering atmosphere and temperature regime. The resulting densification rates, grain morphology and microstructures, and mechanical properties of the sintered material are thus controlled by the appropriate selection of fabrication parameters. In an attempt to obtain full densification and remove processing flaws that practically limit the strength and toughness of SiC ceramics, several additive systems have been investigated in the literature.30-58 In most cases, the additive system is based on aluminum, added as metallic aluminum, aluminum oxide or aluminum nitride, in combination with other elements, such as B and C, or metallic oxides, such as Y2O3 and rare-earth oxides. Aluminum has been primarily used since it has been observed to facilitate liquid phase sintering in SiC. The presence of aluminum lowers the

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8 densification temperature, induces the -tophase transformations, promotes anisotropic grain growth of SiC grains, forms amorphous grain boundary films, and produces various secondary phases in triple junction pockets.53,54,58 The use of aluminum, in combination with boron and carbon, was extensively studied by several investigators from the University of California, Berkeley.2, 53-58 Their work has contributed to what has become known as the aluminum-boron-carbon containing SiC (ABC-SiC) system. The starting SiC powder in these studies is generally of the -SiC polytype and utilizes the -tophase transformation to toughen the resulting monoliths. Toughness values as high as 9.1 MPam1/2 and a 4-point bend strength of around 660 MPa have been achieved.2 This toughness value is the highest reported for SiC so far and was achieved at a relatively low sintering temperature of 1900C using a minimum amount of additives (3 wt.% Al, 0.6 wt.% B and 2 wt.% C). Other liquid phase-forming additives used to densify SiC ceramics reported in the literature include Al2O3, Al2O3-Y2O3, AlN-Y2O3, Al2O3-Y2O3-CaO, Al2O3-rare-earth oxides and Al2O3-lanthanum series oxides.30-52 Al2O3 is believe to react with SiO2, which was always present on the surface of SiC particles, to form a liquid melt and, with increasing temperature, an oxycarbide melt due to dissolution of SiC.13 The use of Al2O3-Y2O3 system was because this compound forms a eutectic melt at a lower temperature.5, 31-33 The resulting microstructure from this additive system usually yielded a yttrium-aluminum-garnet (YAG)-containing grain boundary phase, which was believed to enhance the toughness values of the SiC ceramics. A toughness value of about 8 MPam1/2 was reported with this additive system.5 The difficulty, however, of using either Al2O3-Y2O3 or Al2O3 additives is controlling the partial reduction of Al2O3 at the

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9 sintering temperatures usually used. Al2O3 is reduced at the sintering temperatures near 2000C via the reaction, Al2O3 + SiC Al2O(g) + SiO(g) + CO(g) (1) Below 2000C, yttria alone does not show appreciable gas phase reaction with SiC, nor does it form sufficient melt phase to allow complete densification.23 Thus, the use of SiO2-rich powder bed to prevent the escape of reaction products during sintering has been employed. This complicated process is not desirable for an industrial production and usually results in low mechanical property-reproduction in resulting monoliths. A better alternative was the use of AlN-Y2O3 additive, where the presence of aluminum was maintained.23 The decomposition reaction of AlN, given by AlN 2Al(l) + N2(g) (2) can be effectively suppressed by a nitrogen overpressure during sintering. An added advantage, realized in the later studies,7,36,40 was the formation of solid-solution between AlN and SiC which produces a wurtzite 2H-SiC structure that effectively impinges on the anisotropically growing hexagonal (6H) grains during sintering.8 The 2H-SiC structure does not form a solid solution with 6H-SiC polytype.8 This phenomenon offers a way of controlling the observed anisotropic grain growth and coarsening in SiC ceramics which results in strength reduction of the manufactured ceramics. Addition of CaO in Al2O3-Y2O3-containing SiC ceramics has been used to lower the vapor pressure of product gases and to further lower the sintering temperature.3,50,59 Addition of CaO further decreases the grain growth rate and reduces the mass loss during processing. The use of rare earth and lanthanum series oxides was driven by the similarity of chemical and physical properties of these oxides with Y2O3.42 The

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10 differences in the cationic field strength (z/r2; valence (z), bond length (r)), however, resulted in the differences in properties of the grain boundary phases formed from the different oxides. It has been reported that a decrease in the cationic radius of the rare-earth oxides was accompanied by an increase in Youngs modulus, hardness, and flexural strength of the SiC ceramics, whereas the fracture toughness was improved by incorporating rare-earth oxides of larger cationic radius.42 In the case of lanthanum series oxides, specifically Lu2O3, formation of highly refractory crystalline rare earth disilicates in the grain boundary and intergranular phase leads to an improvement in the high temperature properties of the SiC ceramics produced.25 Effects of Processing Parameters in LPS-SiC For a SiC ceramic to be useful in its desired application, be it as a high-temperature structural material or as a ceramic armor component, it should be able to exhibit relatively high hardness and toughness. Like traditional ceramics, SiC exhibits a hardness/toughness trade-off. Thus, judicious selection of fabrication parameters for this material should be made in order to have mechanical properties suitable for its intended applications. In SiC, microstructure and grain boundary and triple junction phases control the mechanical properties. Knowledge of the fabrication parameters that controls these variables is therefore important. Effect of Sintering Time and Temperature The range of sintering temperature, as reported in the literature, varies from 1750C to 2200C.30-58 The choice of sintering temperature generally depends on the melting temperature of the additives used, in the case of single element additive, and the temperatures of the eutectic melts, in multi-element additive system. For liquid-phase

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11 sintered SiC, the need for enough liquid-phase volume to allow adequate wetting of the solid during densification becomes the primary consideration in sintering temperature selection.13 As an example, SiC sintered with Al2O3 and rare-earth oxides is usually sintered at 1800C. 24 This temperature is higher than the range of eutectic temperatures, 1720 to 1780C, of these additives. Another factor that determines the range of temperature used in sintering of SiC is the formation of a desired phase useful in the toughening of SiC ceramics, although phase formation or transformation also depends on the type of additive used. The -to--SiC phase transformation responsible for most of the toughening observed in SiC ceramics occurs at 1950C.2 The temperature range for the stability of -SiC, however, can be extended through the use of appropriate additive, i.e., AlN or oxynitride-phase forming additives.7,23 Formation of 2H-SiC from -SiC with AlN additive that controls the anisotropic grain growth of 6H-SiC polytype, depending on the amount of AlN additive, occurs at 1600C. Thus, sintering temperature choices are usually higher than the above-mentioned temperatures. The effect of sintering temperature on the microstructure of the resulting ceramics is manifested in the grain growth and morphology. Higher temperature generally leads to bigger grains and, for -SiC starting powders, an elongated grain morphology. It has to be noted, however, that sintering time also plays an important role in determining the resulting microstructure and acts in tandem with the choice of temperature. Sintering times reported for densification of SiC ceramics vary from 30 minutes to a couple of hours. At a given temperature, sintering time controls the grain growth and the

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12 completion of the -toSiC phase transformations. Shorter sintering times generally correspond to less grain growth and extent of phase transformation. In addition to their effect on microstructure and grain morphology, sintering time and temperature also affect the density of the sintered ceramics. In the time and temperature range reported in the literature,30-58 almost all of the SiC materials were sintered to above 95% of the theoretical density. Near theoretical densities, however, are usually attained at relatively higher temperatures and longer sintering times. Attainment of fully densified SiC ceramics, however, also depends on how the sintering process was carried out. For Al2O3-Y2O3 sintered SiC ceramics, use of powder beds with similar composition to that of the materials being sintered is necessary to prevent the evolution of reaction products during sintering that contributes to porosity formation in the final material.8,23 For AlN-Y2O3 sintered SiC, an overpressure of nitrogen is necessary to prevent the evolution of nitrogen gas during processing.8,23 Effect of Choice of Additives Literature data indicate that the choice of additive used in sintering of SiC generally controls the secondary phases (grain boundary, triple junction and residual) retained in the material upon cooling. The secondary phases in the sintered ceramics affect the toughening mechanisms and mode of failure operable in this material. It is desired, in the case of ceramic armor application, that the secondary phase provides a weak interface between the matrix grains to facilitate the toughening mechanisms operable. It is also accepted that the weak interface due to the secondary phases ensures an intergranular failure mode in the SiC ceramics. 18 As mentioned in the earlier section of this chapter, the additive system used in sintering SiC includes Al2O3, Al2O3-Y2O3, AlN-Y2O3, Al2O3-Y2O3-CaO, Al2O3-rare

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13 earth oxides and Al2O3-lanthanum series oxides. Al, Al2O3 and AlN, in conjunction with the SiO2 present in the starting SiC powders, generally form Al and O enriched secondary phases in the sintered materials.5,38,60 The Al and O-enriched grain boundary and triple junction phases provide a sufficiently weakened interface for intergranular mode of failure to predominate in SiC sintered with these additives. In the case of Al2O3-Y2O3 sintered SiC, the resulting grain boundary and triple junction phases generally contains yttrium-aluminum-garnet (YAG).5,59 In most cases, the grain boundary and triple junction phase formed are generally weaker than the SiC grains, thus providing an easier path for crack formation leading to intergranular mode of failure in the material. Aside from the formation of secondary phases, several additives also affect the resulting microstructure of the sintered SiC. In the case of SiC ceramics fabricated from -SiC starting powders, some of the additive used also affect the -tophase transformation responsible for the in-situ toughening in this material.2,53-58 As shown by studies performed by the University of California, Berkeley group,2,53-58 the use of Al, B and C has corroborating effects that promote and enhance the -toSiC phase transformation. They have observed that in terms of developing phase composition, boron is more effective in promoting the -tophase transformation than carbon. Aluminum, on the other hand, retards the -tophase transformation, but promotes 6H-to-4H transformation. In terms of grain morphology, aluminum and carbon promote anisotropic grain growth resulting in a plate-like or elongated microstructure in the final ceramics. Boron, on the other hand, tends to coarsen the grains but reduces the average aspect ratio.

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14 Studies of the ABC-SiC system have also revealed that the combined roles often override the individual roles of Al, B and C during processing of this material.2,53-58 Although both boron and carbon together favor the -tophase transformation associated with grain elongation, the resulting microstructure does not necessarily lead to strongly elongated grains as boron and carbon additives have the opposite effect on anisotropic grain growth. Thus, the final grain configuration in the resulting ceramics is determined by the B/C ratio. The effect of aluminum, however, is such that if the B/C ratio favors the anisotropic grain growth, Al accelerates such growth so that the aspect ratio is further increased. It has also been observed that reduction of the Al/B and Al/C ratio results in a weakened aluminum effect on the final microstructure due to the decrease in liquid phases generated. Even at constant Al/B/C ratios, the resulting grain configurations are significantly altered by a change in the total amount of additives. Thus, in the case of ABC-SiC system, as might be true in other sintering additive systems, literature data imply that both the ratios and total amount of sintering additives affect the final grain morphology and microstructure of the sintered SiC ceramics. Another sintering additive that affects not only the retained secondary phase but the resulting microstructure is AlN.23,35-37,61-64 It has been shown that for -SiC starting powders, use of AlN as sintering additive results in the formation of 2H-SiC polytype, together with the common 4H and 6H SiC structure, in the densified ceramics.35-37 The formation of 2H structure results in equiaxed grain morphology, while 6H and 4H yield elongated grains. At appropriate temperature and amount of additive, the 3C-to-2H SiC transformation can be made to go into completion such that the resulting grains are all equiaxed. If completion of the 3C-to-2H transformation is not achieved, the resulting

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15 microstructure contains elongated and equiaxed grains. Thus, different microstructure can be achieved in this system depending on the amount of AlN and the processing temperature. Effect of the Starting SiC Powders SiC crystallizes in a zinc-blend cubic structure and a number of different polytypes with a hexagonal or rhombohedral symmetry. The cubic structure is known as the -SiC phase and is given the Ramsdell notation of 3C.65 The hexagonal and rhombohedral polytypes are collectively called the -SiC phase and the most common, in Ramsdell notation, are 2H, 4H, 6H and 15R.65 -SiC is a metastable phase and transform at higher temperature to one of the -SiC polytypes, depending on the type of additive used. -SiC transform to a 6H -SiC polytype in the presence of boron, while the 4H modification is dominant in the presence of aluminum.2 Use of aluminum and nitrogen, either in the form of AlN or Al with N2 as sintering atmosphere, generally stabilizes the 2H polytypes.35-37 The rate of transformation, however, in most cases is relatively slow due to the small difference in free energy (about 2 kJ/mol) between the and -SiC phases.23 The polytype of the starting powder is not important for the densification of the SiC ceramics. It has been reported that near theoretical densities can be achieved with the appropriate choice of sintering additives, sintering time and temperatures.5 It has also been shown that the -tophase transformation does not contribute to the densification of the SiC ceramics.4 The starting polytype, however, severely affects the resulting microstructure. SiC sintered from -SiC starting powders usually shows equiaxed grain morphology.4 Although grain growth usually occurs, the aspect ratio of the larger grains remains near unity. This is particularly true for starting -SiC powders of unimodal grain size distribution. In cases where the grain size of the starting powder is intentionally

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16 varied, i.e., the grain size exhibits a bimodal distribution; the resulting microstructure usually shows a bimodal grain distribution distinctly different from the starting powder. In an experiments where -SiC seeded with large -SiC grains are used, the resulting microstructure shows a bimodal grain size distribution consisting of large elongated grains in small equiaxed grain matrix.45 This was attributed to the grain growth mechanism in liquid phase sintering, i.e., solution-reprecipitation. The large grains grow at the expense of the smaller grains, thus, resulting in a microstructure with large elongated grains in a matrix of small equiaxed grains. Microstructure with small equiaxed grains can also be fabricated from a -SiC starting powder. This is possible if the -SiC starting powder is sintered in a temperature and time range where the -tophase transformation was prevented.2 SiC sintered from -SiC starting powders processed in an overpressure of nitrogen, with AlN or AlN-Y2O3 additive also shows the same equiaxed grain morphology when processed at temperatures below 2000C.8,14 This is believed to be due to the effect of the oxynitride phase that forms during processing which extends the stability regime of the -SiC phase. If the -tophase transformation was allowed, the resulting microstructure usually contains plate-like or elongated grain morphology.8 If the starting -SiC powders are seeded with similarly sized or -SiC, the resulting microstructure still shows a unimodal grain size distribution.45 With a different grain size seeds, the resulting microstructure shows a bimodal grain size distribution consisting of large elongated grains in a matrix of small elongated grains.45 This is especially true for -SiC starting powders of grain size in nanometer range seeded with up to 5 wt. % of micron-size grains.51

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17 Effect of Sintering Atmosphere The sintering atmosphere can have several important effects on densification and microstructure development during sintering. In many instances, the atmosphere can have a decisive effect on the ability to reach a high density with controlled grain size. The important effects of sintering atmosphere are associated with the gas solubility and the chemical reaction with the powder system. Sintering atmosphere reactions with powder system is especially true for SiC since volatility of reaction products in this material, depending on the additive, is an issue during sintering. In sintering SiC ceramics, densification to near theoretical density are readily achieved with the use of argon and nitrogen atmosphere.12 Nitrogen atmosphere, in most cases, is used to address the issue of escape of gaseous reaction products during processing.30 For SiC sintered with Al2O3-Y2O3-SiO2 additive (SiO2 is always present on the surface of SiC powders), the reaction of Al2O3 and SiO2 with SiC at high temperatures generate gaseous reaction products.31 As reported in the literature,23 the use of powder beds of composition similar to the powders being sintered is generally required to attain a high density ceramics. A change in additive, i.e., AlN-Y2O3 system, and an overpressure of N2 gas provides a way to control the gaseous reaction products resulting in a high density final ceramics.23,30 In some cases, the use of nitrogen as sintering atmosphere results in the absorption of nitrogen in the grain boundary and triple junction phases. The oxynitride phase is believed to increase the stability of -SiC such that the -tophase transformation can be prevented.8 The effects of argon and nitrogen as sintering atmosphere in the resulting microstructures of SiC ceramics are patent. For -SiC starting powders of uniform size distribution, it has been observed that sintering in argon atmosphere produces equiaxed

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18 grains in the microstructure.11,12 If the starting powder, however, is seeded with large -SiC grains, the resulting microstructure shows a bimodal distribution where large elongated grains exists in a matrix of small elongated grain.11 For a -SiC starting powders sintered in argon atmosphere, the resulting microstructure usually shows presence of elongated grains.4 This is due to the -tophase transformation that easily occurs in argon atmosphere. If the -tophase transformation is made to occur to completion, the resulting microstructure has a unimodal distribution of elongated grains. If the phase transformation is incomplete, the resulting microstructure usually has a bimodal grain distribution wherein the elongated -SiC grains is contained in the matrix of equiaxed -SiC grains. In a nitrogen sintering atmosphere, the resulting microstructure using -SiC starting powders is similar to that observed in materials processed in argon atmosphere.11 The grain morphology is usually equiaxed and grain growth is minimal. In the case of -SiC starting powders, the resulting microstructure is different than those observed in argon.12 The use of nitrogen is believed to extend the stability of -SiC such that the -tophase transformation is shifted to a higher temperature.14 Thus, at low temperature, below 2000C, the resulting microstructure usually shows equiaxed grain morphology. At high temperature, above 2000C, the resulting microstructure generally shows presence of elongated grains. If the additives used, however, can form a solid-solution with SiC, example of which is AlN, the resulting microstructure varies.35-37 If the resulting phase in the microstructure is a combination of 6H-, 4Hand 2H-SiC structure, the resulting morphology shows a bimodal grain size distribution. 6Hand 4H-SiC generally appear elongated and 2H-SiC shows equiaxed grain morphology. If the phase transformation is

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19 driven to form the 2H-SiC structure, the resulting microstructure generally shows equiaxed grains. Effect of Annealing Conditions Annealing in the fabrication of SiC is an additional step in the processing wherein further grain growth and phase transformation can be achieved. It has been demonstrated that appropriate choice of annealing conditions can greatly enhance controllability of the resulting microstructure.5 Lee et al5 fabricated SiC ceramics from -SiC powders sintered with SiO2 and YAG-forming additive. Upon hot pressing at 1850C for one hour under a 25 MPa load in argon atmosphere, the resulting microstructure of the densified material showed relatively fine equiaxed grains. Annealing at 1950C in flowing argon for 4 hours resulted in a self-reinforced microstructure consisting of large elongated grains and relatively fine elongated grains. Introduction of external pressure during annealing affect the microstructure of the sintered ceramics. Kim et al6 hot pressed SiC ceramics at 1750C for 40 minutes under 25 MPa in argon atmosphere and further annealed the resulting monolith at 1850 to 1950C with or without applied pressure. The hot pressed material showed fine, equiaxed grains since the hot pressing temperature is low enough for the -tophase transformation to occur. Upon annealing for 4 hours, without applied pressure, the resulting microstructure showed elongated grains due to the -tophase transformation that occurred. The densities measured are lower than those materials annealed with pressure. The average grain diameter and aspect ratio were observed to increase with annealing temperature, however, the grain growth comes mainly from an increase in aspect ratio. For materials annealed with applied pressure at temperatures below 1950C,

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20 the resulting microstructure showed equiaxed grains with relatively low aspect ratios. At 1950C, elongation of some grains occurred resulting in a duplex microstructure consisting of large elongated grains and small equiaxed grains. The aspect ratios of the elongated grains, however, are lower than those observed in materials annealed without pressure. These results have shown that annealing under an applied pressure generally inhibits grain growth and the -tophase transformation, which can provide a new strategy for control and optimization of mechanical properties in SiC ceramics. Microstructure and Mechanical Property Modification in Sintered SiC It is a known fact that optimization of mechanical properties of materials can be achieved via the control of microstructure in the final product. The mechanical properties desired depend on particular applications that the material is being considered. In the case of SiC, this material has been used, and considered, in several structural applications because of its superior properties in terms of wear, corrosion, high-temperature creep, and oxidation resistance, as well as its high temperature retention characteristics. The drawbacks that make SiC inapplicable in some of its envisioned structural applications, as in most polycrystalline ceramics, are the low fracture toughness and extremely flaw-sensitive strength. The latter is a manufacturing/processing issue and can generally be improved by a post-processing treatment while the former can be addressed via the control of the microstructure of the final product. Discussion on what has been done to address the fracture toughness issue of SiC will be presented here. The post-processing treatment that has been used to address the issue of flaw-sensitive strength will be presented in latter section.

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21 Fracture Toughness Enhancement in SiC Cao et al2 have used -SiC starting powders and took advantage of the -tophase transformation that produces platelike microstructure in the resulting monoliths in producing a high toughness SiC. The -tophase transformation were further enhanced and promoted through the addition of aluminum, boron and carbon additives. The resulting microstructure was controlled through the variations in sintering time and temperature parameter during processing. The measured fracture toughness of the resulting ceramics reached a value of 9.5 MPam1/2, the highest value reported so far. The microstructure of the ceramics exhibiting this fracture toughness value showed mostly elongated grains. Cao et al have also shown that grain growth observed correspond to an increase in the fracture toughness and bend strength. The fracture toughness of the SiC ceramics has also shown a strong dependence on the aspect ratio of the elongated grains. The bend strength, on the other hand, shows no sensitivity towards the aspect ratio. The mechanisms responsible for toughening in ABC-SiC were reported to be due to crack bridging and subsequent grain pullout, with minor contribution from crack deflection. Cao et al observed partially debonded platelike grains, which resided just behind the crack tip, bridging the crack, thereby reducing the effective stress intensity. Based on SEM of crack path propagation on ABC-SiC sample, crack deflection due to the presence of elongated grains was also apparent. These mechanisms were not observed to be operable in materials hot pressed at lower temperature, 1700C, where the microstructure of the resulting ceramics generally shows equiaxed grains. Thus, the toughness measured from these materials is generally lower. For materials sintered at

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22 higher temperatures, 1900 to 1950C, resulting microstructure shows elongated grains with higher aspect ratio. This kind of microstructure was believed to enhance the effect of the toughening mechanisms operable resulting in a significant improvement in the fracture toughness measured. In another paper, Moberlychan et al54 compared the microstructure and mechanical properties of several SiC materials. They have found that higher fracture toughness was associated with intergranular mode of fracture and high aspect ratio of elongated grains in the SiC. It was found that the intergranular mode of fracture is enhanced by the presence of weaker amorphous secondary phases in the grain boundary and triple junctions of the studied material. Thus, Moberlychan et al concluded based on the comparison made that high toughness in SiC ceramic requires an amorphous grain boundary layer with a chemistry to weaken the grain boundary and a grain shape with a high aspect ratio. Lack of any of these ingredients produced inherently brittle materials. The additives used in sintering SiC ceramics affect the resulting microstructure which in turn can affect the fracture mode. As shown in ABC-SiC system reported by Zhang et al,53 a change in the amount of aluminum can result in a different microstructure. For aluminum addition of 3 to 7 wt. %, the resulting microstructure shows equiaxed and elongated SiC grains. The elongated SiC grains show an increasing aspect ratio as the amount of aluminum was increased. The corresponding volume of the elongated grains in the material, however, decreased as the aluminum content was increased such that at higher aluminum content (6-7 wt.%), the resulting microstructure has a much higher volume of fine equiaxed grains. Because of this change in the microstructure, a noticeable change in the fracture mode between 5 and 6 wt. %

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23 aluminum was observed. For Al additive less than 5 wt. %, the observed fracture mode was primarily intergranular, while starting at 6 wt. %, the fracture mode observed are predominantly transgranular. Changes in the microstructure and fracture mode were also observed for ABC-SiC at lower Al content. Flinders et al18 found, for a similarly processed -SiC powders, that a transgranular to intergranular mode of fracture occurs between 1 and 1.5 wt. % aluminum additions. The corresponding microstructure changes from a primarily equiaxed grains to a primarily elongated ones with a reportedly lower aspect ratio than reported by Zhang et al. Other attempts in improving the fracture toughness of SiC ceramics processed from -SiC starting have been reported in the literature.4,7,9 The achieved fracture toughness values, however, are all lower than those measured from ABC-SiC system but still significantly higher than those observed for solid-state sintered SiC. Hilmas and Tien35 measured a fracture toughness of 8.5 MPam1/2 through the use of 4.95 wt.% Al2O3 5.97 wt.%AlN -0.5 wt.% B additives hot pressed at 2100C for one hour. The resulting microstructure showed mostly elongated grains, due to the 3C-to-6H transformation, and minor fine equiaxed grains, due to the 3C-to-2H transformation. The mode of fracture observed was intergranular. Aside from crack bridging and crack deflection mechanism of toughening, interfacial microcracking due to substantial residual stresses at the grain boundaries which resulted from thermal expansion anisotropies have been considered operable in this particular material. This was further supported by the measured hardness value which is significantly lower for SiC ceramics. The low hardness value was attributed by the authors to the possible microcracking that occurred during cooling of the ceramics.

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24 Improvement in the fracture toughness of SiC ceramics were also observed by Lee et al5 who reported a fracture toughness value of 8.5 MPam1/2. In their case, the sintering additive used for the densification of -SiC was 5.7 wt.% Al2O3 3.3 wt.% Y2O3 1 wt.% CaO. The reported toughness value was from a ceramic hot pressed at 1850C for 1 hour under 25 MPa in argon atmosphere. The hot pressed ceramic was further annealed for 4 hours at 1950C under flowing argon. The resulting microstructure was reported to contain mostly of elongated grains with grain size and aspect ratio of 1.25 m and 3.9, respectively. The mode of fracture in the reported material is predominantly intergranular. Hot pressing of -SiC using Al2O3 rare earth oxides also proved to be useful in improving the fracture toughness of sintered SiC, as shown by Zhou et al.10 Although the reported toughness values are relatively lower than those reported for -SiC starting powder, the reported values are only from the hot-pressed specimens. The microstructure of these materials generally showed equiaxed grains and XRD analyses revealed that most of the grains are untransformed -SiC. If the -totransformation were made to occur, the resulting microstructure will be different and will possibly improve the fracture toughness, as was shown by other researchers. The use of similar additive system on the densification of -SiC powder seeded with large -SiC at different additive amount and processing temperature resulted in distinct mechanical properties. Zhan et al have used 7 wt.% Al2O3 2 wt.% Y2O3 1 wt.% CaO as sintering additive on a 90 nm sized -SiC powder.4 The resulting fracture toughness values are relatively lower than those reported by Lee et al.5 The hardness values, however, are relatively higher, even for the as hot-pressed material. The

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25 microstructure for the as hot-pressed material generally showed equiaxed grains. Further annealing, resulted in a microstructure with uniformly elongated grains. Improvement in the fracture toughness of SiC can also be achieved through the use of -SiC starting powders. Although -tophase transformation resulting in elongated grains will not occur when using -SiC, elongated microstructure can also be produced through the anisotropic grain growth. Kim et al were able to achieved a fracture toughness ranging from 5.4 to 6.2 MPam1/2 using -SiC starting powder sintered with YAG and YAG-SiO2 additives.6 The resulting microstructure of the hot pressed materials generally showed equiaxed grains with aspect ratio near unity. Upon annealing at 1950C for 4 hours under a flowing argon atmosphere, the resulting microstructure showed a self-reinforced morphology consisting of large elongated grains in a matrix of small elongated grains. Hot pressing of -SiC using different additives, as reported by Flinders et al and Nader et al,7 also resulted in a possible improvement in the mechanical properties. Flinders et al hot pressed -SiC using Al, AlN, and Al2O3 additive and showed minor improvement in the fracture toughness.26 The values reported, however, were from as-hot pressed specimen and possible improvement can still be achieved through annealing. Nader et al, on the other hand, used Y2O3 AlN and showed that pressureless sintering of -SiC results in moderate toughness values, even at longer sintering times (14 hours). This is in contrast to their result using -SiC starting powders where marked improvement in the fracture toughness of the SiC ceramics, even with the use of pressureless sintering, was achieved.

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26 To summarize, the microstructural toughening of SiC can be achieved through the use of -tophase transformation, the use of liquid phase forming additives, seeding or controlling the initial -SiC, and annealing for controlled grain growth. The -tophase transformation results in a development of in-situ or self-reinforced microstructure which enhances the toughening mechanisms operable. The use of liquid phase forming additive and seeding generally enhances the -tophase transformation resulting in a toughened ceramics. The low amount of additive used in some of the studies reported generally results in an improvement in the resulting toughness and hardness values. The use of annealing, even in ceramics sintered from -SiC powders, results in a pseudo-self reinforced microstructure which also shows promise in improving the toughness of the material. Crack Healing as a Tool for Mechanical Property Improvement in SiC The phenomenon of crack healing has been observed in a variety of different ceramic materials including single crystal,66 polycrystalline ceramics,67-71 inorganic glasses,72 and ceramic composites.73-92 It has also been observed in metals93 and polymer materials.94 This phenomenon is usually described as a process where the surface crack of a given material closes, either by crack-face rebonding or by filling up of the crack face, after a given heat treatment. Associated with the crack closure is an observed increase in the flexural strength of the material. The recent increase in the number of studies, in the late 1990s,73-92 on crack-healing is driven by the promise of recovering the mechanical properties of structural materials that are compromised due to inherent flaws introduced during processing and handling. Application of a crack healing treatment as a post-production process will reduce, if not entirely eliminate, the flaws/cracks introduced during processing and

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27 handling. Crack healing will greatly enhance the reliability and integrity of the structural material in question. It will also decrease the machining, maintenance, and inspection costs and increase the lifetime of the ceramic material. Most of the crack healing studies in the literature were done in an oxidizing environment. 73-92 In most cases, complete recovery of the flexural strength of the indented sample can be achieved within an hour of heat treatment. The mechanism for the crack healing is usually attributed to the formation of the oxidation products in the crack surface, bonding the crack walls and rebonding the crack faces. The increase in the flexural strength observed is usually attributed, aside from the reduction in flaw sizes, to the stress relaxation upon heat treatment and possible crack tip blunting. On the belief that crack healing by heat treatment is a phenomenon which occurs for all material that sinter, F.F. Lange, K.C. Radford and T. K. Gupta showed that crack healing was possible for ZnO, SiC and polycrystalline Al2O3.66,68-69,95 In their experiments, deep surface cracks were introduced to the specimens by quenching into water at room temperature. The thermally shocked specimens were then heat treated in air at a temperature just below the sintering temperature of the starting SiC, ZnO, and Al2O3 powder. For the case of SiC and Al2O3, the specimens were heat treated for one hour at 1400C and 1700C, respectively. In all cases, the flexural strength measured after heat treatment showed strength values higher than the un-oxidized thermally shocked specimens suggesting that crack healing occurred. For longer heat treatment times, strengths higher than that of the as-received specimen were observed. Microscopic examination of the crack patterns after heat treatment also revealed that the crack patterns

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28 disappeared, and in the case of SiC, formation of an oxide layer on the crack surface was evident. In the study of the crack healing behavior of SiC, F. F. Lange used two types of SiC, a low density (80 % of theoretical) and a high density form.69 For the low-density SiC, flexural strength measured after oxidation yielded a higher value than that of the un-oxidized quenched specimens. Flexural strengths measured after 90 hours of exposure were even higher than the flexural strength of the as-cut specimens. Results for the high density material, on the other hand, showed a minimal (10% of the unoxidized quenched specimen) flexural strength increase for the same oxidation time. Rationalization for the difference in strength recovery was provided by the difference in the oxidation behavior of the material, i.e., the high density form has experienced very little oxidation. These observations, coupled with the observed oxidation of the cracked surface, suggest that the formation of the oxide film was responsible for crack healing. Crack-healing behavior also depends on how the ceramic material was manufactured. This was shown by the studies performed on a reaction-bonded and liquid phase-sintered silicon carbide.95-96 Both of these materials showed crack healing ability, however, the strengthening achieved for liquid phase-sintered material was higher. In both materials, the crack healing ability was attributed to the flow of the preexisting glass phase across the crack plane, and the reaction of the crack surface with the environment (oxidation). The difference in strengthening was attributed to the different residual stress produced by the thermal expansion mismatch between the oxidation product and the bulk material. For the reaction bonded silicon carbide, the residual stress is lower since above

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29 500C, silicon becomes ductile making dislocation movement possible. Thus, the residual stress contribution for this material will be due to cooling from this temperature. Systematic studies of the crack healing behavior of several ceramic and ceramic composites geared towards improving the reliability and integrity of these materials were only accomplished in the late 1990s by several Japanese investigators.73-92 They performed crack-healing behavior evaluations on the following materials; silicon carbide, silicon carbide-reinforced silicon nitride composite, mullite/silicon carbide ceramics and silicon carbide whiskers toughened alumina. Several issues including: a) the effect of chemical composition on the crack healing ability, b) the effect of healing conditions on the strength of the healed-zone, c) determination of the maximum crack size which can be healed, d) knowledge of the high-temperature strength of the healed zones, e) understanding of the crack healing mechanisms and f) assessment of the cyclic and static fatigue strengths of crack-healed ceramic member, were also addressed in their investigations. Results of the investigation of the Japanese group are summarized here. In most of the materials that they studied, they found out that surface crack sizes of up to 200 m can be healed, if the crack-healing was done under the optimal healing conditions. Most of these crack-healing conditions are in the temperature range of 1200 to 1500C, in air, for a heat treatment time of 1 hour. In the best healing conditions, the cracked-healed materials flexural strength is either higher or comparable to the flexural strength of the smooth samples. Failure during the flexural strength testing occurred mostly outside the Vickers indentation zone, suggesting that the healed zone has sufficient strength. Variations of the flexural strength of the crack-healed specimens are generally the same

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30 as that of the smooth specimens. It was also shown that most of the crack-healed specimens are not sensitive to static fatigue testing below an applied stress of 70-75% of the average monotonic bending strength. Cyclic fatigue testing of Si3N4/SiC (R = 0.2, frequency = 8 Hz, N = 2 x 106 cycles) also showed that the maximum stress at which the crack-healed specimens did not fracture is usually about 3 3.5 times that of the crack specimens, and higher than those for smooth specimens. It was also shown that most of the materials tested are capable of crack-healing even in the presence of cyclic stress. These results suggest that an improvement in structural reliability can be achieved via crack-healing, with the primary mechanisms being the formation of oxidation products which bonds and strengthen the crack surface. The observed flexural strength increase, on the other hand, was generally attributed to stress relaxation and crack tip blunting.

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CHAPTER 3 STATEMENT OF THE OBJECTIVE The general objective of this investigation is to verify the correlation between grain boundary and triple junction chemistry and mechanical properties of several SiC ceramics sintered with minimum additives for armor applications. In order to accomplish this objective, the chemistry of the triple junction and grain boundary phase will be determined through the combination of the following techniques: High resolution transmission electron microscopy to determine the presence of amorphous intergranular films, energy dispersive spectroscopy and energy filtered transmission electron microscopy to determine the composition of the grain boundary/intergranular films and triple junction phase, and electron energy loss spectroscopy to determine the exact nature of the triple junction phases. Several sets of SiC ceramics sintered with the addition of aluminum, either in the form of aluminum metal or aluminum compounds, will be investigated. SiC ceramics fabricated from -SiC starting powders with fixed boron and carbon content and varying amounts of aluminum addition will be studied to determine the effect of amount of aluminum on the grain boundary and triple junction characteristics. Mechanical properties of these materials, particularly toughness, hardness, and mode of fracture, will also be determined. Correlation between the mechanical properties and the grain boundary and triple junction characteristics will also be explored. SiC ceramics fabricated from -SiC starting powders with equimolar aluminum additions (from Al, AlN and Al2O3) will be investigated to determine the effect of 31

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32 sources of aluminum on the grain boundary and triple junction chemistry. Effect of boron and carbon addition on the microstructure and grain boundary and triple junction chemistry will also be explored. SiC fabricated from -SiC starting powders with aluminum or aluminum nitride addition in differing amounts and hot pressed at different temperature will also be studied. This is to determine the effect of amount of additive and processing temperature on the resulting microstructure and mechanical properties. Grain boundary and triple junction characteristics will also be investigated in these materials to determine the effects of aforementioned variables. The implications of the observations on these materials on the microstructure and property tailoring in SiC ceramics will also be discussed. Comparison with a commercially available state-of-the-art ceramic armor will also be made. Strength improvement through crack healing in SiC ceramics will also be demonstrated. Although many investigations have been carried out on the microstructure and property tailoring of SiC, data on grain boundary and triple junction phase characteristics of this material are very limited. The investigations performed in this dissertation aim to contribute to the literature in this area. SiC ceramics sintered with minimum amount of aluminum (Al, AlN and Al2O3) are used to demonstrate that high toughness and hardness are achievable at low level of additives.

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CHAPTER 4 MATERIALS AND METHODS The materials processing and characterization methods used in this investigation are presented in this chapter. The materials used in this study were all fabricated by Ceramatec, Inc. The mechanical properties (hardness, toughness and mode of fracture), chemical analysis and SiC polytypes characterization were also performed by Ceramatec, Inc. Electron microscopy such as energy dispersive spectroscopy (EDS) and high resolution transmission electron microscopy (HRTEM) were done at MAIC at the University of Florida. Energy filtered transmission electron microscopy was performed at MCF at the University of Central Florida. Materials The materials used in this investigation are presented in Table 4-1. The materials were grouped in five classifications according to the type of starting SiC powders, kind of additive and intended experiments. ABC-SiC system was investigated to determine the effect of amount of Al additive on the grain boundary and triple junction. It was also intended to find any correlation between the grain boundary and triple junction characteristics, and the mechanical properties and mode of fracture in these materials. Materials fabricated from -SiC were used to investigate the effect of different sources of aluminum additive, hot pressing temperature and amount of additive and find an alternative way to tailor properties similar to commercially available ceramic armor material (SiC-N). Additions of B (in form of B4C) and C were also performed to 33

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34 investigate effect of these additives on the mechanical properties and grain boundary and triple junction characteristics. Table 4-1. Compositions and processing conditions of the materials investigated. Compositions (wt. %) Hot Pressing Conditions Temperature Holding Time Atmosphere/Pressure ABC-SiC System -SiC-0.6B-2C 2100C 1 hr Argon/28 MPa -SiC-0.5Al-0.6B-2C 2100C 1 hr Argon/28 MPa -SiC-1Al-0.6B-2C 2100C 1 hr Argon/28 MPa -SiC-1.5Al-0.6B-2C 2100C 1 hr Argon/28 MPa -SiC-4Al-0.6B-2C 2100C 1 hr Argon/28 MPa -SiC with Al -SiC-1.65Al 1900C 1 hr Argon/28 MPa -SiC-1.65Al 2100C 1 hr Argon/28 MPa -SiC-1.65Al 2200C 1 hr Argon/28 MPa -SiC-3.3Al 2000C 1 hr Argon/28 MPa -SiC-3.3Al 2200C 1 hr Argon/28 MPa -SiC-1.65Al-0.5B4C-2C 2100C 1 hr Argon/28 MPa -SiC with AlN -SiC-2.5AlN 2100C 1 hr Argon/28 MPa -SiC-2.5AlN 2200C 1 hr Argon/28 MPa -SiC-5AlN 2000C 1 hr Argon/28 MPa -SiC-5AlN 2200C 1 hr Argon/28 MPa -SiC-2.5AlN-0.5B4C 2100C 1 hr Argon/28 MPa -SiC-2.5AlN-0.5B4C-2C 2100C 1 hr Argon/28 MPa Comparison Materials SiC-N Proprietary Proprietary Proprietary -SiC-3.1Al2O3 2100C 1 hr Argon/28 MPa -SiC-0.5B4C-2C 2100C 1 hr Argon/28 MPa Crack Healing -SiC-0.6B-2C 2100C 1 hr Argon/28 MPa -SiC-1.5Al-0.6B-2C 2100C 1 hr Argon/28 MPa -SiC-0.43Al-0.5Y2O3 2200C 1 hr Argon/28 MPa The starting powders used for processing are listed in Table 4-2 along with the data provided by the suppliers. Ceramics fabricated from different starting powders were processed differently. In the case of materials fabricated from -SiC starting powders, -SiC powders were mixed with Al, B and C additives and dispersed with polyamine

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35 polyester polymer by adding 1 wt. % of the polymer, based on solids, to 400 grams of reagent toluene. The aluminum was added in 0.5, 1, 1.5 and 4 wt. % whereas the born and carbon were fixed at 0.6 and 2 wt. %, respectively. The carbon was added as 4 wt.% apiezon wax which would result in 2 wt. % C incorporation upon pyrolysis. The slurries were then deagglomerated for two hours with paint shaker and rolled overnight before drying. The powders were sieve through a 44 m screen before hot pressing at 28 MPa in stagnant argon inside a graphite die. The hot pressing temperature used is 2100C and holding time is 1 hour. Table 4-2. Raw materials used in powder processing. Powder Supplier Grade Information from Supplier (wt.%) -SiC Superior Graphite HSC-059 SA=15-17 m2/g -SiC H.C Starck UF-15 O=1.0%, C=29.6%, d50=0.5 m, SA=15m2/g Al Valimet H-3 Fe=0.12%, Al=99.98%, d50=5.0 m AlN Tokuyama Soda F O=0.78%, C=0.03%, SA=3.4 m2/g Al2O3 Sasol North Am. SPA-0.5 Si=14 ppm, d50=0.47 m, SA=7.6 m2/g Y2O3 Molycorp 5600 B H.C Starck S-432B B4C H.C. Starck HS O=1.5%, B/C=3.80, d50=0.96 m, SA=18.0 m2/g C Apiezon W 50 wt.% yield after pyrolysis C Capital REsin CRC-720 40 wt.% yield after pyrolysis In the case of SiC ceramics fabricated from -SiC starting powders, all compositions were prepared by batching 600 grams of powder in two-liter high-density polyethylene (HDPE) jars filled with 1.6 kg solid state sintered SiC media and 700 g reagent grade acetone. The slurries were mixed for 16 hours on a ball mill in order to disperse the agglomerates. The powders were stir dried before screening through an 80-mesh screen. Compositions batched with B4C and C were pyrolyzed by heating in N2 to 600C. Billets (45 mm x 45 mm x 6 mm) were hot pressed at 28 MPa inside graphite

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36 dies by heating to 1500C in a vacuum of 1 torr, holding at 1500C for one hour to remove SiO and CO from the samples, and then backfilling with Ar and heating to 2100C for one hour. The vacuum hold was performed to eliminate porosity in the materials. SiC-N is a commercially available SiC from Cercom, Inc. produced via the pressure assisted densification method. Characterizations Mechanical Properties and SiC Polytypes Toughness and hardness measurements were performed on a test bar of dimension 3 mm x 4 mm x 45 mm. The bars were pre-cracked with a commercially-available fixture (Maruto model MBK-603C) after using a single 98 N Knoop indenter to initiate the crack. The tests were stopped after an audible pop-in noise was detected, with loads ranging from 4 kN to 18 kN using bridging spans of 4, 5, or 6 mm. The original crack was marked by a dye (black ink from an inkjet printer), using vacuum infiltration after unloading the sample. The dye was oven dried at 50C overnight and cooled before testing. Fracture toughness measurements were conservative since the crack location used in the calculation was prior to the small amount of stable crack growth associated with the SEPB tests. All crack planes were parallel to the hot pressing direction. Each data point reported is the mean of 3-7 bars tested, with error bars representing one standard deviation. All microhardness data (Leco Model LM-100) were obtained using a one kilogram load Vickers (HV1) on polished surfaces. The dwell time at load was nominally 15 seconds. Rietveld analysis97-98 was used to determine SiC polytypes present in the

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37 densified samples with X-ray diffraction patterns collected from a diffraction angle of 30-80, with a step size of 0.02/step and a counting time of 4 sec/step. Microstructure In the case of SiC materials fabricated from -SiC starting powders, the grain size and aspect ratio reported are based on measurements from the polished and etched surface of the sample. Polished samples of materials containing aluminum (0.5, 1, 1.5 and 4 wt. %) were plasma-etched by evacuating and back-filling with 400 millitorr of CF4-10%O2 and etched for 20 to 40 minutes. The solid state sintered material (no aluminum) was etched in molten KOH at 550C for 10 to 15 seconds. Grains size was determined by line-intercept method, where the multiplication constant ranged from 1.5 (for equiaxed grains) to 2.0 (for elongated, platelike grains).99 The mean grain size were measured from 200 to 300 grains for each composition and the aspect ratio was estimated from the five most acicular grains. In the case of SiC materials fabricated from -SiC starting powders, the grains size and aspect ratio reported were based on the measurements from a low magnification TEM images taken from the sample. Average grain sizes were determined from 20 grains and the aspect ratio was measured from at least 5 elongated grains. Chemical Analysis Chemical analysis for nitrogen and oxygen constituents on selected starting compositions and sintered samples were performed for materials sintered from -SiC starting powders. The analyses were performed based on the specification of ASTM method E1409 for the determination of oxygen by the inert gas fusion/thermal

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38 conductivity detection technique with a Leco TC-136 model analyzer. The detection range for oxygen of this analyzer is from ppm level to below 5 wt. %. Transmission Electron Microscopy Transmission electron microscopy is unique among materials characterization techniques in that it enables essentially simultaneous examination of microstructural features through high-resolution imaging and the acquisition of chemical and crystallographic information from small (usually in submicrometer range) regions of the specimen. The signals generated, collected, and analyzed in TEM are produced by interactions between the electrons in the high-energy incident beam and the material in the thin film target. Some of these signals are illustrated in Figure 4-1. The direction shown for each signal is not necessarily the physical direction but indicates where the signals are generally detected. Figure 4-1. Signals generated from high-energy electron beam/thin specimen interaction.100

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39 Electron beam/specimen interactions, or scattering events, can be divided into two categories: elastic and inelastic events. Elastic events affect the trajectories but do not significantly affect the velocities or kinetic energies of the incident electrons. This process leads to the emission of forward diffracted and transmitted electrons as well as backscattered electrons. Inelastic events, on the other hand, results in a transfer of energy to the solid with very little change in the electron trajectory. This process leads to generation of secondary electrons, Auger electrons, characteristic and continuum x-rays, visible light, electron-hole pairs, lattice vibrations and electron oscillations. A number of these signals produced by the interaction of the incident electron beam with the thin specimen are used in TEM. Elastically and inelastically scattered electrons, secondary electrons and backscattered electrons are used for imaging. Inelastically scattered electrons are also used for chemical microanalysis due to the characteristic energy loss during inelastic scattering in a technique called electron energy loss spectroscopy (EELS). The characteristic x-rays are also used for microanalysis through a technique called energy dispersive x-ray spectroscopy (EDS). TEM sample preparation Specimen used in TEM studies requires that the samples are electron transparent. This translates to a thickness of typically less than 100 nm. To achieve such, bulk specimens must be sectioned and eletrothinned or ion milled to produce regions that permit transmission of the electron. TEM specimen from the bulk SiC samples were prepared by the standard mechanical thinning method. A 3mm x 4 mm rectangle with a thickness of 1 mm was cut from each specimen test bar with a low speed diamond saw. The cut sample was then crystal-bonded to an aluminum polishing stub and wet polished to a thickness of ~ 100

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40 m. A 3mm diameter disc was then cut from the polished sample using an ultrasonic drill. The 3 mm disc was further polished to a thickness of ~ 50 m using a 3 m diamond suspension in a precision dimpling machine and a flattening tool. The specimen center was then subsequently thinned to a thickness of 20 m using a dimpling tool. At this point, the specimen center typically appears translucent. To make the specimen electron transparent, argon ion milling was then performed in a Gatan duo-mill ion milling machine. The typical parameters used for the SiC materials are: 4 kV accelerating voltage, 1 A ion gun current and a beam angle of 13. Ion milling time generally depends on the type of SiC materials and usually takes from 3 to 10 hours. Ion milling was carried out until a small perforation on the sample was observed. Due to the brittle nature of the thin section of the specimen (the thin section falls off during TEM investigation), several samples are usually made. High resolution TEM (HRTEM) High resolution transmission electron microscopy (HRTEM) is generally used to obtain lattice images. To produce a lattice image, the transmitted beam and one diffracted beam are passed through the objective aperture and combined. The phase interference between these two beams yields the periodic intensity fringes present in the image. Under the appropriate imaging conditions, there is a one-to-one correspondence between the intensity fringes in the image and the atomic planes from which the electrons were diffracted. In this case, the spacing of the fringes in the image is equivalent to the spacing of the atomic planes. Using lattice images, it is possible to examine the detail of grain boundaries, phase interfaces and to image edge dislocations. In this dissertation, the lattice images are

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41 primarily used for the examination of the grain boundaries and to determine presence of intergranular films. The lattice images needed for this analysis therefore requires that it contains details from the two grains and the interface. To achieve this, lattice image of one of the adjacent grain is acquired first and the sample is tilted and oriented such that lattice image of the other grain is obtained. In the absence of secondary phase, a direct transition from the lattice images of the two grains can be observed. If the intergranular film is present in amorphous form, absence of lattice fringes on the interface would be apparent. If the intergranular film is crystalline, lattice fringes of distinctly different appearance (different from the two lattice image of surrounding grains) could be observed. Obtaining lattice images was a tedious and time consuming process. It also requires that the grain boundaries being studied are head-on. In the case of the materials studied in this dissertation, the nature of the microstructure (interpenetrating elongated grains) severely limited the number of grain boundary investigated. For this reason, only three or four high resolution images from each specimen were taken. The high resolution images presented in this dissertation was generated on a JEOL JEM-2010F FEG with a point to point resolution of 0.19 nm. The images were acquired with the help of Kerry Seibien of the Major Analytical Instrumentation Center (MAIC) at the University of Florida. Energy dispersive spectroscopy (EDS) The EDS data presented in this dissertation was taken with a JEOL 2010F equipped with an Oxford INCA 200 system. The EDS spectra were generated under STEM mode and used a probe size of approximately 5 nm. The electron beam hitting the sample, however, is probably about 8 nm (based on the burn patterns observed on investigated

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42 samples) due to the inhomogenieties in the electron delivery system. The size of the probe used also limits the results obtained to a merely qualitative guide, i.e., only indicate the elements present. This is particularly true for the compositions obtained on the grain boundary since the grain boundary width (or thickness of the intergranular film) observed is generally ~ 1 nm. The advantage of EDS relative to other microanalytical technique is its ability to generate x-ray spectra from, and therefore determine the composition of, very small volumes of the specimen. Characteristic x-rays are generated from the interaction of energetic electron and an atom. A sufficiently energetic electron can interact with an atom and cause the ejection of tightly bound inner-shell electron, leaving the atom in an ionized and highly energetic state. During subsequent de-excitation, an electron transition occurs in which an electron from an outer shell drops inward to fill the inner-shell vacancy. The change in energy (or the energy released) due to this electron transition will be in the form of x-ray or an ejected outer-shell electron. Since the electron structure of each atom is unique, the x-rays generated will also be unique and a characteristic of a given element. The decision to use the EDS results presented in this dissertation for qualitative purposes only lies on the limitation of this technique. Since the EDS were taken from TEM specimens, the number of x-ray counts generated will be minimal. The x-ray generation volume will be smaller due to the thickness of the specimen. Thus, the elemental detectability limits will be greatly diminished. This problem is compounded by the elemental constituents of the materials being investigated. In the case of ABC

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43 SiC, the amount of B in the sample is only 0.5 wt. % (0.6 wt. % in -SiC), which was borderline with the quoted detectability limits in JEOL 2010F. Energy filtered transmission electron microscopy (EFTEM) EFTEM is a technique based on electron energy loss spectroscopy (EELS). EELS involves analysis of the energy distribution of the inelastically scattered electrons in the transmitted beam. Information pertaining to the chemical and bonding from the atoms in the sample can be determined through this technique. During EELS, all the inelastically scattered electrons are detected making the signal intensity much higher than in EDS where characteristic x-rays just comes from small portion of inelastic scattering events. It is thus a better technique than EDS for quantitative analysis. It also offers an advantage in detecting low Z elements. In EFTEM, an imaging filter is used to generate an image consisting of electrons with specific energy losses. The image filter only allows electrons with an energy loss of E E/2, where E is the selected energy and E is the slit width of the energy window. It is therefore possible to make chemically selective images from a given sample. Due to the non-element specific background in an EELS spectrum, however, EFTEM images can not be directly interpreted as elemental maps. Background removal has to be accomplished first. A method for removing background EELS spectrum and achieving elemental maps from EFTEM images used in this dissertation is a technique called three window elemental mapping. Three window elemental mapping allows generation of quantitative images wherein the intensity is proportional to the EELS spectrum intensity under a certain excitation edge. The schematic overview of this technique is shown in Figure 4-2.

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44 The technique uses three images made up of electrons coming from three different energy regions in the spectrum. Two pre-edge images are used to estimate the non-specific background in each pixel of the image, by estimating the parameters of the background model AE-r for each pixel. The third image (post-edge) is obtained from the energy region which contains the excitation edge of interest. The estimated background from each pixel is then subtracted from the post-edge image and an excitation specific image remains. E Figure 4-2. Schematic overview of the three-window technique.101 The EFTEM elemental maps presented in this dissertation was taken using a Tecnai F30 equipped with FEG and GATAN GIF. The accelerating voltage used in the TEM is

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45 300kV. All the elemental maps were taken with the help of Dr. Helge Heinrich in Materials Characterization Facility at the University of Central Florida.

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CHAPTER 5 RESULTS OF MECHANICAL PROPERTY, GRAIN BOUNDARY, AND TRIPLE JUNCTION CHARACTERIZATION Due to the large number of materials investigated, the results chapter of this manuscript will be subdivided into four main sections based on the types of additive used. The first section will cover the results for -SiC sintered with 0.6 wt.% B 2 wt.% C and 0 4 wt.% Al. The second section will cover the results for -SiC sintered with Al additive, while the third section will deal with the results obtained for -SiC sintered with AlN. The last section will present the result for -SiC sintered with 3.1 wt.% Al2O3, the solid state sintered -SiC and the commercially available material SiC-N. Discussions pertinent to the results presented in this chapter will be relegated to the next chapter of this dissertation. -SiC Sintered with 0.6 wt.% B, 2 wt.% C and 0 4 wt.% Al (ABC-SiC) Processing and Polytypes The SiC ceramics fabricated from -SiC starting powders with Al, B and C additions sintered to near-theoretical density of pure SiC. The theoretical density of SiC was calculated to be 3.2 g/cc. As shown in Table 5-1, the material with only B and C additions gave a value near the theoretical density of SiC. Increasing addition of aluminum has yielded density values which are progressively lower, with the highest amount of aluminum showing the lowest density values. The polytypes of SiC present in the manufactured ceramics, based on the Reitveld analysis performed on the XRD patterns obtained from the material, are also shown in 46

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47 Table 5-1. Hot pressing at 2100C for one hour enabled the complete -tophase transformation as shown by the absence of the cubic 3C polytype in the sintered material. When only B and C were used as additives, the -tophase transformation was reported to begin at temperatures greater than 1950C.2,102-104 The addition of a small amount of metallic aluminum or aluminum compounds can lower this phase transformation temperature. As reported by Shinozaki et al.105 and William et al.,106 the onset of the -tophase transformation at approximately 1800C is even possible. Since the temperature used in hot pressing the -SiC in this study is higher than 1950C, it is not surprising that complete -tophase transformation was achieved. Table 5-1. Density and polytypes of SiC ceramics from -SiC starting powders. Sample (wt. %) Density (g/cc) Phase Assemblage (wt. %) 3C 4H 6H 15R -SiC 0.6B -2C 3.17 0.01 0.0 0.0 93.1 6.9 -SiC 0.5Al 0.6B -2C 3.08 0.01 0.0 0.0 87.5 12.5 -SiC 1Al 0.6B -2C 3.15 0.01 0.0 22.4 68.2 9.4 -SiC 1.5Al 0.6B -2C 3.13 0.01 0.0 34.0 58.9 7.1 -SiC 4Al 0.6B -2C 3.12 0.01 0.0 81.4 14.5 4.1 The polytype present in the manufactured ceramics varies with the amount of aluminum added. When only B and C are added, the resulting polytypes are mostly 6H, with the remainder being that of the rhombohedral 15R structure. As the amount of aluminum is increased, the amount of 4H polytype in the fabricated ceramics also increases, with a corresponding decrease in the amount of 6H polytype present. This observed trend is similar with those reported for -SiC with Al-B-C addition processed at 1900C.2

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48 Microstructure and Mechanical Properties The resulting microstructure of the hot-pressed -SiC generally shows interlocking platelike grains. As reported in the literature,2-4 formation of the platelike grains in specimens that utilized -SiC starting powders is due to the -tophase transformation that occurred during sintering. As shown in the previous section, complete -tophase transformation was achieved in the materials investigated, thus, the microstructure is expected to consist of mainly elongated grains. Examples of the observed microstructure in sample with 0, 0.5 and 1 wt.% aluminum are shown in Figure 5-1. c) b) a) Figure 5-1. Transmission electron micrographs of the microstructure observed in SiC ceramics fabricated from a -SiC starting powder. a) microstructure found in -SiC-0.6B-2C, b) microstructure found in -SiC-1Al-0.6B-2C, and c) microstructure found in -SiC-1.5Al-0.6B-2C. The marker shown in each micrograph is for 2 m. The measured grain size, aspect ratio, hardness, toughness and observed failure mode of the SiC materials sintered from -SiC starting powders are shown in Table 5-2. It is observed that correlation between the grain size and the amount of aluminum added does not exist for the materials investigated. The grain size is smallest (4.10.5 m) for the sample with 1 wt.% Al added. This value, however, is not significantly different than those measured for -SiC with 0.5 wt.% Al addition if one considers the standard deviation of the measured grain size. For the material with no Al additive, the grain size

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49 measured is similar to those of otherwise similar material with 1.5 and 4 wt. % Al additions and is higher than those measured from -SiC with 0.5 and 1 wt. % Al. Thus, considering that the grain morphology in these materials is generally comprised of elongated grains, correlation between the measured grain size and the amount of aluminum added is really non-existent. Table 5-2. Grain size, aspect ratio, hardness, toughness and observed fracture mode in SiC ceramics with -SiC starting powders. Sample (wt. %) Grain Size (m) Aspect Ratio Toughness (MPa m1/2) Hardness HV1 (GPA) Fracture Mode -SiC 0.6B 2C 5.41.1 5.41.1 2.6 0.2 25.5 0.7 Transgranular -SiC 0.5Al 0.6B 2C 4.60.2 5.91.3 2.6 0.1 20.9 0.9 Transgranular -SiC 1Al 0.6B 2C 4.10.5 5.40.9 2.7 0.1 21.5 0.9 Transgranular -SiC 1.5Al 0.6B 2C 5.20.4 3.90.6 6.1 0.3 15.6 1.2 Intergranular -SiC 4Al 0.6B 2C 5.30.5 5.11.0 6.1 0.2 20.8 0.9 Intergranular The aspect ratios of the elongated grains are deemed important in determining the toughness of a given ceramic. Literature data suggests that the higher the aspect ratio of the elongated grains in a microstructure, the higher the toughness achievable.9 This is of course dependent on the toughening mechanisms operating in the material system. In the case of SiC with a microstructure consisting of elongated grains, the toughening is usually proposed to be associated with crack bridging, microcracking and crack deflection mechanisms. And in these mechanisms, the aspect ratio of elongated grains plays an important role. As shown in Table 5-2, the aspect ratio of the elongated grains in the materials investigated are almost similar, with the exception of those measured in -SiC with 1.5 wt. % Al. The toughness value, however, does not show the same trend.

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50 Additions of 0 to 1wt.% Al yielded an SEPB toughness value of about 2.6 MPa m1/2, while additions of 1.5 and 4 wt. % Al showed a higher toughness value of 6.1 MPa m1/2. This trend in toughness value has a better correlation with the observed fracture mode of the materials investigated. At low toughness values, the materials exhibited a transgranular mode of failure while the higher toughness value materials showed an intergranular fracture mode. Hardness/toughness trade-off common in liquid phase sintered SiC ceramics was also observed in the materials investigated. As shown in Table 5-2, high toughness material generally shows a slightly lower hardness value while the low toughness material shows a higher hardness. The hardness/toughness trade-off is readily apparent when one considers the case of -SiC with no Al addition and -SiC with 1.5 wt. % Al additive. In -SiC with no aluminum addition, a Vickers hardness of ~25.5 GPa and SEPB toughness of ~2.6 MPam1/2 were measured. In -SiC with 1.5 wt. % Al, Vickers hardness of ~15.6 GPa and the SEPB toughness of ~6.1 MPam1/2 were reported. These two materials exhibits the extreme in the hardness and toughness measured from the materials investigated. The low hardness value of -SiC with 1.5 wt.% Al additive might be due to the presence of amorphous grain boundary and the intergranular mode of failure observed in the material which results in a possible grain boundary sliding during indentation. Achievement of high toughness at the same level of hardness was also demonstrated in the -SiC ceramics investigated. The Vickers hardness measured in the -SiC with 0.5, 1 and 4 wt. % Al are statistically similar while the toughness values are markedly different. The SEPB toughness for -SiC with 0.5 and 1 wt.% Al are 2.6 and

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51 2.7 MPam1/2, respectively, while that for -SiC with 4 wt.% Al addition is ~6.1 MPam1/2. This observation suggests the possibility of manufacturing a toughened hard SiC ceramics. Grain Boundary, Triple Junction and Other Secondary Phases To ascertain the correlation between the grain boundary and triple junction chemistry and crystallinity and mechanical properties of the SiC ceramics fabricated from a -SiC starting powder with Al, B and C additions, high resolution transmission electron microscopy (HRTEM), energy dispersive spectroscopy (EDS), and energy filtered transmission electron microscopy (EFTEM) were performed. High resolution imaging was performed to determine the presence and crystallinity of intergranular films in the SiC ceramics, while EDS and EFTEM were done to determine the chemical constituents of the grain boundary, triple junction and other secondary phases found in the SiC ceramics. Grain boundary width and crystallinity The results of HRTEM performed on the materials investigated are shown in Figure 5-2 and summarized in Table 5-3. The grain boundary width and crystallinity was observed to vary with the amount of aluminum added. At 0 to 1 wt.% Al additions, the resulting grain boundary shows no intergranular phase and direct transitions of the fringe pattern in-between grains are observed. Increasing aluminum content, from 1.5 to 4 wt. %, leads to the formation of amorphous intergranular film with thickness of approximately 1 nm. This is shown in Figure 5-2d and Figure 5-2e where the fringe patterns of the two grains shown are delineated by a region where a fringe pattern was non-existent.

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52 Formation of secondary phases is expected in the presence of sintering additives in concentrations well above their solubility limits in SiC. The solubility limit reported for aluminum in SiC is 0.26 and 0.50 wt. % at 1800 and 2000C, respectively. The solubility of boron in SiC is 0.1 wt. % at 2500C.107 Thus, based on the solubility limit of aluminum in SiC, formation of an intergranular film at the grain boundary of the -SiC with 1 wt. % Al addition should occur. Since this is not the observed phenomenon, the presence or lack thereof of the intergranular film cannot be explained in terms of the solubility limit alone. Discussion along these lines will be presented in the next chapter of this dissertation. e) d) c) b) a) Figure 5-2. High resolution transmission electron micrographs of grain boundaries in -SiC with Al, B and C additives. a) clean grain boundary in -SiC-0.6B-2C. b) clean grain boundary in -SiC-0.5Al-0.6B-2C. c) clean grain boundary in -SiC-1Al-0.6B-2C. d) amorphous grain boundary in -SiC-1.5Al-0.6B-2C and e) amorphous grain boundary in -SiC-4Al-0.6B-2C. The markers in a), b) and c) are 2 nm while the markers in d) and e) are 5 nm.

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53 Determination of the chemical constituents in the grain boundary, triple junction and other secondary phases found in -SiC materials investigated is presented in the next section. As will be shown, the intergranular amorphous film was found to contain Al and O. Table 5-3. Grain boundary width and crystallinity of -SiC with Al, B, and C additions. Sample (wt. %) Grain Boundary Width Intergranular Film -SiC 0.6B -2C 0 none -SiC 0.5Al 0.6B -2C 0 none -SiC 1Al 0.6B -2C 0 none -SiC 1.5Al 0.6B -2C ~ 1 nm Yes, Amorphous -SiC 4Al 0.6B -2C ~ 1 nm Yes, Amorphous Grain boundary, triple junction and secondary phase composition It is known that the use of sintering additives may lead to the formation of secondary phases in a given ceramic.13 The sintering additives attract impurities in the starting powder, react with the native oxide on the particle surface and form a mass transport medium during densification.13 Upon cooling, some of the mass transport medium remain and become a crystalline or glassy phase in the resulting microstructure. In the case of ABC-SiC, it has been reported that B and C does not form secondary phases in the absence of aluminum.2 Carbon reacts with the native oxide on the surface of the starting powder and form SiO and CO gas.2 This reaction reduces the amount of oxygen available in the system. Thus, in the case of -SiC-0.6B-2C, formation of secondary phases is not expected. However, if the amount of carbon additive used is more than the amount of native oxide available for the formation of SiO and CO, then, presence of residual carbon is expected. This phenomenon is shown by the results obtain for -SiC-0.6B-2C.

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54 In the presence of aluminum above its solubility limit in SiC, formation of secondary phases is expected. Due to its low melting point (~660C), the aluminum metal will melt at the processing temperature used. B, C and O will then be transported into the aluminum melt and form secondary phases. The liquid phase will then flow and fill the porosity between SiC particles and engulf many grains. Thus, upon cooling, secondary phases can be found in triple and multiple grain junctions. Due to the low vapor pressure of aluminum, it is also possible that aluminum vapor will coat the SiC particle surfaces and react with the native oxide on the powder surface as well as the added boron and carbon and form a liquid grain boundary phase. Based on the above discussion, the amount of secondary phases available in the final ceramics should vary with the amount of aluminum added when the amount of B and C are fixed. -SiC 0.6B -2C Based on the above discussion, presence of secondary phases in this material is not expected, except for possible residual carbon in the system. Formation of intergranular phases are also expected to be non-existent since boron and carbon does not form intergranular phase in the absence of aluminum and will just get incorporated in the SiC grains. As shown in Figures 5-3, 5-4 and 5-5, variation in the grain boundary and triple junction composition was not observed. However, as shown in Figure 5-3, presence of residual carbon, indicating that excess carbon additive was used in this material, was detected. Figure 5-3 and 5-4 also suggests that the grain boundary and triple junction are relatively clean in this material except for several residual carbon phase usually found in the triple junction. Although not shown here, several metallic impurities were also

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55 observed to segregate in the triple junction. This segregation of metallic impurities, however, was observed for almost all of the materials investigated in this dissertation. Figure 5-5 is an EFTEM of grain boundaries and associated triple junction and as shown, variation in the elemental map suggesting possible elemental segregation in grain boundary and triple junction was not observed. The absence of grain boundary and triple junction phases might have contributed to the low toughness value of this material. The presence of residual carbon might also act as a failure origins which may result in low fracture strength achievable in this material. Bulk Grain and Grain Boundary R es i dua l Ca r bo n Figure 5-3. STEM of a grain boundary and a secondary phase in -SiC-0.6B-2C. The accompanying EDS identify the secondary phase as residual carbon. -SiC 0.5Al 0.6B 2C Results of the EDS and EFTEM studies performed in this material are shown in Figures 5-6 to 5-9. The triple junction in this material was found to contain Al and O-rich phases, as shown in Figure 5-6. The grain boundaries observed in this material generally do not contain secondary phase. However, the region shown in Figure 5-7 show some indication that secondary phase in the grain boundary might be present. Still, the amount of aluminum and oxygen are too small to be believable. In addition, the EDS

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56 of this particular grain boundary was taken near a triple junction, thus, it is not unexpected that a greater amount of aluminum and oxygen will be seen. Bulk Grain and Grain Boundar y Tri p le Junction Figure 5-4. STEM of a triple junction in -SiC-0.6B-2C. Variation in the triple junction, bulk grain and grain boundary composition is non-existent. Figure 5-5. EFTEM of a triple junction and grain boundary in -SiC-0.6B-2C. Variation in the elemental map is non-existent. The assertion that no grain boundary phase is present can be clearly seen in Figure 5-8. As shown, the elemental maps generated by EFTEM do not indicate any variation in

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57 composition across the grain boundary. This is particularly true for the map of aluminum, oxygen and boron. As observed in -SiC-0.6B-2C and most of the materials studied, metallic impurities are also present in this material as shown in Figure 5-9. In addition, the presence of B and C-rich secondary phase was also observed. Bulk Grains Triple Junction Figure 5-6. STEM of a triple junction in -SiC-0.5Al-0.6B-2C. The accompanying EDS were taken from the triple junction and bulk grains. Bulk Grains Grain Boundary Figure 5-7. STEM of a grain boundary in -SiC-0.5Al-0.6B-2C. The accompanying EDS were taken from grain boundary and bulk grains.

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58 Figure 5-8. EFTEM of a grain boundary in -SiC-0.5Al-0.6B-2C. Variation in composition across the grain boundary is non-existent. B-C rich grains Metallic Impurities Figure 5-9. STEM of a multiple grain junction in -SiC-0.5Al-0.6B-2C containing metallic impurities as shown by the corresponding EDS. B-C rich grains are also shown. Based on the value of solid solubility of aluminum in SiC reported by Tajima and Kingery,107 formation of secondary phases either in the triple junction or grain boundary

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59 of this material should be non-existent since the amount of aluminum used is close to the solubility limit. However, smaller solubility value has been reported for aluminum in SiC. Kinoshita et al.108 have reported that solubility of aluminum in SiC is ~0.2 wt. % Al based on their work in -SiC with Al2O3 addition. The solubility limit was calculated based on the change in d-spacing in SiC which was assumed to be proportional to the amount of Al. Once the solubility limit was reached, then the d-spacing based on the x-ray diffraction results should be constant. Constant d-spacing values was reached at ~0.4 wt. % Al2O3 and the corresponding Al amount is ~0.2 wt. % Al. Thus, based on this value of solubility limit, which was deemed more appropriate for the present case, -SiC-0.5Al-0.6B-2C can have secondary phases in the triple junction and grain boundary. The formation of secondary phase in the triple junction was observed but a secondary phase in grain boundary was not. This suggests, as mentioned earlier, that solubility limit alone cannot explain the formation of intergranular film in the grain boundary. -SiC-1Al-0.6B-2C This material has an aluminum additive of amount well above the solubility limits of aluminum in SiC and is expected to have secondary phase in the triple junction and grain boundary. As shown in Figures 5-10 and 5-11, the triple junction of this material contains Al-O rich composition. Based on the elemental map shown in Figure 5-11, aluminum and oxygen segregates in the triple junction. Corresponding depletion of silicon and carbon are also shown. The appearance of the triple junctions in this material is different than those observed for -SiC-0.5Al-0.6B-2C. In addition, the Al and O content of the triple junction is greater. This observation supports the earlier assertion that the amount of secondary phases in the triple junction and grain boundary will vary with the amount of aluminum additive.

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60 Tri p le Junction Grain Boundar y Bulk Grains Figure 5-10. STEM of a triple junction in -SiC-1Al-0.6B-2C. Accompanying EDS was taken from bulk grains, grain boundary and triple junction. Figure 5-11. EFTEM of a triple junction in -SiC-1Al-0.6B-2C. Segregation of Al and O in the triple junction is observed.

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61 The grain boundary of this material, however, does not show the expected formation of intergranular phase. As shown in Figure 5-10, the grain boundary EDS shows the presence of Al and O, however, the amount present is comparable to the amount observed in the bulk. Further evidence of non-formation of secondary phase in the grain boundary is shown in Figure 5-12. The EFTEM elemental map does not show any variation in composition across the grain boundary that may suggests possible segregation of secondary phase. Figure 5-12. EFTEM of a grain boundary in -SiC-1Al-0.6B-2C. Compositional variation across the grain boundary is non-existent. -SiC 1.5Al 0.6B 2C As discussed in the grain boundary width and crystallinity section of this chapter, this material shows the presence of amorphous intergranular film in the grain boundary. It was also shown previously that an increase in toughness and change in fracture mode from transgranular to intergranular occurred in this specimen. Since the amount of

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62 aluminum in this specimen is well in excess of the solubility limit of aluminum in SiC, it is expected that formation of secondary phases in grain boundary and triple junction should occur. As shown in Figures 5-13 to 5-16, formation of secondary phases in triple junction and grain boundary was observed in this material. Figure 5-13 and Figure 5-14 shows that the triple junction is filled with Al and O rich composition. In Figure 5-14, the elemental maps of aluminum and oxygen show definite segregation of these elements in the triple junction. Corresponding depletion of carbon and silicon was also observed. The boron map, however, only indicates that boron is almost undetectable in the grain and the triple junction region. The amount of Al and O in the triple junction and in the grain boundary in this material is considerably greater than those observed in the -SiC specimen discussed previously. The Al and O-rich composition in the triple junction is believed to be in the form of Al2O3, as will be shown later. Grain Boundary Bulk Grain Triple Junction Figure 5-13. STEM of a triple junction in -SiC-1.5Al-0.6B-2C. The accompanying EDS was taken from the triple junction, grain boundary and the bulk grain.

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63 Figure 5-14. EFTEM of a triple junction in -SiC-1.5Al-0.6B-2C. Elemental maps indicate segregation of aluminum and oxygen and depletion of silicon and carbon at the triple junction. Bulk Grains Grain Boundary Figure 5-15. STEM of a grain boundary in -SiC-1.5Al-0.6B-2C. The accompanying EDS was taken from the grain boundary and bulk grains. Definite aluminum and oxygen segregation in the grain boundary is shown.

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64 Figure 5-16. EFTEM of a grain boundary in -SiC-1.5Al-0.6B-2C. Elemental maps shows segregation of aluminum and oxygen in the grain boundary. As shown in Figure 5-15 and 5-16, the grain boundary of this material definitely contains Al and O. Segregation of these elements in the grain boundary is readily apparent in the elemental maps presented in Figure 5-16. Thus, based on the results of EDS and EFTEM performed on the grain boundary, the amorphous intergranular film shown in Figure 5-2d and 5-2e is due to the presence of Al and O, which is probably in the form of Al2O3. Other secondary phases were also observed in this material. Aside from the common residual carbon and metallic impurities that segregate primarily to triple junctions, presence of Al-O rich grains was also observed. This is shown in Figure 5-17 where the EDS of the lower grain indicate the presence of Al and O in that particular grain. This phenomenon was not observed in the previous sample studied.

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65 Al-O rich grain Grain Boundary SiC grain Figure 5-17. STEM of a grain boundary and associated bulk grains in -SiC-1.5Al-0.6B-2C. EDS taken at the lower grain shows presence of Al-O rich grain not found in previous samples. -SiC 4Al 0.6B 2C The triple junction and grain boundary of this material also contains secondary phases rich in Al and O. This observation is shown in Figures 5-18 to 5-20. Aside from the usual metallic impurities and residual carbon, other secondary phases present in triple junction and multiple grain junctions were also observed. Some of these secondary phases are shown in Figure 5-18, 5-20 and 5-21. Residual carbon, which was common in all the materials investigated, is shown in Figure 5-18. Figure 5-20 shows a multiple grain junction rich in Al, B and C, whereas Figure 5-21 shows Al-C and B-C rich grains found in the material. The presence of large number of secondary phases in this material is not unexpected. The amount of aluminum in this material is quite high, 4 wt. % Al, and as discussed in the previous section, the amount of secondary phases will be proportional to the amount of aluminum additive.

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66 Bulk Grain Grain Boundary Triple junction Carbon Figure 5-18. Residual carbon in triple junction of -SiC-4Al-0.6B-2C. Figure 5-19. EFTEM of a grain boundary in -SiC-4Al-0.6B-2C. Elemental maps confirms segregation of Al and O in grain boundary.

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67 Figure 5-20. EFTEM of a multiple grain junction in -SiC-4Al-0.6B-2C. Box shows region rich in Al and C. Al-C rich g rain B-C rich g rain Figure 5-21. STEM and EDS of secondary phases found in -SiC-4Al-0.6B-2C. The top two EDS identifies the Al-C and B-C rich grain.

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68 -SiC Sintered with Al Additive Processing and Polytypes Table 5-4 summarizes the density, chemical analysis and phase assemblage of the -SiC sintered with aluminum addition. -SiC-1.65Al-0.5B4C-2C was included in this group to determine the effect of B4C and C. As shown, the fabricated ceramics sintered to a high density. It is apparent that an increase in the Al additive, (comparison of -SiC-1.65Al/2200C and -SiC-3.3Al/2200C), leads to a decrease in the achievable density. Increase in the hot pressing temperature, at the same aluminum additive level, also results in a decrease of the density of the final ceramics. Addition of B4C and C results in a much lower density as compared to those materials with Al addition only. The density measured, however, is still about 98% of the theoretical density of SiC. Table 5-4. Density, chemical analysis and phase assemblage of -SiC sintered with Al additive. Sample (wt.%) Density (g/cc) Chemical Analysis (wt. %) Phase Assemblage (wt. %) Oxygen Nitrogen 3C 2H 4H 6H 15R -SiC-1.65Al, 1900C 3.22 1.01 0.006 0 0 8.1 83.5 8.5 -SiC-1.65Al, 2100C 3.21 0.76 0.008 0 0 61.6 33.0 5.4 -SiC-1.65Al, 2200C 3.21 0.46 0.006 -SiC-3.3Al, 2000C 3.19 0.80 0.01 0 0 12.7 79.7 5.9 -SiC-3.3Al, 2200C 3.18 0.27 0.007 -SiC-1.65Al-0.5B4C-2C, 2100C 3.16 0.98 0.008 0 0 21.2 71.7 7.1 The results of the chemical analysis performed on the fabricated ceramics may not be an exact representative of the bulk due to the small sample volume used, less than 1

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69 gram of material. The results, however, are still valid considering that much smaller volume are being used for grain boundary and triple junction characterization. In order to perform a meaningful comparison, the amount of oxygen on the starting powders used in processing these materials should be considered. The amount of oxygen in the starting -SiC was analyzed to have 1.2 wt. % oxygen, whereas the Al starting powder was measured at 0.76 wt. % oxygen. As shown in Table 5-4, the amount of oxygen retained in the fabricated specimens varies as a function of amount of aluminum added and the hot pressing temperature used. The retained oxygen decreases as a function of hot pressing temperature, as observed for -SiC with 1.65 wt. % Al additive. Doubling the amount of aluminum additive at the same temperature, as shown by comparing -SiC-1.65Al/2200C and -SiC-3.3Al/2200C, results in a decreased retained oxygen in the system. In the case of -SiC-1.65Al-0.5B4C-2C, the amount of oxygen retained is higher than what was retained in a material with the same aluminum content, -SiC 1.65Al. As will be shown later, addition of B4C and C effectively removes the oxygen in the system. However, whenever Al is present, either in the form of Al or AlN, considerable amount of oxygen is retained. The amount of retained oxygen in the fabricated ceramics may have an important implication on the densification of these materials. The decreasing amount of retained oxygen with increasing hot pressing temperature suggests that closed porosity was not obtained until the hold temperature was reached. The presence of porosity provides an escape route for the oxygen vapor generated during processing. Since heating at a higher temperature will take more time than heating at lower temperature, oxygen losses will be higher for materials hot pressed at a higher temperature. Higher retained oxygen content

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70 when B4C and C were added with aluminum also implies a change in the densification kinetics in these materials. It seems that the addition of B4C and C speed up the densification rate and the material achieved a state of closed porosity earlier. The starting -SiC powders used in sintering the materials studied are primarily of 6H polytype. The use of aluminum additive in sintering SiC is known to promote the 6H to 4H polytype transformation.2 As shown in Table 5-4, the 6H to 4H transformation occurred in the materials fabricated. Comparison of the phase assemblage of -SiC-1.65Al hot pressed at 1900C and 2100C also indicates an increase in the amount of phase transformation as a function of temperature. Microstructure and Mechanical Properties The resulting microstructure of the hot-pressed -SiC with aluminum additions are shown in Figure 5-22. In the case of -SiC with 1.65 wt. % Al additive, an increase in the hot pressing temperature results in a change in grain morphology and grain size distribution. Hot pressing at 1900C results in a near-equiaxed grains with an average grain size of ~ 1 m and an aspect ratio range of 1 to 2. Since the average grain size of the starting powder is about 0.5 m, it is apparent that a grain growth had occurred during processing. The small aspect ratio also indicates that anisotropic grain growth, as illustrated by the grain elongation, occurred in a very small degree. Hot pressing at 2100C resulted in distinctly different grain morphology. Elongation of some of the grains occurred while some of the grains remained equiaxed. In effect, a bimodal grain morphology consisting of large elongated grains in a matrix of small equiaxed grains was achieved at this hot pressing temperature. The small equiaxed grains average size is about 1 m and the aspect ratio of the elongated grains at this temperature ranges from 2

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71 to 5. Hot pressing at 2200C did not dramatically change the microstructure observed. Large elongated grains in a matrix of small equiaxed grains are still apparent, however, the aspect ratio of the elongated grains increased. The small equiaxed grains average size is still about 1 m but the aspect ratio range increases to about 3 to 6. Figure 5-22. Transmission electron micrographs of the microstructures observed in -SiC sintered with Al additive. The magnifications used are not the same due to the differences in grain sizes observed. Increasing the amount of Al additive in the material did not drastically change the resulting microstructure. As shown in Figure 5-22, the microstructure of -SiC with 3.3 wt. % Al additive hot pressed at 2000C is similar to the microstructure observed in -SiC-1.65Al hot pressed at 1900C. The only difference is the slightly bigger grain sizes observed in the former material. The bigger grain size, however, is more suitably correlated with the higher temperature. A more suitable comparison to determine the effect of increased aluminum additive would be that of -SiC-1.65Al and -SiC-3.3Al

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72 hot pressed at the same temperature of 2200C. In both materials, the microstructure shows large elongated grains in a matrix of fine equiaxed grains. Although the particular TEM of the microstructure shown for -SiC-3.3Al hot pressed at 2200C consisted of larger volume of elongated grains, on the average, the volume of elongated grains in both materials is about the same. It is evident, however, that the aspect ratio of some of the elongated grains in -SiC-3.3Al is larger than the aspect ratio observed for the elongated grains in -SiC-1.65Al. The effect of B4C and C addition in the resulting microstructure of -SiC sintered with Al is readily apparent. As shown in Figure 5-22, the microstructure of -SiC-1.65Al-0.5B4C-2C exhibits large grains indicating a high grain growth in the material. The marker bar for the transmission electron micrograph of the microstructure of this material is set at 5 m. Average grain size is about 5 m and the aspect ratio of the elongated grains is about 1 to 2. As anticipated, the addition of boron and carbon promoted normal grain growth in -SiC ceramics. The measured hardness, toughness and observed fracture mode from the materials sintered with -SiC with aluminum additions are shown in Table 5-5. The fracture modes of the materials studied are gleaned from the appearance of the Vickers indentation and fracture surfaces which was shown in Figure 5-23 and 5-24, respectively. The observed fracture mode in -SiC sintered with Al addition remains mixed independent of hot pressing temperature and the amount of aluminum added. Direct correlation between the observed microstructure and measured hardness and toughness exists for -SiC sintered with aluminum additives. In the case of -SiC with 1.65 Al additions, hot pressing at 1900C resulted in the highest hardness measured and

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73 the lowest toughness value. The high hardness can be attributed to the grain size and morphology observed in this material, i.e., equiaxed grains with an average grain size of about 1 m which was the smallest grain observed for -SiC sintered with Al additive. Increasing the hot pressing temperature resulted in a decreased hardness and increased toughness in the resulting material. As discussed above, increasing the hot pressing temperature in -SiC-1.65Al also resulted in a change in microstructure from small equiaxed morphology to a self-reinforced microstructure consisting of large elongated grains in a matrix of small equiaxed grains. Table 5-5. Mechanical properties of -SiC sintered with Al additive. Sample (wt. %) Hardness (HV1, GPa) Toughness (MPa m1/2) Fracture Mode -SiC-1.65Al, 1900C 25.30.7 4.00.2 Mixed -SiC-1.65Al, 2100C 22.10.7 4.70.4 Mixed -SiC-1.65Al, 2200C 20.80.3 5.70.1 Mixed -SiC-3.3Al, 2000C 22.10.6 4.20.1 Mixed -SiC-3.3Al, 2200C 20.50.5 6.80.1 Mixed -SiC-1.65Al-0.5B4C-2C 20.30.3 3.10.1 Transgranular Figure 5-23. Optical micrographs of Vickers indentation in -SiC sintered with aluminum additions. Marker bars shown are 10 m.

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74 Figure 5-24. SEM of fracture surfaces of -SiC sintered with aluminum additions. The fracture mode remains mixed for all materials. The decreasing hardness as a function of increasing hot pressing temperature in -SiC-1.65Al seems to correlate with the anisotropic grain growths which have resulted in a larger average grain size at higher processing temperature. This observation is in stark contrast with the ABC-SiC system presented earlier, SiC-YAG and solid-state SiC which has shown little or no hardness dependence on grain size.9 The differences in hardness dependence in grain size may be due to the fracture mode, which remains mixed in all processing temperature, observed in -SiC-1.65Al. As suggested by Rice et al,10 hardness dependence on grain size is possible for ceramics fracturing intergranularly when the grain size is sufficiently smaller than the Vickers indent size. As shown in Figure 5-23, this condition may be true for -SiC-1.65Al ceramics at the hot pressing temperatures used. Another possibility on why the hardness decreases with increased in grain size may be due to the trend in ceramic materials wherein the inherent flaw sizes present scales up with the grain size. Since the mechanical properties in ceramics are highly dependent on flaw sizes in the material, it is not surprising that hardness decreases

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75 with increasing grain size since increasing grain size would imply an increase in the inherent flaw size present. As shown in Table 5-5, the toughness measured in -SiC-1.65Al also correlates with the hot pressing temperature and the resulting microstructure. Increasing toughness values were measured at increasing hot pressing temperature. As discussed previously, an increase in the aspect ratio of the elongated grains in the materials processed at higher temperature was also observed. Literature report indicates that the presence of key elongated grains in the microstructure greatly affects the toughness of toughened ceramics.11-13 As observed by Lee et al13 an increase in aspect ratio and volume of key elongated grain results in an increase toughness in silicon carbide possessing a self-reinforced microstructure. Although increased in volume of elongated grains in the -SiC-1.65Al was not observed, increased in the aspect ratio of elongated grains were. Thus, the resulting increased in aspect ratio of the elongated grains at higher processing temperature is deemed primarily responsible for the toughening observed for -SiC sintered with Al additive. This also applies for the observed toughness value in -SiC-3.3Al processed at 2200C since the resulting microstructure in this material also showed elongated grains of aspect ratio higher than those observed for -SiC-1.65Al processed at the same temperature. As shown in Table 5-5, B4C and C additions in -SiC-1.65Al resulted in low hardness and toughness values. This observation was primarily due to the microstructure achieved in this material. As shown, the grain size is relatively large and the grains are mostly equiaxed. Presence of elongated grains with high aspect ratio in this material was

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76 not observed thus the toughening mechanisms responsible for high toughness in -SiC with pure aluminum additions are not operable in this material. Grain Boundaries, Triple Junctions and Secondary Phases -SiC 1.65 wt. % Al hot pressed at 1900C Figure 5-25 shows several grain boundaries studied in -SiC-1.65Al hot pressed at 1900C. As shown, presence of amorphous intergranular film in the grain boundary of this material is apparent. However, some of the grain boundaries observed, as shown in Figure 5-25d, also show absence of intergranular film. The amount of aluminum additive used in this material is well in excess of the solubility limit of aluminum in SiC. The amount is even higher than the amount of aluminum used in -SiC where intergranular amorphous film in the grain boundaries were observed. Thus, it is expected that formation of intergranular film in the grain boundary of this material should occur. a) b) d) c) Figure 5-25. High resolution transmission electron micrographs of several grain boundaries studied in -SiC-1.65Al hot pressed at 1900C. a), b), and c) Shows amorphous grain boundary with a width of about 1 nm. d) Grain boundary with no amorphous intergranular film.

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77 The composition of grain boundary and triple junction phases in this material are shown in Figures 5-26, 5-27 and 5-28. The grain boundary, as typified by Figure 5-26, is generally filled with Al and O rich phase. The Al and O concentration in the grain boundary, as shown by the EDS data, are generally higher than those found in the nearby grains. This indicates that definite segregation of Al and O in the grain boundary really occurred. As shown in Figure 5-27 and 5-28, segregation of Al and O in the triple junctions of this material was also observed. Electron energy loss spectra (EELS) were taken from the triple junction and an example of which is shown in Figure 5-19. The Al and O peaks were then compared with literature data for crystalline Al2O3 and found to be consistent. The crystallinity of the triple junction is further ascertained via high resolution lattice imaging of the triple junction shown in Figure 5-30. And as shown, lattice fringes consistent with crystalline structure were observed. Bulk Grain Grain Edge Grain Boundary Figure 5-26. STEM and EDS of a grain boundary and associated grains in -SiC-1.65Al hot-pressed at 1900C.

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78 Bulk Grain Triple Junction Triple Junction Figure 5-27. STEM and EDS of a triple junction in -SiC-1.65Al hot-pressed at 1900C. The triple junction phases observed are generally Al and O rich. Figure 5-28. EFTEM of a triple junction and grain boundary in -SiC-1.65Al hot pressed at 1900C. Al and O, in the form of Al2O3, segregates at the triple junction.

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79 O K-peak in Al2O3 Al K-peak in Al2O3 Figure 5-29. EELS taken from a triple junction in -SiC-1.65Al hot pressed at 1900C. The Al and O peaks are consistent with crystalline Al2O3. b) a) O Al c) Figure 5-30. High resolution images and EDS of triple junctions in -SiC-1.65Al hot pressed at 1900C. a) and b) are high resolution images, c) is an EDS taken from the triple junction. The appearance of the triple junction phase in the manufactured ceramics may yield information about the process in which the liquid phase distributes itself to cover the surfaces of the particulate solids. By measuring the dihedral angle, based on STEM images and EFTEM elemental maps, one can determine if complete penetration of the triple junction and grain boundary by the liquid phase occurred. In the case of -SiC-1.65Al hot pressed at 1900C, the measured dihedral angles range from 15 to 60. This implies that a continuous liquid phase penetration of all triple grain junctions and partial penetration of the grain boundary occurred in this material. This also implies that the composition of the triple junction phase and the grain boundary phase (or the intergranular amorphous film) is the same, i.e., Al2O3.

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80 -SiC 1.65 wt. % Al hot pressed at 2100C Figure 5-31 shows high resolution micrographs of several grain boundaries observed in -SiC-1.65Al hot pressed at 2100C. As shown, grain boundaries with distinct grain transitions which indicates absence of intergranular phase, and grain boundary with amorphous intergranular film are also present in this material. For grain boundaries with amorphous intergranular film, the width of the film observed is typically about 1 nm. The composition of the amorphous intergranular film or the secondary phase in the grain boundary is generally Al and O. As shown in Figure 5-32, segregation of Al and O in the grain boundary is generally observed. It thus seem conflicting that the high resolution images of the grain boundary shows absence of intergranular films while the EDS and EFTEM elemental maps suggests that segregation of Al and O occurs which implies that it should be present as an intergranular film. d) c) b) a) Figure 5-31. HRTEM of several grain boundaries in -SiC-1.65Al hot pressed at 2100C. a), b), and c) are grain boundaries with no intergranular phase. d) shows a grain boundary with amorphous intergranular film and a crystalline triple junction.

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81 Figure 5-32. EFTEM of a grain boundary in -SiC-1.65Al hot pressed at 2100C. Elemental maps shows segregation of Al and O in the grain boundary. The triple junction phase composition observed in -SiC-1.65Al hot pressed at 2100C is shown in Figures 5-33, 5-34 and 5-35. In all these figures, the composition of the triple junction phase was shown to be Al and O. EELS of the triple junction also revealed that the phase present is in the form of crystalline Al2O3. The three figures also show different appearances of the triple junction phases. Measured dihedral angle based on the STEM images and EFTEM of triple junction phases in -SiC-1.65Al hot pressed at 2100C ranges from 15 to 90. Based on this range of dihedral angles, presence of liquid phase penetrating all three grain junctions, and liquid phase that only partially penetrates the three grain junctions, are possible. In addition, partial liquid phase penetration of the grain boundary would still occur. This explains the observed absence of secondary phase in some of the grain boundaries investigated.

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82 Bulk Grain Grain Boundary Triple Junction Triple Junction Figure 5-33. STEM and EDS of triple junction in -SiC-1.65Al hot pressed at 2100C. Grain Boundary Bulk Grain Triple Junction Figure 5-34. STEM of triple grain junction phase with lower dihedral angle in -SiC-1.65Al hot pressed at 2100C.

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83 Figure 5-35. EFTEM of a triple junction in -SiC-1.65Al hot pressed at 2100C. -SiC 1.65 wt. % Al hot pressed at 2200C The high resolution images of the grain boundary in this material, as shown in Figure 5-36, show presence of clean grain boundary and grain boundary with amorphous intergranular films. The thickness of amorphous intergranular film (IGF) in this material is about 1 nm, similar to the IGF observed in specimen with the same composition hot pressed at lower temperatures. EDS and EFTEM elemental maps taken from the grain boundary of -SiC-1.65Al hot pressed at 2200C showed that the grain boundary phase is Al and O rich and is similar to the previous materials studied. The triple junction phase composition, as shown in Figures 5-37 and 5-38, is also consistent with that of Al2O3. The triple junction phase, based on high resolution imaging, is also crystalline. The grain

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84 boundary and triple junction phases are deemed similar, i.e., both are Al2O3, based on the dihedral angles measured which ranges from 40 to 90. a) c) b) Figure 5-36. HRTEM of grain boundaries in -SiC-1.65Al hot pressed at 2200C. Bulk Grain Grain Boundar y Tri p le Junction Tri p le Junction Figure 5-37. STEM of a triple junction in -SiC-1.65Al hot pressed at 2200C. EDS shown were taken from triple junction, grain boundary and bulk grains.

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85 Figure 5-38. EFTEM of a triple junction and grain boundary in -SiC-1.65Al hot pressed at 2200C. -SiC 3.3 wt. % Al hot pressed at 2000C Three distinct grain boundary appearances have been observed in this material. As shown in Figure 5-39, presence of a clean grain boundary, a grain boundary with less than 1 nm amorphous IGF, and a grain boundary with about 3 nm amorphous IGF were observed. Based on the EDS taken from several grain boundary of this material, the grain boundary phase exhibiting a thickness of less 1 nm is believed to be Al and O rich. The amorphous IGF with thickness of about 3nm, however, is primarily aluminum with some oxygen incorporation. The relative amount of Al and O in this thick IGF is not consistent with Al2O3. The differences in the grain boundary phases composition can be seen by comparing the EDS taken from the grain boundary in Figure 5-40 and Figure 5-41 showing the aluminum-rich and normal Al-O rich composition, respectively.

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86 c) d) b) a) Figure 5-39. HRTEM of grain boundaries in -SiC-3.3Al hot pressed at 2000C. a) and b) shows a clean grain boundary. c) Shows an amorphous IGF of thickness less than 1 nm, and d) Shows a grain boundary with amorphous IGF of thickness of about 3 nm. Grain Boundary Bulk Grain Triple Junction Figure 5-40. STEM of a grain boundary near a triple junction in -SiC-3.3Al hot pressed at 2000C. EDS were taken from triple junction, grain boundary and bulk grain. Triple junction shows presence of metallic impurities. Grain boundary composition is primarily Al-rich.

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87 Bulk Grains Grain Boundary Al-rich TJ phase Metallic Impurities Figure 5-41. STEM of a triple junction in -SiC-3.3Al hot pressed at 2000C. EDS were taken from triple junction edge, triple junction, grain boundary and bulk grains. Triple junction phase is primarily Al-rich. Grain boundary phase is shown to be the normally observed Al-O-rich phase. It seems that the increase in the amount of the aluminum additive used has a significant effect on the grain boundary and secondary phases present in the material. As discussed above, presence of a primarily Al-rich grain boundary phase not observed previously was determined. As shown in Figure 5-40 and 5-41, the appearance of metallic impurities also differed from previous materials studied. In -SiC-3.3Al hot pressed at 2000C, the number of triple junction sites containing metallic impurities seemed to increase. The volume of the metallic impurities, however, is similar to the previous materials studied. Another difference observed was the present of Al-rich phases, existing in the triple junction, as shown in Figure 5-41, and existing as a grain,

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88 shown in Figure 5-42. Primarily Al-rich secondary phases were not observed in previous materials studied. Normal SiC grain Al-rich grain Figure 5-42. STEM of two grains with different composition found in -SiC-3.3Al hot pressed at 2000C. Although presence of primarily Al-rich triple junction phase are numerous in this material, presence of the normal Al2O3 in the triple junction was still observed, as shown in Figure 5-43. Based on the elemental map, the triple junction phase in this particular triple junction is still Al2O3. The grain boundary phase, however, based on the comparison of Al and O elemental maps are primarily Al-rich. The presence of primarily Al-rich phase can be attributed to the larger amount of aluminum additive and the dihedral angle of the liquid phases that formed. Based on several EDS and EFTEM elemental maps, the dihedral angle in this material ranges from 40 to 120. Dihedral angle of 40 to 60 implies that continuous liquid phases penetration of all three-grain junctions and partial penetration of the grain boundary are possible.[14] This is of course assuming that enough volume of liquid phases is available. Since complete penetration of triple junction in -SiC-1.65Al was observed, it is safe to assume that enough liquid

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89 phases are available in -SiC-3.3Al. Thus, for a dihedral angle of 40 to 60, grain boundary with Al2O3 phase is possible and the triple junction phase can be continuous within the grain boundary. In the case of dihedral ranging from 60 to 120, isolated liquid phases partial penetration of triple grain junction is possible, but the liquid phase will not penetrate the grain boundary. This implies that if another liquid phase is present with lower dihedral angle, that liquid phase may penetrate the grain boundary. In Figure 5-43 the dihedral angle is measured to be 120. As observed, the triple junction phase is not continuous into the grain boundary. The fact, however, that aluminum segregation in the grain boundary was observed, as shown by the Al elemental map, suggests that the dihedral angle of liquid Al phase in SiC grains is lower. Figure 5-43. EFTEM of a triple junction in -SiC-3.3Al hot pressed at 2000C. Comparison of the Al and O elemental maps indicate that the grain boundary phase is primarily Al-rich.

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90 -SiC 3.3 wt. % Al hot pressed at 2200C The grain boundaries observed in this material are shown in Figure 5-44. Presence of clean grain boundary and grain boundary with amorphous IGF was observed. As shown, the amorphous intergranular film thickness observed is about 1 nm. The composition of the grain boundary phase is generally Al-O-rich, as shown in Figure 5-45 and 5-46. The triple junction phase and appearance, however, was the same as those observed from the material hot pressed at 2000C, as shown in Figure 5-47. Al-rich secondary phase in the triple junction, Figure 5-47, and in the form of a grain, as shown in Figure 5-48, were also observed. Isolated triple junction phase, containing Al2O3, with dihedral angle range similar to previous material, are also present, as shown in Figure 5-49. The presence of primarily Al-rich phase in the grain boundary was not observed in this material through EDS but was seen through EFTEM as shown in Figure 5-49. d) c) b) a) Figure 5-44. HRTEM of grain boundaries in -SiC-3.3Al hot pressed at 2200C. a) and b) shows amorphous intergranular film of about 1 nm thickness. c) and d) shows clean grain boundaries.

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91 Grain Boundar y Bulk Grain Figure 5-45. STEM of a grain boundary in -SiC-3.3Al hot pressed at 2200C. Figure 5-46. EFTEM of a grain boundary in -SiC-3.3Al hot pressed at 2200C. Elemental maps shows segregation of Al and O.

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92 Triple Junction Bulk Grain Triple Junction edge Grain Boundary Figure 5-47. STEM and EDS of a triple junction containing primarily Al-rich phase and metallic impurities in -SiC-3.3Al hot pressed at 2200C. The grain boundary phase is consistent with the normal Al-O enrichment. The triple junction shows the presence of primarily Al-rich phase. SiC Grain Al-rich Grain Figure 5-48. STEM and EDS of Al-rich grain found in -SiC-3.3Al hot pressed at 2200C.

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93 Figure 5-49. EFTEM of a triple junction in -SiC-3.3Al hot pressed at 2200C. Segregation of aluminum in the grain boundary is seen in the Al elemental map. -SiC 1.65 wt. % Al 0.5 wt. % B4C 2 wt. % C High resolution images of the grain boundaries in this material revealed presence of clean grain boundary and grain boundary with amorphous intergranular film. As shown in Figure 5-50, the thickness of the amorphous intergranular film in this material varies from 2 nm to about 4 nm. The observed amorphous films, in these particular high resolution images, however, are primarily carbon. Presence of segregation of Al, O and B in the grain boundary of this material is also evident, as shown in Figure 5-51. Segregation of Al and O in the grain boundary was also observed, as shown in Figure 5-52. The absence of B in the grain boundary shown in Figure 5-52, however, may just be due to the inability of EDS to detect boron at relatively low amount.

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94 d) c) b) a) Figure 5-50. HRTEM of grain boundaries observed in -SiC-1.65Al-0.5B4C-2C. a) and b) shows clean grain boundaries observed. c) and d) shows amorphous IGF of thickness ranging from 2 to 4 nm. Figure 5-51. EFTEM of a grain boundary in -SiC-1.65Al-0.5B4C-2C. Segregation of Al, O and B in the grain boundary is shown.

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95 Grain Boundary Bulk Grain Triple Junction Figure 5-52. STEM of a triple junction in -SiC-1.65Al-0.5B4C-2C. EDS were taken from triple junction, grain boundary and bulk grains. The triple junction phase composition observed in -SiC-1.65Al-0.5B4C-2C was also shown in Figure 5-52. Based on the EDS data taken from most of the triple junction, the triple junction phase are generally Al and O rich. Results from EFTEM elemental maps, however, also show the presence of boron together with the observed Al and O. As gleaned from the grain boundary results, it is highly possible that B was not detected through EDS and the triple junction phase composition is really Al, O and B rich. Aside from the usual metallic impurities present, other secondary phases were observed in this material. As shown in Figure 5-53, presence of residual carbon and boron-carbon rich phase were detected. The residual carbons are usually found in the triple/multiple grain junctions and in some cases, in between two SiC grains. The B-C rich phases usually exist as a definitive grain and are also found in triple/multiple grain junctions. The size

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96 of these secondary phases, residual carbon and B-C-rich grains, may have a considerable effect on the mechanical property of this material. Bulk Grain Residual carbon B-C rich grain Figure 5-53. STEM and EDS of other secondary phases in -SiC-1.65Al-0.5B4C-2C. -SiC Sintered with AlN Additive Processing and Polytypes Table 5-6 summarizes the density, chemical analysis, and phase assemblage of the -SiC sintered with AlN additive. -SiC-2.5AlN-0.5B4C and -SiC-2.5AlN-0.5B4C-2C are included in this set of sample to determine the effect of B and C addition. Most of the -SiC with AlN additive, except for material with further addition of B4C and C, sintered to a high density. The density values seem to have little dependence on the hot pressing temperature and the amount of additive used. The results of the chemical analysis performed on the materials studied suggest that the amount of retained oxygen and nitrogen are dependent on the hot pressing temperature and the amount of AlN. At 2.5 wt. % AlN addition, the retained oxygen and

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97 nitrogen increases with increasing hot pressing temperature. Hot pressing at 2100C resulted in a retained oxygen and nitrogen of 0.44 and 0.2 wt. %, respectively, while hot pressing at 2200C resulted in a retained oxygen and nitrogen of 0.64 and 0.42 wt. %, respectively. Higher AlN addition resulted in a reversed trend. In the case of -SiC-5AlN, increasing the hot pressing temperature from 2000 to 2200C resulted in a decreased oxygen and nitrogen content from 0.72 to 0.49 and 1.45 to 0.89 wt. %, respectively. Addition of B4C in -SiC-2.5AlN lowers the amount of retained oxygen but increases the nitrogen to twice the amount observed for pure AlN addition. Simultaneous addition of B4C and C in -SiC-2.5AlN increased the amount of oxygen but lowers the nitrogen to an amount still higher than what pure AlN additions have achieved. Table 5-6. Density, chemical analysis and phase assemblage of -SiC sintered with AlN additive. Sample (wt.%) Density (g/cc) Chemical Analysis (wt. %) Phase Assemblage (wt. %) Oxygen Nitrogen 3C 2H 4H 6H 15R -SiC-2.5AlN, 2100C 3.22 0.44 0.20 0 0 26.8 62.2 7.7 -SiC-2.5AlN, 2200C 3.21 0.64 0.42 0 0 39.7 53.0 7.2 -SiC-5AlN, 2000C 3.19 0.72 1.45 0 0 0.8 89.5 9.7 -SiC-5AlN, 2200C 3.22 0.49 0.89 0 0 6.6 82.5 10.9 -SiC-2.5AlN-0.5B4C 3.20 0.38 0.84 0 0 52.3 44.0 3.7 -SiC-2.5AlN-0.5B4C-2C 3.16 0.72 0.59 0 0 49.4 45.7 4.9 It has been reported that formation of a solid solution between SiC and AlN exists.15-18 The reported polytype in compositional range where solid solution was observed is generally that of 2H-SiC. As shown in Table 5-6, addition of AlN in -SiC

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98 did not result in the formation of 2H polytype. The resulting polytypes in the sintered ceramics are still primarily 6H, 4H and 15R. Additions of B4C and C resulted in the formation of greater amount of 4H structure. Microstructure and Mechanical Properties The resulting microstructure of -SiC sintered with AlN additive are shown in Figure 5-54. As expected, increasing the hot pressing temperature resulted in an increased grain sizes in the final structure as a result of normal grain growth. The grain sizes achieved in this material, however, are still smaller than those achieved in -SiC with Al addition processed at the same hot pressing temperature. In the case of -SiC-2.5AlN, average grain sizes of the materials hot pressed at 2100C is about 1.5 m while those hot pressed at 2200C is about 2 m. Increased AlN addition, -SiC-5AlN, resulted in smaller average grain sizes. For -SiC-5AlN hot pressed at 2000C the average grain size is about 0.5 m. Increasing the hot pressing temperature to 2200C resulted in an average grain size of about 1 m only. The use of B4C and C addition also resulted in an increased average grain size. Again, the average grain size did not reach the average achieved when using Al additive together with B4C and C. The grain morphology observed for -SiC sintered with AlN additive remained equiaxed in all the hot pressing temperature and amount of additive used. This is distinctly different than those observed in -SiC sintered with Al additive wherein formation of elongated grains at higher hot pressing temperature was observed. Based on the microstructure observed in -SiC sintered with AlN additive, the effect of nitrogen in the sintering of SiC is patent. Although good densification can be achieved, use of AlN

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99 retards the grain growth in SiC. This effect has been observed in the literature and can be used in the future in controlling the resulting microstructure in SiC ceramics.8, 35-36 Figure 5-54. TEM of the microstructures observed in -SiC sintered with AlN additive. The measured hardness, toughness and observed fracture mode in -SiC sintered with AlN additives are summarized in Table 5-7. The fracture mode was based on the appearance of the Vickers indent on the polished surfaces and fracture surfaces of the studied materials shown in Figure 5-55 and 5-56, respectively. The observed fracture modes in the -SiC sintered with AlN are generally intergranular, except in the case of -SiC-5AlN hot pressed at 2000C where of mixed mode of fracture were observed. This was probably due to the presence of very fine equiaxed grains in the microstructure of this material. Addition of B4C also changes the mode of fracture from intergranular to mixed mode while further addition of C changes the fracture mode to a transgranular one.

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100 Table 5-7. Mechanical properties of -SiC sintered with AlN additive. Sample (wt. %) Hardness (HV1, GPa) Toughness (MPa m1/2) Fracture Mode -SiC-2.5AlN, 2100C 21.90.3 3.50.2 Intergranular -SiC-2.5AlN, 2200C 22.90.1 3.50.4 Intergranular -SiC-5AlN, 2000C 25.00.4 3.90.2 Mixed -SiC-5AlN, 2200C 24.10.6 4.20.1 Intergranular -SiC-2.5AlN-0.5B4C 21.00.6 3.40.4 Mixed -SiC-2.5AlN-0.5B4C-2C 22.70.3 2.80.2 Transgranular Figure 5-55. Optical micrographs of Vickers indentation on a polished surface of -SiC sintered with AlN additive. The marker bar for all micrographs is 10 m. The toughness values measured in -SiC sintered with AlN additive did not show a significant variation as a function of hot pressing temperature or the amount of additive used. The SEPB toughness in these materials is ~4 MPa m1/2, independent of the hot pressing temperature or the amount of additive used. Addition of B4C and C slightly lowers the measured toughness but this may be due to the change in fracture mode observed when these additives are used. As shown earlier, presence of B4C and C changes the fracture mode to a transgranular one.

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101 The amount of AlN added seems to affect the values of hardness measured in -SiC sintered with AlN additive. At high AlN content, -SiC-5AlN, the measured hardness are higher than those measured in -SiC-2.5AlN. This is primarily due to the grain sizes observed in these materials. At higher AlN content, grain growth was retarded resulting in a lower average grain size as compared to low AlN level. As stipulated in the case of -SiC-Al and in general ceramics, the hardness values scales with the inherent flaw sizes which in turn scales with the grain sizes in the material. Figure 5-56. SEM of fracture surfaces of -SiC sintered with AlN additive. Magnification for all images is 2000X. Grain Boundaries, Triple Junctions and Secondary Phases -SiC 2.5 wt. % AlN hot pressed at 2100C Observed grain boundaries in this material are shown in Figure 5-57. High resolution TEM revealed the presence of clean grain boundary and grain boundaries with amorphous intergranular films. The thickness of the amorphous IGF in this material was measured to be less than 1 nm. The composition of the grain boundary phase was found

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102 to be Al and O rich. Nitrogen was not detected in either the grain boundary or the triple junction of this material, as shown in Figure 5-58 and 5-59. d) c) b) a) Figure 5-57. HRTEM of grain boundaries observed in -SiC-2.5AlN hot pressed at 2100C. a) and c) shows amorphous intergranular film with thickness of less than 1 nm. b) and d) are clean grain boundaries observed in this material. Bulk Grain Gr ain B oundar y Metallic Impurities Triple Junction Figure 5-58. STEM and EDS taken from -SiC-2.5AlN hot pressed at 2100C. EDS were taken from triple junction, grain boundary and the bulk grains.

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103 Figure 5-59. EFTEM of a triple junction in -SiC-2.5AlN hot pressed at 2100C. Segregation of nitrogen in the triple junction and grain boundary was not observed. The triple junction phase composition in this material is also Al and O rich, as shown in Figure 5-58 and 5-59. Nitrogen segregation in the triple junction was not detected in either EDS or EFTEM. The triple junction phase was determined via EELS to be Al2O3 and is present in crystalline form, as shown by the high resolution image of the triple junction in Figure 5-60. The dihedral angle measured in this material ranges from 40 to 90. This implies that some of the triple junction phase will be continuous into the grain boundary and some will be isolated pockets in the triple junction. This explains the different grain boundaries shown in Figure 5-57. For dihedral angle of 40 to 60, partial penetration of triple junction phase into the grain boundary is possible thus

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104 the grain boundary shows the amorphous intergranular film. At higher dihedral angle, the triple junction phase is isolated thus the surrounding grain boundaries are clean. b) a) Figure 5-60. Low and high resolution image of a triple junction in -SiC-2.5AlN hot pressed at 2100C. a) low magnification image, b) high magnification image showing lattice fringes in the triple junction phase. -SiC 2.5 wt. % AlN hot pressed at 2200C High resolution images of the grain boundaries observed in -SiC-2.5AlN hot pressed at 2200C are shown in Figure 5-61. Presence of clean grain boundary and grain boundary with crystalline intergranular film were observed in this material. The crystalline intergranular film has a thickness of about 1 nm. As shown by the EDS taken from a grain boundary, Figure 5-62, the grain boundary phase composition is Al and O rich and is possibly in the form of Al2O3. No nitrogen segregation in the grain boundary was observed for this specimen. The triple junction phase composition, as shown in Figure 5-63 and 5-64, is generally Al and O rich. The Al and O are in the form of crystalline Al2O3. Segregation of nitrogen in triple junction was not observed either by EDS or EFTEM. Dihedral angle measured in this material ranges from 15 to 120. Thus it is expected that some of the

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105 triple junction phase will be continuous into the grain boundary. At higher dihedral angle, however, the triple junction phase will not penetrate the grain boundary which explains the presence of clean grain boundary. c) b) a) Figure 5-61. HRTEM of grain boundaries in -SiC-2.5AlN hot pressed at 2200C. a) shows a crystalline intergranular film of about 1 nm thickness. b) and c) shows clean grain boundary. Bulk Grains Grain Boundary Figure 5-62. STEM and EDS of grain boundary in -SiC-2.5AlN hot pressed at 2200C. EDS shows segregation of Al and O in the grain boundary.

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106 Bulk Grains Tri p le Junction ed g e Tri p le Junction Figure 5-63. STEM of a triple junction in -SiC-2.5AlN hot pressed at 2200C. EDS of the triple junction phase shows Al and O in the form of Al2O3. Aluminum Map Oxygen Map Grain Boundary and Tri p l e Ju n ct i o n Silicon Map Oxygen Map Nitrogen Map Figure 5-64. EFTEM of a triple junction in -SiC-2.5AlN hot pressed at 2200C. Elemental maps shows segregation of Al and O in the triple junction.

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107 -SiC 5 wt. % AlN hot pressed at 2000C High resolution TEM of several grain boundaries in -SiC-5AlN hot pressed at 2000C are shown in Figure 5-65. Presence of clean grain boundary and grain boundary with amorphous intergranular film were also observed in this material. As shown in Figure 5-66 and 5-67, the composition of the grain boundary phase are generally Al, O and N rich. Grain boundary with only Al and O were also seen, but this may be due to the inability of EDS to detect nitrogen at relatively low level. d) c) b) a) Figure 5-65. HRTEM of grain boundaries in -SiC-5AlN hot pressed at 2000C. a), b) and c) shows clean grain boundary. d) shows grain boundary with amorphous intergranular film. The triple junction phase composition of this material was also observed to contain Al, O and N, as shown in Figure 5-66 and 5-67. The triple junction shown in Figure 5-66 is not completely filled with secondary phase, as shown by the STEM image and the EDS taken from the triple junction where the Al and O peaks are minimal. It seems that nitrogen is present everywhere, as shown by the EDS taken from the grain boundary, triple junction and the bulk grain. Presence and segregation of nitrogen in the triple

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108 junction and grain boundary was also confirmed by the EFTEM elemental maps, shown in Figure 5-67. The presence of nitrogen is not unexpected. As shown by the chemical analysis performed in this material, retained nitrogen amount is highest in this specimen and may have reached the detection limit even in EDS. Dihedral angle measurements performed on STEM and EFTEM elemental maps yielded a range of 15 to 60. This means that the triple junction phase completely penetrates the triple junction and partial penetration in the grain boundary is also possible. Since grain sizes are small, it is not unexpected that the triple junction phase is continuous into the grain boundary, as observed through EDS and EFTEM. Bulk Grain Grain Boundary Bulk Grain Grain Boundary Triple Junction Figure 5-66. STEM of triple grain junction in -SiC-5AlN hot pressed at 2000C. EDS taken from the triple junction, grain boundary and bulk grain shows presence of nitrogen.

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109 Grain Boundary and Tri p l e Ju n ct i o n Oxygen Map Aluminum Map Silicon Map Carbon Map Nitrogen Map Figure 5-67. EFTEM of triple junctions in -SiC-5AlN hot pressed at 2000C. Triple junction phase composition is continuous into the grain boundary. -SiC 5 wt. % AlN hot pressed at 2200C High resolution TEM of several grain boundaries observed in this material are shown in Figure 5-68. The grain boundaries observed are generally clean, however, abrupt transitions in the lattice fringes of two grains are not generally observed. Results of EDS, as shown in Figure 5-69, suggests that presence of a secondary phase in the grain boundary should exists. As shown in the EDS taken from the grain boundary, segregation of Al and O is present. Based on the appearance of the high resolution images and the EDS results, it is possible that the grain boundary was not properly oriented to the electron beam when the images were taken. Triple junction phase in this material were also found to be Al and O rich, as shown in Figure 5-69 and 5-70. The Al and O are also in the form of crystalline Al2O3.

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110 Dihedral angles measured from STEM and EFTEM images of triple junction in this material ranges from 40 to 90. c) d) b) a) Figure 5-68. HRTEM of grain boundaries in -SiC-5AlN hot pressed at 2200C. a), b), c), and d) shows clean grain boundary with no direct lattice fringe transition. The lattice fringe direction, however, is different in each grain. Bulk Grains Grain Boundar y Tri p le Junction Figure 5-69. STEM of a triple junction in -SiC-5AlN hot pressed at 2200C. EDS were taken from grain boundary, triple junction and bulk grains.

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111 Figure 5-70. EFTEM of a multiple grain junction and triple grain junction in -SiC-5AlN hot pressed at 2200C. Segregation of nitrogen in triple/multiple grain junction was not observed. -SiC 2.5 wt. % AlN 0.5 wt. %B4C HRTEM of several grain boundaries in -SiC-2.5AlN-0.5B4C are shown in Figure 5-71. Most of the grain boundaries observed did not show any indication that secondary phase in the grain boundary are present. As shown, the fringe pattern of the grains did not show abrupt transition at the grain boundary, but distinct lattice fringe representing two grains were observed. The EDS data taken from the grain boundary, as shown in Figure 5-72, indicates that small amount of Al and O in some, and in some cases, only Al was observed in the grain boundary. EFTEM data taken from the grain boundary, as shown in Figure 5-73, confirms that Al and O, together with B, segregates in some of the grain boundary observed. Thus, it is possible that the HRTEM images were just not properly oriented with respect to the electron beam during imaging. The possibility that

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112 what we have observed is really a clean grain boundary cannot be ruled out either. As will be shown later, the dihedral angle measured in this specimen suggests that partial penetration of the grain boundary is possible. Thus, it is highly probable that some grain boundary really did not contain any secondary phase. c) b) a) Figure 5-71. HRTEM of several grain boundaries observed in -SiC-2.5AlN-0.5B4C. a), b), and c) shows grain boundary with no indication of any secondary phase present. The triple junction phase in this material is also Al and O rich, as shown by the EDS taken from the triple junction in Figure 5-72. EFTEM data taken from the triple junction, as shown in Figure 5-74, however, shows that aside from Al and O, B also segregates in the triple junction. Again, the EDS and EFTEM is not in conflict, it is mainly due to the detection limit of EDS that B, in small amount, did not show up in the energy dispersive spectrum taken. Comparison of the EFTEM data taken from the grain boundary and triple junction, shown in Figure 5-73 and 5-74, indicates that the triple junction phase is sometimes continuous into the grain boundary. This is of no surprise since the dihedral angle measured from the specimens ranges from 20 to 80, which

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113 indicates that partial penetration of triple junction phase into the grain boundary is possible. Figure 5-72. STEM of triple junction observed in -SiC-2.5AlN-0.5B4C. EDS of the triple junction and grain boundary shows segregation of Al and O. Figure 5-73. EFTEM of a grain boundary observed in -SiC-2.5AlN-0.5B4C. Elemental maps show segregation of Al, O, and B in the grain boundary. Triple Junction Bulk Grain Grain Boundary

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114 Figure 5-74. EFTEM of a triple junction observed in -SiC-2.5AlN-0.5B4C. Elemental maps show segregation of Al, O, and B in the triple junction. Seveunction phase, were Figure 5-75. STEM of secondary phases observed in -SiC-2.5AlN-0.5B4C. EDS indicates the presence of residual carbon and B-N-rich phase. ral secondary phases, aside from the grain boundary and triple j also observed in -SiC-2.5AlN-0.5B4C. The typical metallic impurities segregation in the triple junction, present in all of the materials studied, were also observed in this specimen. As shown in Figure 5-75, residual carbon and B-N rich phase were also present. Bulk Grains Residual Carbon BN rich p hase

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115 -SiC 2.5 wt. % AlN 0.5 wt. % B4C 2 wt. % C High resolution TEM of several grain boundaries observed in -SiC-2.5AlN-0.5B4C-2C are shown in Figure 5-76. Presence of clean grain boundary and grain boundary with amorphous intergranular films are observed in this material. The thickness of the intergranular film was observed to vary from 1 nm to 10 nm. The compositions of the intergranular film, however, are different for different thicknesses. Grain boundaries with IGF of about 1 nm are observed to be primarily of Al and O, while thickness of about 10 nm is generally of carbon composition. Figure 5-76. HRTEM of several grain boundaries in -SiC-2.5AlN-0.5B4C-2C. a) and b) shows grain boundary with amorphous IGF of about 1 nm. c) shows a clean grain boundary. d) shows a grain boundary with amorphous IGF of about 10 The shows variation in composition. Some of the grain boundary observed showed primarily Al composition, while some of the grain boundary shows presence of Al and O. EFTEM of the grain boundary near a triple junction, as shown in Figure 5-77 and 5-78, illustrate a) b) c) d) nm. Composition of IGF in a) and b) are generally Al and O rich while that in d) is primarily carbon. grain boundary phase composition observed via EDS in this material generally

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116 the observed differences. In Figure 5-77, the grain boundary shows segregation of Al, O, and B. In Figure 5-78, the grain boundary only shows a defined segregation of Al. Figure 5-77. EFTEM of a triple junction and grain boundary in -SiC-2.5AlN-0.5B4C-2C. Elemental map shows segregation of Al, O, and B in the grain boundary and triple junction. Figure 5-78. Another EFTEM taken from a different triple junction in this material. Elemental map shows segregation of Al in the grain boundary. The triple junction phase composition, however, shows presence of Al, O and B.

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117 The triple junction phase composition of this material, as shown in Figure 5-77rimarily Al, O, and B rich, similar to those observed in -SiC-2.5AlN-0. and 5-78, are p5B4C. The don, in Figure 5-79. STEM of several secondary phases observed in -SiC-2.5AlN-0.5B4C-2C. EDS shows presence of C-N-rich and primarily Al-rich phases. SiC-N, -SiC-3.1Al2O3, -SiC-0.5B4C-2C Thesr comparison purposes. SiC-N is conr material against heavy ihedral angle measured in this material also varies from 20 to 80. This would explain the observed absence of intergranular phase in some of the grain boundaries studied. Aside from the common metallic impurities that segregate in the triple junctipresence of residual carbon and other secondary phases was also observed. As shownFigure 5-79, C-N rich grains and primarily Al-rich phase are present in this material. C-N rich grain Al-rich phase Bulk Grain e groups of materials are included in this dissertation primarily fo sidered to be the current state-of-the-art armo threats.19 This material is produced by Cercom via pressure assisted densification and contains AlN as an additive. The mechanical properties of this material would thus

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118 serve as a good baseline for ceramic armor materials. -SiC-3.1Al2O3 is included to determine if different Al sources will affect the mechanical properties, and grain boundary and triple junction compositions. -SiC-0.5B4C-2C is basically a solid statesintered material since formation of liquid phase is not expected due to the absencaluminum. Processing and Polytypes Table e of 5-8 summarizes the density measurement, chemical analysis, and the phase -N, -SiC-3.1Al2O3 and -SiC-0.5B4C-2C. These materials sinter alue 15R mblage of this material to SiC containing Al (Aed. assemblage observed in SiC ed to a high density, comparable to the theoretical density of pure SiC. In the case of SiC-N, the resulting polytype is 6H and chemical analysis revealed that considerable amount of oxygen and nitrogen was retained after processing. In this material, oxygen and nitrogen was present at 0.56 and 0.17 wt. %, respectively. In the case of -SiC-3.1Al2O3, the resulting polytype is also primarily 6H, however, presence of 4H and 15Rare now detected. Chemical analysis for retained oxygen in this material yielded a vof 1.08 wt. % which is relatively higher than what was observed in -SiC-1.65Al hot pressed at 2100C. It has to be noted, however, that introduction of Al2O3 introduces more oxygen in the system than Al addition alone. Thus, it is expected that higher amount of oxygen is retained in this material. For -SiC-0.5B4C-2C, the resulting polytype is also primarily 6H, with minorstructure present. Comparison of the phase asse l, AlN and Al2O3) reveals the effect of Al addition on the resulting polytype. Presence of Al promotes the formation of 4H SiC polytype. As shown in Table 5-8, when only B and C were used as additives, presence of 4H polytype was not observ

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119 Chemical analysis performed on -SiC-0.5B4C-2C revealed that retained oxygen is almost non-existent. This is due to the effectiveness of C in removing oxygen in the system through the formation of CO and SiO gas. Table 5-8. Density, chemical analysis and phase assemblage of SiC-N, -SiC-Al2O3 aSample (wt.%) Density Chemical Analysis nd -SiC-0.5B4C-2C. (g/cc) (wt. %) Phase Assemblage (wt. %) Oxygen Nitrogen 3C 2H 4H 6H 15R SiC-N 3.22 0.56 0.17 0 0 0 100 0 -SiC3.1Al2O3 02 3.21 1.08 0.004 0 9.7 83.17. -SiC-0.5B4C-3.19 0.02 2C 0.008 0 0 0 94.8 5.2 Microstructure anhanicopertieepresentative microstructure in this group of materials is shown in Figure 5-80. quiaxed with and average size of about 2 m. I f f fracture was determined via the obser d Mec al Pr s R In the case of SiC-N, the grains are generally e n -SiC-0.5B4C-2C, the grains are generally equiaxed with some small volume of small elongated grains. The average grain size is also about 2 m and the aspect ratio of the elongated grains varies from 1 to 2. In the case of -SiC-3.1Al2O3, the microstructureshows a bimodal grain distribution consisting of small elongated grains and fine equiaxedgrains. The average size of the equiaxed grains is about 0.5 m and the aspect ratio of the small elongated grains also varies from 1 to 2. The measured hardness, toughness, and observed mode of fracture in this group omaterials are summarized in Table 5-9. The mode o vation of the Vickers indent on the polished surface and the appearance of the fracture surface of the material shown in Figure 5-81 and 5-82, respectively. In the case of SiC-N, the measured hardness and toughness are 22.80.4 GPa and 4.80.1 MPa

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120 m1/2, respectively. This material was also observed to fracture intergranularly. The highest hardness and lowest toughness in all of the materials studied was observed in SiC-0.5B4C-2C. This material was also observed to fracture in transgranular mode. -SiC-3.1Al2O3, the fracture mode is mixed. This material exhibits high hardness primarily due to the small grain sizes observed. -For Figure 5-80. TEM of microstructure observed in SiC-N, -SiC-0.5B4C-2C and -SiC-3.1Al2O3. Table 5-9. Hardness, toughness, and mode of fracture observed in SiC-N, -SiC-3.1Al2O3 Sample (wt. %) (HV1, GPa)Toughness (MPa m1/2)Fracture Mode and -SiC-0.5B4C-2C. Hardness SiC-N 22.80.4 4.80.1 Intergranular -SiC-3.13.80.2 Mixed Al2O3 23.50.7 -SiC-0.5B4C-2C 25.70.4 2.40.1 Transgranular Figure 5-81. Optical micrographs of Vickers indent on a polished surface of SiC-N, -SiC-0.5B4C-2C and -SiC-3.1Al2O3. Marker bars shown are 10 m.

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121 Figure 5-82. SEM of fracture surfaces of SiC-N, -SiC-0.5B4C-2C and -SiC-3.1Al2O3. Magnification for SiC-N is 4000X while in -SiC-0.5B4C-2C and -SiC-3.1Al2O3 is 2000X. Grain Boundaries, Triple Junctions and Secondary Phases SiC-N High resolution TEM of several grain boundaries observed in SiC-N are shown in Figure 5-83. Presence of clean grain boundary and grain boundary with crystalline and amorphous intergranular films were observed in this material. The thickness of the intergranular film observed is about 1 nm. The composition of the grain boundary phase is generally Al and O rich, as shown in Figure 5-84 and 5-85. Although the exact chemistry of the grain boundary phase was not determined, it is believed that the Al and O rich phase is in the form of AlO. The triple junction phase composition, as shown in Figure 5-86 and 5-87, are also generally Al and O rich. EELS taken from the triple junction confirm that the Al and O rich phase is in the form of crystalline AlO. The dihedral angles measured, based on the STEM and EFTEM images of the triple junction ranges from 40 to 70. It is al angles is m 23 23noticeable that the variation of the dihedrinimal in this specimen.

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122 a) b) c) Figure 5-83. HRTEM of several grain boundaries observed in SiC-N. a) grain boundary with crystalline intergranular film of thickness ~1 nm. b) clean grain boundary. c) grain boundary with amorphous intergranular film of thickness ~1 nm. Grain Boundary Bulk Grains Figure 5-84. STEM of a grain boundary in SiC-N. Segregation of Al and O at the grain boundary was observed.

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123 Figure 5-85. EFTEM of a grain boundary in SiC-N. Elemental maps show segregation of Al and O at the grain boundary. Grain Boundary Bulk Grains Triple Junction Figure 5-86. STEM of a triple junction in SiC-N. Accompanying EDS were taken from the triple junction, grain boundary and bulk grains. EDS of the triple junction shows segregation of Al and O.

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124 Figure 5-87. EFTEM of a triple junction in SiC-N. Elemental maps show segregation of Al and O in the triple junction. -SiC 3.1 wt. % Al2O3 High resolution TEM of the grain boundaries observed in this material is shown in Figure 5-88. Presence of grain boundary with amorphous intergranular film was observed. As determined by EDS and EFTEM, shown in Figure 5-89 and 5-90, the grain boundary film composition is generally rich in Al and O and is believed to be in the form of Al2O3. Most of the composition data for the triple junction phase of this material was derived from the EDS taken at the triple junction. This is primarily due to the observed fall-off of triple junction phase in the thin section of the specimen. Although triple junction phase fall-off is common in most of the materials, it was observed that it was more pronounced in -SiC-3.1Al2O3. As shown in Figure 5-89, the triple junction phase

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125 is also Al and O-rich and is believed to be of Al2O3. Dihedral angles measured from the STEM of the triple junction phase vary from 20 to 70. c) b) a) Figure 5-88. HRTEM of several grain boundaries in -SiC-3.1Al2O3. a) and b) shows grain boundary with IGF of about 1 nm. c) shows a grain boundary with amorphous IGF of less than 1 nm. Bulk Grains Grain Boundary Triple Junctions Figure 5-89. STEM of a triple junction in -SiC-3.1Al2O3. Accompanying EDS were taken from the triple junction, grain boundary and the bulk grains.

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126 Figure 5-90. EFTEM of a grain boundary and a triple junction in -SiC-3.1Al2O3. The particular triple junction shown is not filled with secondary phase. Elemental maps show segregation of Al and O at the grain boundary. -SiC 0.5 wt. % B4C 2 wt. % C Grain boundaries observed in this material are generally clean. As shown in Figure 5-91, distinct transition between lattice fringes of two grains are generally observed indicating the absence of secondary phase in the grain boundary. Since B and C will not form secondary phases in the absence of aluminum,2 the observed absence of secondary phases in the grain boundary is expected. The same is true for the absence of triple junction phases observed in this material. As shown in Figure 5-92, variation in the composition between triple junction, grain boundary and bulk grains does not exists. Presence of metallic impurities and residual carbon, however, are expected due to the excess carbon available and metallic impurities of the starting powder. These secondary phases, which are usually found in triple junctions, are shown in Figure 5-93. Presence

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127 of B-C-rich grains, possibly in the form of B4C, was also observed to exist, as shown in Figure 5-94. c) b) a) Figure 5-91. HRTEM of grain boundaries observed in -SiC-0.5B4C-2C. Bulk Grains Triple Junction Grain Boundary Figure 5-92. STEM of a triple junction in -SiC-0.5B4C-2C. EDS taken from triple junction, grain boundary, and bulk grains does not indicate any compositional variation.

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128 Residual Carbon Bulk Grains Metallic Impurities Figure 5-93. STEM of a triple junction containing residual carbon and metallic impurities in -SiC-0.5B4C-2C. Figure 5-94. EFTEM of triple junction and grain boundary in -SiC-0.5B4C-2C. Elemental maps show presence of B-C-rich grains. Composition variation across grain boundary was not observed.

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CHAPTER 6 SUMMARY AND DISCUSSION: MECHANICAL PROPERTY, GRAIN BOUNDARY, AND TRIPLE JUNCTION CHARACTERISTICS A summary of observation for the materials studied will be provided in this chapter to facilitate the discussions. Although the observations from the ABC-SiC system studied is applicable to the materials sintered from -SiC starting powders, delineation between the materials will still be observed. -SiC Sintered with 0.6 wt.% B, 2 wt.% C and 0 4 wt.% Al (ABC-SiC) A summary of observations for the materials sintered from the -SiC starting powders is given in Table 6-1. As presented in the previous chapter, correlation between the grain boundary and triple junction characteristics, and the toughness and mode of fracture in this material exists. It has been observed that the formation of amorphous intergranular film in these materials correlates with the change in fracture mode from transgranular to intergranular. The change in fracture mode, in turn, correlates with an increase in the fracture toughness from 2.7 to 6.1 MPa m1/2. The correlation between the change in fracture mode and the grain boundary and triple junction characteristics can be explained by the formation of amorphous intergranular films in the grain boundary. The presence of amorphous intergranular film, which was observed starting at 1.5 wt. % Al addition, provides a presumably weak secondary phase where crack propagation could easily occur. Since crack propagation is a minimum energy process, the cracks will seek out a path with minimum energy requirement to propagate. Since the average grain sizes observed for the materials with 129

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130 intergranular phase (with 1.5 and 4 wt. % Al addition) are relatively large, ~5 m, the energy requirement for the crack to propagate through the grains will be larger than what will be required for propagation through the intergranular phase. Table 6-1. Summary of observations for SiC ceramics sintered from -SiC starting powders. Materials (wt. %) Toughness (MPa m1/2) Hardness HV1 (GPA Mode of Fracture Secondary Phases Grain Boundary Triple Junction Others -SiC 0.6B -2C 2.6 0.2 25.5 0.7 T None None C -SiC 0.5Al 0.6B -2C 2.6 0.1 20.9 0.9 T None Amorphous Al-O C -SiC 1Al 0.6B -2C 2.7 0.1 21.5 0.9 T None Partially crystalline Al-O C -SiC 1.5Al 0.6B -2C 6.1 0.3 15.6 1.2 I 1 nm, amorphous, Al-O Crystalline Al-O C -SiC 4Al 0.6B -2C 6.1 0.2 20.8 0.9 I 1 nm, amorphous, Al-O Crystalline Al-O C, Al8B4C7 T = Transgranular, I = Intergranular Influence of Residual Stress Aside from energy consideration, effect of residual stress due to the formation of secondary phases also contributes to the observed intergranular fracture and toughening. The residual stresses are generally attributed to the thermal expansion mismatch between the grains and the secondary phases.109,110 Looking at the particular case of the intergranular films, an analysis similar to a fiber in an infinite matrix could be applicable. Following the derivation by Green,110 the pressure at equilibrium is given by, mffmffAfmmfEETTEEP2121110 (6-1)

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131 where P is the equilibrium pressure, E is the Youngs modulus, is the coefficient of thermal expansion, TO is the temperature at which the misfit stress starts to arise, TA is the ambient temperature, and is the poisons ratio. The subscripts m and f refers to matrix and fiber, respectively. The residual stress (radial and tangential) on the fiber can be calculated from the equilibrium pressure and is given by rr = = -P. By assuming that the fiber in this case is the intergranular film (Al-O-rich composition in the form of Al2O3) and the matrix is the SiC grains, the residual stress on the intergranular film can be calculated as a function of temperature, as shown in Figure 6.1a. The values of the variables used in the calculation are as follows: EAl2O3 = 400 Gpa, Al2O3 = 8.3 x 10-6/C, Al2O3 = 0.23, ESiC = 480 GPa, SiC = 4.4 x 10-6/C, and SiC = 0.17.109 A similar approach can be applied at the triple junction phase. In this case, an assumption that the triple junction phase is a spherical particle in an infinite matrix has to be made. And in this case, the pressure at the interface is given by, mppmApmmpEETTEEP121220 (6-2) where the subscript p and m now refers to the particle and matrix, respectively. The residual stress on the triple junction phase can be calculated from the equilibrium pressure and is given by rr = = -P. The result of the calculations for the residual stress on the triple junction phase is shown in Figure 6-1b. As shown in Figure 6-1, the residual stress on the intergranular film and triple junction phase is positive. This indicates that the intergranular film and triple junction phase will be in tension. Since the calculated residual state is high, it is highly probable

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132 that failure may initiate near, or at, the intergranular film and triple junction phase. This failure at, or near, the intergranular film and triple junction phase will enable intergranular mode of fracture in the material and activates the toughening mechanisms such as crack deflection and crack bridging. Activation of crack deflection and crack bridging would then increase the fracture toughness in the SiC ceramic. Figure 6-1. Calculated residual stresses due to the presence of secondary phase. a) Residual stress in the intergranular film. b) Residual stress in the triple junction phase. It has to be noted that the calculation presented for the residual stress due to the presence of secondary phases was highly simplified. Superposition of stress fields was not considered. The equations used are applicable only for a case wherein perfect interfacial bonding between the filament and the matrix exists. Due to the amorphous nature of the intergranular films observed on the materials studied, perfect interfacial bonding may not be the case. Thus, the values of the calculated residual stress only serve as an approximation and means of relative comparison. The nature of the residual stress (tensile) on the intergranular film and triple junction phase, however, is still valid and intergranular cracking are still expected to occur.

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133 Absence of Intergranular Films in Materials with Al Additive in Excess of Solubility Limit Another issue that crop up during the investigation performed in SiC sintered from -SiC starting powders is the observation that even with the solubility limit of Al in SiC was exceeded, formation of the grain boundary film was still not observed. This phenomenon was observed in the case of -SiC-0.5Al-0.5B-2C and -SiC-1Al-0.5B-2C specimen. The amounts of aluminum additive used in these materials are greater than [or equal to, in the case of -SiC-0.5Al-0.5B-2C] the solubility limit of Al in SiC. The solubility limit of Al in SiC are reported to be 0.2 wt. % Al, based on the work of Kinoshita et al,108 and 0.5 wt. % Al at 2000C, based on the work of Tajima and Kingery.107 The expectation that formation of intergranular film once solubility limit of additives in liquid phase forming system was based on the reasoning that the excess additive should be accommodated in the grain boundary or triple junction phases. And if it is in the triple junction or the grain boundary, it is reasonable to assume that the excess additives should remain upon cooling, unless evaporation of the excess additive occurs. Experimentally, in the presence of liquid phase forming additives, formation of intergranular phase in the grain boundary was almost always observed.111 Theoretical considerations pertaining to the thickness of the intergranular films due to the liquid forming phase in sintered ceramics have been reported in the literature.111,112 Lange113 applied a theory developed for plates separated by a liquid layer to the case of spherical particles undergoing liquid-phase sintering. His analysis showed that although the thickness of the liquid layer decreases progressively with time, a liquid layer of finite thickness always remains between the particles within any experimental time frame as

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134 long as the liquid completely wets the solid. Clarke111 even extended the analysis to address the issue of the equilibrium thickness of intergranular films. Based on two continuum approach, one based on interfacial energies and the other on the force balance normal to the boundary, he concluded that there will exists a stable thickness for the intergranular film and that it will be of the order of 1 nm. He had shown that the origin of equilibrium thickness is the result of two competing interactions, an attractive Van der Waals-dispersion interaction between the grain on either side of the boundary acting to thin the film and a repulsive term, due to the structure of the intergranular liquid, opposing this attraction. The main assumption on the theoretical considerations is the presence of complete wetting of the solid by the liquid phase. Wetting characteristics in the presence of liquid phase-forming system are dependent on the interfacial energies and dihedral angle, and is given by the relation: slss22cos (6-1) where is the dihedral angle, ss and sl are the solid/solid and solid/liquid interfacial energies, respectively. Dihedral angles observed for -SiC with higher aluminum content (1.5 and 4 wt. % Al) suggests that partial penetration of the triple junction phase into the grain boundary is possible. And the fact that formation of intergranular films in these materials were observed suggests that dihedral angle effect is not the main cause of the absence of grain boundary phase in the low Al containing materials.

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135 The solution to the apparent non-formation of grain boundary phase in materials sintered with Al above its solubility limit in SiC (0.5 and 1 wt. % Al) may be quite simple. Embedded in the assumption for complete wetting of the solid phase by the liquid phase in the theoretical treatment is the fact that there should be enough volume of liquid phase in the system. For the production of advance ceramics by liquid phase sintering, the amount of liquid phase produced at the firing temperature is typically in the range of 5 to 10 vol. %.29 In the case of -SiC with 0.5 and 1 wt. % Al addition, the volume of the liquid phase generated may just be too small to fill the grain boundaries between the SiC particles. The fact that segregation of Al and O in the triple junction of these materials was observed indicates that the liquid phase are present, however, the volume of the liquid phase is insufficient. SiC Ceramics Fabricated from -SiC Starting Powders The results that will be discussed are the ones reported under the following sections in the experimental results chapter: -SiC sintered with Al additives, -SiC sintered with AlN additives, and SiC-N, -SiC-3.1Al2O3 and -SiC-0.6B4C-2C. In these materials, except for the solid state sintered -SiC-0.6B4C-2C, the amount of aluminum additive (in the form of Al, AlN and Al2O3) used is greater than 1.5 wt. % Al. Per the results of investigation in SiC ceramics sintered from -SiC starting powders, 1.5 wt. % Al addition corresponds to the transition from transgranular to intergranular mode of fracture. Use of higher amount of aluminum additive in the same system, as illustrated by the 4 wt. % Al addition, results in the same failure mode. Following the argument presented earlier that enough liquid phase is formed at 1.5 wt. % Al addition in -SiC starting powders, it is expected that the materials being considered in this section would fracture intergranularly.

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136 Comparison of Mechanical Properties and Grain Boundary and Triple Junction Characteristics of SiC Sintered with Equimolar Al Addition Table 6-2 presents the summary of observations for the materials with equimolar Al addition pertaining to the grain boundary and triple junction characteristics and mechanical properties. SiC ceramics with B4C and C additions are included to illustrate the effect of B and C. All of these materials are hot pressed at 2100C for one hour under an argon atmosphere at an applied load of 20 MPa. Table 6-2. Mechanical properties, grain boundary and triple junction characteristics of SiC materials with equimolar aluminum addition. Composition Intergranular Film Triple Junction Other Secondary Phases Toughness (MPam1/2) Hardness HV1 (GPa) Fracture Mode 1.65 Al Al-O, am, 0-1 nm Al-O, cr MI 4.70.4 22.10.7 M 2.5 AlN Al-O, am, 0-1 nm Al-O, cr MI 3.50.2 21.90.3 I 3.1 Al2O3 Al-O, am, 1nm Al-O, cr MI 3.80.2 23.50.7 M 1.65Al-0.5B4C-2C Al-O-B, Al, am, 0-4 nm Al-O-B, cr MI, C, B-C 3.10.1 20.30.3 M 2.5AlN-0.5B4C Al-O-B, Al 0 nm Al-O-B, cr MI, C, B-N 3.40.4 21.00.6 M 2.5AlN-0.5B4C-2C Al-O-B, Al, am, 0-10 nm Al-O-B, cr MI, C, C-N 2.80.2 22.70.3 M 0.5B4C-2C None None C, B-C 2.40.1 25.70.4 T am = amorphous, cr = crystalline. I = intergranular, T = transgranular, M = mixed mode, MI = metallic impurities. Al-O composition in triple junction is usually in the form of Al2O3. Compositions are given in wt. %. Starting SiC powder is 6H -SiC polytype. Comparison of SiC Ceramics with Equimolar Al Addition As shown in Table 6-2, different sources of Al do not affect the grain boundary and triple junction characteristics of the hot pressed SiC ceramics. The triple junctions observed are all crystalline Al2O3 independent of the type of additive used (Al, AlN or Al2O3). Dihedral angles measured from the STEM and EFTEM images of the triple junction in these materials also showed a similar range of about 15 to 90. This indicates

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137 that it is highly probable that the triple junction phase is continuous into the grain boundary. The grain boundary (or the intergranular films) also consistently contains Al-O-rich phase. Although presence of clean grain boundaries was observed for -SiC with Al and AlN additions, the width of the intergranular films, whenever present, is typically around 1 nm. This confirms the existence of equilibrium thickness of intergranular film in liquid phase sintered ceramics calculated by Clarke.111 Although the grain boundary and triple junction characteristics are the same, differences in the mechanical properties of SiC sintered with equimolar aluminum addition exist. This is particularly true for the SEPB toughness measured. As shown in Table 6-2, the toughness value of -SiC-1.65Al is considerably higher than the toughness in -SiC-2.5AlN and -SiC-3.1Al2O3. The hardness values, however, are statistically similar. The mode of fracture is purely intergranular in -SiC-2.5AlN and mixed in the other two materials. It should be noted that the fracture mode observation was based on the Vickers indent appearance on the polished surface and the appearance of the fracture surface. Thus, determining if it is primarily intergranular or transgranular would be difficult unless quantification of the amount of intergranular and transgranular mode per unit volume is made. Unfortunately, that data is still not available as of this writing. Per the discussion in previous chapter, the differences in toughness observed can be attributed to the grain size and morphology observed. This is particularly true since the grain boundary and triple junction characteristics of these materials are similar. As shown in the previous chapter and reproduced in Figure 6-2, the microstructure achieved is distinct for each additive. In -SiC-1.65Al, a microstructure consisting of large elongated grains in a matrix of small equiaxed grains was observed. Average grain size

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138 of the equiaxed grains is about 1 m and the aspect ratio of large elongated grains ranges from 2 to 5. The width of the elongated grains also varies from about 1 m to 3 m. In the case of -SiC-AlN, the microstructure is generally consisting of fine equiaxed grains. The average grain size is about 1.5 m, with the grain size distribution covering a relatively large range, i.e., presence of grains as small as 0.5 m and as large as 4 m was observed. In the case of -SiC-3.1Al2O3, the resulting microstructure is similar to those observed in -SiC-1.65Al. The grain sizes and aspect ratio, however, are relatively smaller. The average grain size of the equiaxed grains is about 0.5 m while the aspect ratio of the small elongated grains ranges from 1 to 2. Figure 6-2. TEM of microstructure observed in -SiC sintered with equimolar Al addition. The presence of elongated grains in the microstructure have been observed to facilitate an increase in the toughness of SiC ceramics.9 These elongated grains can act as a reinforcing phase that promotes crack bridging and deflection, resulting in improved toughness. The differences in the toughness of the materials with equimolar Al addition would then be explainable if the toughening mechanisms in these materials differed due to the differences in the microstructure. Figure 6-3 shows the crack path introduced by Vickers indentation in these materials. Although the grains are not clearly

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139 distinguishable due to difficulty in etching these materials, the appearance of the crack indicates the existing toughening mechanisms. As shown, crack deflections are observed in all the materials compared. Crack bridging, however, were only observed in -SiC-1.65Al and -SiC-3.1Al2O3. Figure 6-3. SEM of crack profiles in SiC materials with equimolar Al addition. a) Crack path in -SiC-1.65Al showing crack deflection and grain bridging. b) crack path in -SiC-2.5AlN showing crack deflection only. c) Crack path in -SiC-3.1Al2O3 showing crack deflection and grain bridging. In summary, it was shown that differences in the grain boundary and triple junction characteristics of SiC ceramics sintered with equal molar concentration of Al from different Al sources does not exists. The change in the mechanical properties, specifically the toughness measured, was attributed to the change in the microstructure of

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140 the materials observed. The change in microstructure results in activation of different toughening mechanism in the material. As shown, for purely equiaxed grains, the toughening mechanism is primarily crack deflection and in the presence of elongated grains, crack bridging becomes activated resulting in additional toughness improvement. Effect of B and C Additions Table 6-2 also shows the summary of observations for materials containing B (in the form of B4C) and C in addition to the Al or AlN additive. Changes in the grain boundary and triple junction characteristics, as compared to materials with pure Al or AlN addition, was evident. The triple junction in the materials containing B4C and/or C shows that segregation of Al-O-B occurred. The triple junction phase, as observed from its appearance and presence of thickness fringes, however, is still crystalline. The grain boundary phase (or intergranular films) composition also changed from a pure Al-O-rich phase (for materials with pure Al or AlN additions) to a composition containing B (Al-O-B-rich phase). Presence of a grain boundary containing pure Al was also observed in all the materials containing B4C and/or C. The thickness of the amorphous intergranular film also varies, and the ranges are greater than those observed for materials with pure Al or AlN addition. The thick intergranular films, however, are mostly amorphous carbon. Changes in the mechanical properties of materials with B4C and/or C additions were also observed. The hardness and toughness values measured in these materials are generally lower than those observed in SiC with Al or AlN additions only. The mode of fracture also changes to a mixed mode with primarily transgranular fracture, as shown in Figure 6-4. Presence of crack deflection on the crack path was observed to be at a minimum in all of the samples.

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141 The addition of B (B4C) and C also results in a change in the microstructure of the final ceramics. As shown in the previous chapter and reproduced in Figure 6-5, grain sizes of these materials are generally larger than those observed in materials with pure Al or AlN addition only. In the case of -SiC-0.5B4C-2C, the average grain size is about 2 m and the aspect ratio of the elongated grains varies from 1 to 2. For -SiC-2.5AlN-0.5B4C and -SiC-2.5AlN-0.5B4C-2C, the average grain size is almost similar (~2.5 m) and the grains are generally equiaxed. The largest grain size in all of the materials studied was observed to occur in -SiC-1.65Al-0.5B4C-2C. The average grain size in this material is about 5 m and the aspect ratio of the elongated grains varies from 1 to 2. Figure 6-4. SEM of crack profiles in materials containing B4C and C. a) crack path in -SiC-0.5B4C-2C. b) crack path in -SiC-1.65Al-0.5B4C-2C. c) crack parth in -SiC-2.5AlN-0.5B4C, and d) crack path in -SiC-2.5AlN-0.5B4C-2C. Minimum crack deflections were observed. The grain growth observed in the materials containing B (in B4C) and C was not unexpected. It has been reported in the literature that boron and carbon promotes grain growth.53 In the presence of Al, further enhancement of grain growth can be achieved due to its cooperative effect with B and C.53 Presence of the liquid phase, due to the addition of aluminum, also provides another densification route. The large grain size observed in

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142 -SiC-1.65Al-0.5B4C-2C can thus be attributed to the cooperative effect of Al, B and C on grain growth. The case of AlN-containing SiC, however, is different. Presence of nitrogen, either in AlN or N2 form, retards the growth of the grains in liquid phase sintered materials.8,35 Thus, smaller grain sizes were observed for -SiC-2.5AlN-0.5B4C and -SiC-2.5AlN-0.5B4C-2C. b Figure 6-5. TEM of microstructure observed in -SiC sintered with further addition of B4C and C. Two obvious effects of C and B addition on the materials studied have been observed. The first one is the change in the grain boundary and triple junction phase characteristics in the resulting materials. The other is the enhanced grain growth observed in the resulting microstructure which was possibly due to the enhanced densification during processing. Which of these two effects are responsible for the

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143 mechanical property variations, however, cannot be ascertained at this time. Comparison of the properties of the materials with equimolar Al additions and materials with B and C additions, however, may provide a certain degree of clarity. As shown in Table 6-2, materials with equimolar Al additions only, possess better mechanical property combination than materials with B and C. Although highest hardness was observed for the solid state sintered material (-SiC-0.5B4C-2C), the measured toughness was minimal. This combination of property is not desirable for armor applications wherein achievement of good hardness and toughness was desired. The low toughness and high hardness value of this material is due to the absence of intergranular films that would have enabled intergranular failure and activate other toughening mechanisms. The toughness and hardness values measured for materials containing B and C (excluding -SiC-0.5B4C-2C) are generally lower than those observed for materials with equimolar Al additions only. Although the change in the chemistry of the grain boundary and triple junction phases of B and C containing materials may not be that important, its effect on the mechanical properties can not be ruled out. The mere presence of the intergranular films, at the very least, enables some degree of intergranular fracture which increases the toughness. The change in the microstructure of B and C containing materials may have a more important effect on the mechanical properties. The larger grain sizes in these materials caused the lower hardness values observed. The larger grain sizes and the equiaxed nature of the grains may also have caused the lower toughness values. Presence of elongated grains is one of the key in activating toughening mechanisms such as crack deflection, crack bridging and microcracking.9 Although

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144 elongated grains are present in -SiC-1.65Al-0.5B4C-2C, the thickness of these grains and the low aspect ratio are just not suitable for activations of aforementioned toughening mechanisms.9 Comparison of -SiC Ceramics Sintered with Al or AlN Summary of observations pertaining to grain boundary and triple junction characteristics and mechanical properties of -SiC sintered with Al or AlN additions are presented in Table 6-3. All materials are hot pressed at the indicated temperatures under an argon atmosphere at 20 MPa load, except for SiC-N. SiC-N is the current state of the art ceramic armor material for heavy threat.19 It is manufactured by Cercom, Inc. via pressure assisted densification. The additives and processing temperature used are currently unknown and are proprietary to Cercom, Inc. Ceramatec Inc, however, stipulates that it contains 2.5 wt. % AlN. SiC-N is included in this group of sample for comparison purposes. It will be shown later that fabrication of materials sintered with Al or AlN with properties similar to SiC-N can be achieved based on the observations made. Effect of Hot Pressing Temperature and Amount of Additives As shown in Table 6-3, hot pressing temperature did not alter the grain boundary and triple junction chemistry of the materials studied. Presence of amorphous intergranular films with composition rich in Al and O were observed in most of the materials investigated. The triple junction phases are also observed to be primarily crystalline Al2O3. Increased amount of additive have shown to be effective in changing the grain boundary and triple junction chemistry. In the case of materials with higher Al addition (3.3 wt. % Al), presence of grain boundary phase primarily rich in Al, aside from the grain boundary phase which are Al-O-rich, were observed. Al-rich grains were also

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145 found in these materials. In the case of -SiC-5AlN hot pressed at 2200C, no intergranular film formation was observed through HRTEM. It is believed, however, that intergranular films of Al-O-rich composition are also present in this material as shown by EDS and EFTEM results. It is possible that the grain boundary, when observed through HRTEM, was just not properly oriented as indicated by the absence of direct transition between the lattice fringes of each grain. In the case of -SiC-5AlN hot pressed at 2000C, presence of Al-O-N and Al-O-rich grain boundary phase were observed. The variation in the grain boundary phase composition, however, does not seem to significantly alter the mechanical properties. Table 6-3. Mechanical properties, grain boundary and triple junction characteristics of SiC sintered with Al or AlN addition. Composition/ Hot Pressing Temperature Intergranular Film Triple Junction Other Secondary Phases Toughness (MPam1/2) Hardness HV1 (GPa) Fracture Mode 1.65Al, 1900C Al-O, Am, 0-1 nm Al2O3, Cr MI 4.00.2 25.30.7 M 1.65Al, 2100C Al-O, Am, 0-1 nm Al2O3, Cr MI 4.70.4 22.10.7 M 1.65Al, 2200C Al-O, Am, 0-1 nm Al2O3, Cr MI 5.70.1 20.80.3 M 3.3Al, 2000C Al-O, Al, Am, 0-3 nm Al2O3, Cr MI, Al grains 4.20.1 22.10.6 M 3.3Al, 2200C Al-O, Al, Am, 0-1 nm Al2O3, Cr MI, Al grains 6.80.1 20.50.5 M 2.5AlN, 2100C Al-O, Am, 0<1 nm Al2O3, Cr MI 3.50.2 21.90.3 I 2.5AlN, 2200C Al-O, Cr, 0-1 nm Al2O3, Cr MI 3.50.4 22.90.1 I 5AlN, 2000C Al-O, Al-O-N, Am, 0-1 nm Al-O-N, Cr MI 3.90.2 25.00.4 M 5AlN, 2200C Al-O, 0 nm Al2O3, Cr MI 4.20.1 24.10.6 I SiC-N Al-O, Cr,Am 0-1 nm Al2O3, Cr MI 4.80.1 22.80.4 I Am = amorphous, Cr = crystalline, MI = Metallic impurities, M = mix, I = Intergranular. All compositions are given in wt. %. Starting powders are all -SiC, except for SiC-N.

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146 Since the grain boundary and triple junction phase in -SiC sintered with Al or AlN addition are not significantly altered, the change in microstructure, resulting from the change in hot pressing temperature and amount of additives, is responsible for the changes in the mechanical properties observed. As shown in Table 6-3, the toughness values measured from -SiC with Al additions are generally higher than those observed in -SiC sintered with AlN additive. This trend can be attributed to the microstructure observed in these materials. As shown in previous chapter (Figure 5-22), the microstructure resulting from -SiC with Al addition generally shows presence of elongated grains in a matrix of small equiaxed grains when hot pressed at higher temperature. In the case of -SiC-1.65Al, hot pressing at 1900C resulted in a microstructure containing mostly equiaxed grains with an average grain size of 1 m. Increasing the hot pressing temperature to 2100C resulted in a microstructure consisting of elongated grains in a matrix of small equiaxed grains. Further increase of hot pressing to 2200C increases the aspect ratio and volume of elongated grains but retaining the average size of the equiaxed grains. Increasing the Al additive amount to 3.3 wt. % Al did not significantly changed the observed microstructure. In the case of -SiC with AlN additive, the microstructures are generally of equiaxed grains, as shown in Figure 5-54 in the previous chapter. Increasing the hot pressing temperature only results in normal grain growth thus no grain elongations were observed. Increasing the AlN content to 5 wt. % resulted in a decreased grain sizes which resulted in a high hardness values observed in these materials.

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147 The differences in microstructure of -SiC with Al or AlN additions may have caused different toughening mechanisms. Crack path observations in these materials, shown in Figure 6-6 and 6-7, could give an idea on what toughening mechanisms are active. The grains in these figures are not highly distinguishable but the crack paths indicate that toughening is achieved mostly through crack deflection. Crack bridging was only observed for -SiC sintered with Al additive processed at higher temperature where formation of elongated grains occurs. Figure 6-6. SEM of crack profiles in -SiC sintered with Al additives. a) -SiC-1.65Al hot pressed at 1900C. b) -SiC-1.65Al hot pressed at 2100C. c) -SiC-1.65Al hot pressed at 2200C. d) -SiC-3.3Al hot pressed at 2000C. e) -SiC-3.3Al hot pressed at 2200C. Achievement of SiC-N-like Properties Fabrication of SiC ceramics with properties (mechanical and grain boundary and triple junction characteristics) similar to that observed in SiC-N is one of the underlying objectives of this dissertation. The mechanical properties (toughness, hardness and mode

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148 of fracture) of SiC-N are shown in Table 6-3. This material has a good combination of hardness and toughness, 22.80.4 GPa and 4.80.1 MPam1/2, respectively, and fails in purely intergranular mode. The triple junction phase composition of this material is a crystalline Al2O3, and was observed to be continuous into the grain boundary. The intergranular film (or the grain boundary phase) thickness is usually about 1 nm and are present in crystalline and amorphous form. The composition of which is generally Al-O-rich and is believed to be Al2O3 due to the dihedral angles measured from the appearance of the triple junction phase. Figure 6-7. SEM micrographs of crack path in SiC-N and -SiC sintered with AlN addition. a) -SiC-2.5AlN hot pressed at 2100C. b) -SiC-2.5AlN hot pressed at 2200C. c) -SiC-5AlN hot pressed at 2000C. d) -SiC-5AlN hot pressed at 2200C. e) crack path in SiC-N. Examination of Table 6-3 would show that similar properties are achievable in -SiC ceramics sintered with Al or AlN. The triple junction and grain boundary phases of -SiC sintered with Al or AlN addition are similar to those observed in SiC-N. The only

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149 difference is the presence of crystalline intergranular films (aside from the usually observed amorphous one) in SiC-N. In the case of -SiC sintered with Al addition, toughness and hardness values of -SiC-1.65Al hot pressed at 2100C and -SiC-3.3Al hot pressed at 2000C are comparable with that of SiC-N. Hot pressing of -SiC-1.65Al and -SiC-3.3Al at 2200C results in a higher toughness but lower hardness. The mode of fracture in -SiC sintered with Al addition, however, are primarily mixed. The microstructure observed in these materials also contained elongated grains (at hot pressing temperature of 2100C and 2200C) whereas microstructure in SiC-N is primarily of equiaxed grains. In the case of -SiC sintered with AlN, the toughness values in these materials are generally lower than in SiC-N. The hardness in -SiC-2.5AlN, independent of hot pressing temperature, is comparable but an increase in AlN content (to 5AlN) resulted in much higher hardness values. The microstructure observed in -SiC-AlN is similar to that in SiC-N and the mode of fracture is also generally the same, except for -SiC-5AlN hot pressed at 2000C where a mixed mode of fracture occurred. As have been discussed in the previous section, the mechanical properties of -SiC sintered with Al or AlN additions are highly dependent on the microstructure. The microstructure, in turn, is dependent on the processing temperature and the type of additive used. To achieve SiC-N-like properties, the SiC ceramics must primarily have microstructure similar to that of SiC-N. And, as discussed above, the use of AlN as an additive is a good way to start. Of particular interest is the -SiC-5AlN hot pressed at 2200C. The toughness of this material differs to that of SiC-N by only 12.5%. The hardness, however, is much

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150 higher than in SiC-N. As shown in the literature,11 further annealing or heat treatment can increase the toughness of a SiC ceramics. Crystallization of the grain boundary phase is also possible upon additional heat treatment or annealing steps. The only drawback is that the hardness of the material will decrease due to the expected grain growth. Since the hardness of -SiC-5AlN is higher; the expected decrease in hardness due to annealing will bring the hardness of this material to the level of SiC-N. Thus, achievement of SiC-N-like characteristics is possible through an additional annealing step. Further annealing, however, will not work for -SiC sintered with Al. As shown by the microstructure observed in this material and results reported in the literature,2,53 excessive grain elongation will occur. Achievement of SiC-N-like properties in these materials, however, is still possible by circumventing the grain growth. One of the first approaches will be to use the low hot pressing temperature (1900C) as the annealing temperature. This, however, may not be enough since grain growth was reported in the literature at this temperature.3 Another possible approach would be to use an applied pressure during annealing. Studies shows that grain growth are greatly inhibited at this condition.11 The use of nitrogen atmosphere during annealing will also reduce grain growth in -SiC sintered with Al additive. To summarize, SiC-N like properties can be achieved from -SiC sintered with either the addition of Al or AlN. The advantage of using -SiC-AlN, however, lies on the effect of AlN or the presence of nitrogen on the grain growth behavior in this material. Thus, a single additional annealing or heat treatment stage can added to the fabrication route and produce desirable mechanical properties such as those present in SiC-N.

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CHAPTER 7 CRACK HEALING OF SELECT SPECIMEN The phenomenon of crack healing has been observed in a variety of different ceramic materials including single crystal,66 polycrystalline ceramics,67-71 inorganic glasses,72 and ceramic composites.73-92 It has also been observed in metals93 and polymer materials.94 This phenomenon is usually described as a process where the surface crack of a given material closes, either by crack-face rebonding or by filling up of the crack face, after a given heat treatment. Associated with the crack closure is an observed increase in the flexural strength of the material. The recent increase in the number of studies, in the late 1990s,73-92 on crack-healing is driven by the promise of recovering the mechanical properties of structural materials that are compromised due to inherent flaws introduced during processing and handling. Application of a crack healing treatment as a post-production process will reduce, if not entirely eliminate, the flaws/cracks introduced during processing and handling. Crack healing will greatly enhance the reliability and integrity of the structural material in question. It will also decrease the machining, maintenance, and inspection costs and increase the lifetime of the ceramic material. Most of the crack healing studies in the literature were done in an oxidizing environment.73-92 In most cases, complete recovery of the flexural strength of the indented sample can be achieved within an hour of heat treatment. The mechanism for the crack healing is usually attributed to the formation of the oxidation products in the crack surface, bonding the crack walls and rebonding the crack faces. The increase in the 151

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152 flexural strength observed is usually attributed, aside from the reduction in flaw sizes, to the stress relaxation upon heat treatment and possible crack tip blunting. Reports of crack healing in a reducing environment are minimal in the literature.73-92 The flexural strength increase associated with crack healing in a reducing environment is small when compared to that gained from crack healing in an oxidizing environment. Since flexural strength increase associated with crack healing in a reducing environment is smaller compared to flexural strength in an oxidizing environment, it is possible that even if it was observed, it was not usually reported. In this study, it will be demonstrated that crack healing, even in inert atmosphere, can be done in SiC ceramics. Materials and Methods The materials used to demonstrate crack healing phenomena in SiC ceramics is shown in Table 7-1. These materials were manufactured from a -SiC starting powder and were used due to the differences in their mechanical properties. SS-SiC has the highest hardness and lowest toughness, ABC-SiC shows the lowest hardness, and YAG-SiC exhibits a good combination of hardness and toughness. All these materials exhibit a microstructure consisting of large elongated grains due to the anisotropic grain growth associated with -to--SiC phase transformation during processing. Table 7-1. Composition, processing conditions and mechanical properties of materials used for crack healing. Sample Composition (wt. %) Processing Condition Density (g/cc) Hardness (GPa) Toughness (MPam1/2) SS SiC -SiC-0.6B-2C 2100C/1hr 3.170.01 25.50.7 2.60.2 ABC-SiC -SiC-1.5Al-0.6B-2C 2100C/1hr 3.130.01 15.61.2 6.10.3 YAG-SiC -SiC-0.5Y2O3-0.43Al 2200C/1hr 3.280.01 23.20.7 7.31.0

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153 The specimens used in the crack healing experiments were in the form of a rectangular test pieces of dimension 3 mm x 4 mm x 45 mm. Prior to crack healing, a semi-elliptical crack was introduced to the specimen via Vickers indenter using a load of about 98 N (~ 10 kg). The indentation was done in the center of the specimen which corresponds to the center of the surface in tension during four point bend testing. The load used and position of the indentation was done to make sure that the materials will fail right at the indentation during fracture. As shown in Figure 7-1, the surface cracks radius varies depending on the material although the indentation load was the same. The surface crack diameter observed in SS SiC is about 200 m, ~100 m in ABC-SiC and ~150 m in YAG-SiC. Figure 7-1. Optical micrographs of the Vickers indentation using a 98 N load. a) SS-SiC. b) ABC-SiC. c) YAG-SiC. The crack healing experiments were performed in three different atmospheres. Experiments in an argon atmosphere were done to simulate an inert environment. In this case, the specimens were wrapped in a graphite foil inside a glove box purged with nitrogen gas. This was done to ensure the lowest oxygen partial pressure possible and avoid passive oxidation of the SiC materials. The experiments in argon atmosphere will also prove if crack healing of SiC ceramics can be achieved in an inert environment. Experiments in zero grade air (80 vol. % N2, 20 vol. %) were done to simulate possible effect or correlation of passive oxidation on crack healing. The use of water-containing

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154 atmosphere, on the other hand, could show possible effect of an environment where both active and passive oxidation of SiC will occur. The water vapor was introduced into the gas stream using a humidifier and utilizing argon as the carrier gas. The resulting gas mixture delivered into the reaction tube contains approximately 50 vol. % H2O. The crack healing experiments were done using a high temperature tube furnace capable of temperature of up to 1700C. The tube furnace has a heat zone of about 6 inches with a temperature variation of about 5 C. The crack healing temperature used in all the experiments performed was 1300C. The exposure time used was fixed at 3 hours. Crack healing of Si3N4/SiC composites have been reported possible in this conditions.74-77 The specimens were loaded into the tube furnace (alumina tube) in an alumina boat. The tubes were then evacuated and purge with argon gas for three hours prior to each experiments. The furnace was then heated to 1300C at a heating rate of 10C/min. The desired atmosphere were then introduced upon reaching 1300C and maintained for three hours. After the desired exposure time, the gas used was then turned off and the furnace was allowed to cool down to room temperature at a rate of 10C/min. The specimens weight was measured before and after each experiment. All fracture tests were performed using a four-point loading system shown in Figure 7-2. The support span of the four-point jig assembly is 40 mm easily accommodating the specimen size used. The cross-head speed used in all the fracture tests was 0.5 mm/min. All four-point bend testing was done in room temperature.

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155 L Load Span P/2 P/2 L/4 L/4 Support Span P/2 P/2 Figure 7-2. Schematic diagram of the four-point bend testing fixture for the flexural strength measurement. For the Instron machine used, the load span is 20 mm, while the support span is 40 mm Results and Discussion The weight changes as a result of exposure in different atmosphere are shown in Figure 7-3. As expected, active oxidation was observed during exposure at argon environment. This is manifested by the weight losses in the materials studied. Active oxidation in SiC results in evolution of CO and SiO gas.114 Exposure in air and water-containing environment showed that the material gains weight. Passive oxidation of SiC is usually observed in environment containing high partial pressure of oxygen.114 The weight gain observed was a result of formation of SiO2 during passive oxidation. Figure 7-3. Weight changes of the materials studied as a result of exposure in different atmosphere.

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156 The fracture strengths of the SiC ceramics investigated are shown in Figure 7-4. The fracture strengths of the polished specimens were included to provide a baseline for comparison. As shown in Figure 7-4 (indented condition), introduction of a semi-elliptical surface cracks on the specimens reduced the fracture strength. Average strength reduction upon indentation in SS SiC, ABC-SiC, and YAG-SiC are 71, 48 and 65%, respectively. Figure 7-4. Fracture strengths of the SiC materials before and after crack healing. Fracture strengths of the SiC materials prior to crack healing are shown by the polished and indented conditions. Fracture strengths after crack healing are shown by the argon, air and water conditions. The fracture strengths of the materials after crack healing were also shown in Figure 7-4. Argon, air and water conditions are the respective atmosphere where the crack healing experiments were performed. Indications of crack healing can be gleaned from the comparison of the fracture strengths measured from the indented samples and the respective exposure environment. As shown, increase in the fracture strengths of all the materials investigated after crack healing was observed. In the case of SS-SiC, the resulting increase in fracture strength upon crack healing in argon, air and water

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157 environments are 31, 61, and 216 % respectively. Although there was a large increase in fracture strength upon crack healing in water-containing environment, the achieved strength was still below the fracture strength of the polished material. In the case of ABC-SiC, the observed fracture strength upon crack healing in argon, air and water vapor environment are 24, 37, and 74 %, respectively. In the case of YAG-SiC, the respective fracture strength increases are 50, 153, and 191 %. In this material, the fracture strength measured upon crack healing in water vapor containing environment is even higher than the fracture strength measured in the polished specimen. It has to be noted, however, that the fracture strength reported for this material was only based on one specimen where the fracture was observed to initiate at the indentation site. Based on the results presented in Figure 7-4, maximum fracture strength recovery can be achieved in a water vapor-containing crack healing environment. Aside from the fracture strength, crack sizes and fracture toughness of the SiC ceramics were also measured. The cracks observed from the fractured surfaces of the SiC ceramics are generally semi-elliptical. Changes in crack appearance depending on the crack healing environment, however, were observed. As shown in Figure 7-5, the cracks in the fracture surfaces appear more semi-circular than semi-elliptical in materials crack-healed in air atmosphere. A semi-elliptical configuration, however, was used in all the calculations. The important dimensions measured are shown in Figure 7-5(a). The results of the crack radius calculation based on the long axis diameter and short axis radius of the semi-ellipse are presented in Figure 7-6. The crack radius was calculated by assuming a semi-circular crack of equal area, i.e., crack radius c = ab Comparison of the crack radius of the indented and crack-healed specimen showed a

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158 decrease in crack radius upon exposure in different atmosphere. ABC-SiC and YAG-SiC showed a very large decrease in crack size upon exposure in argon atmosphere. Figure 7-5. SEM of fracture surfaces in materials exposed in air. a) surface crack in SS-SiC and schematic of the dimensions measured. 2b is the diameter of the semi-major axis and a is the radius of the semi-minor axis of the ellipse. b) surface crack in ABC-SiC. c) surface crack in YAG-SiC. Figure 7-6. Crack radius calculated from the measured long axis diameter and short axis radius. Fracture toughness of the materials studied in indented and crack healed conditions is shown in Figure 7-7. The fracture toughness was calculated from the crack radius and the fracture strength of each specimen using the equation for brittle fracture given by, cY K IC (7-1)

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159 where KIC is the fracture toughness, is the fracture strength, Y is the stress intensity shape factor for the origin and c is the radius of surface crack. The stress intensity factor (Y) value used in all the calculations performed was 1.29 for semi-elliptical crack, following the ASTM standard C 1322 for fractography and characterization of fracture origins in advance ceramics.115 Figure 7-7. Fracture toughness calculated from the crack radius of the SiC materials. Figure 7-7 also indicates that toughening in the SiC materials studied can also be achieved through heat treatment in air and water vapor-containing environment. In the case of SS-SiC, an increase in fracture toughness in air and water vapor-containing environment of 37 and 149 %, respectively, were observed. ABC-SiC and YAG-SiC have shown an increase of 37 and 48 %, and 36 and 54 %, respectively, in the same environment. The increase in toughness in ABC-SiC and YAG-SiC in these environments is understandable due to the additional formation of SiO2 upon oxidation. On the other hand, the enormous increase in the fracture strength of SS-SiC in water vapor-containing environment seems to be baffling. It is, however, possible that formation of secondary phases may have occurred in this material during heat treatment.

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160 And as shown in previous section, formation of secondary phases (with a different thermal expansion coefficient) could result in presence of residual stress. The presence of residual stress in the material could result in a large increase in toughness, as shown in the results of experiments in -SiC in the previous chapter. Conclusions It was demonstrated that improvement in fracture strength and toughness through an additional heat treatment process is possible in SS-SiC, ABC-SiC and YAG SiC ceramics. Exposure in argon, air and water vapor-containing environments resulted in strength recovery in all materials studied. Decreases in crack sizes were also observed upon crack healing. Toughening in SS-SiC in water vapor containing environment due to the possible formation of secondary phases was also demonstrated.

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CHAPTER 8 CONCLUSION The grain boundary and triple junction phase characteristics in ABC-SiC were observed to vary with the amount of Al additive. The grain boundary and triple junction phases are generally Al-O-rich and at higher Al addition (1 and 4 wt. %), were found to be in the form of Al2O3. The triple junction phase was observed to form upon addition of aluminum. Crystalline triple junction phase, however, were only observed from 1 to 4 wt %. Al additions. The intergranular films only forms upon addition of 1.5 and 4 wt. % Al, and are amorphous at these levels of Al additives. The mechanical properties of ABC-SiC were observed to correlate with the formation of secondary phases. A decrease in hardness coincidental with the formation of triple junction phase in the material was observed. A two-fold toughness increase was observed upon the formation of intergranular films. The formation of intergranular film also induces the change in fracture from transgranular to an intergranular mode. The observed change in mode of fracture and toughening in the ABC-SiC system was attributed to the residual stress in the material. Residual stress arises due to the thermal expansion mismatch between the SiC grains and the Al2O3 present in the grain boundary (as intergranular films) and triple junction pockets. Investigations performed on SiC sintered with equimolar Al additions revealed that the intergranular films and triple junction phases was not affected by the choice of the source of Al. The intergranular films composition remains primarily Al-O-rich. The film also remains amorphous in each specimen and the thicknesses are ~ 1 nm. The triple 161

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162 junction phases are also similarly crystalline Al2O3. The mechanical properties of these materials, however, show some variations. Highest toughness was observed for -SiC-1.65Al while highest hardness was observed in -SiC-3.1Al2O3. These variations in the mechanical properties can be attributed to the microstructure observed. The grain sizes in -SiC-3.1Al2O3 are generally smaller than those observed in the other material, explaining the hardness value. The microstructure in -SiC show presence of elongated grains resulting in additional toughening. The microstructure in SiC-2.5AlN is generally equiaxed. Addition of B4C and C generally resulted in increased grain growth. The grain growth was more patent in -SiC-1.65Al-0.5B4C-2C due to the cooperative effect of Al, B and C. Increase in the aspect ratio of the elongated grains, however, did not occur, thus the toughness in this material decreased. The hardness is also lower than those observed in the other material. The grain growth in AlN containing material (-SiC-2.5AlN-0.5B4C and -SiC-2.5AlN-0.5B4C-2C) was not as pervasive as observed in Al-containing ones. This is due to the inhibitive effect of nitrogen on grain growth in SiC ceramics. The grain boundary and triple junction phases were also influenced by the addition of B and C. The chemistry observed changed from primarily Al-O-rich to one exhibiting presence of B (Al-O-B-rich composition). Presence of other secondary phases (B-C, C-N, B-N-rich composition) was also observed. The effect of this change in chemistry on the mechanical properties, however, may have been overshadowed by the effect of the microstructure (grain growth). Fabrication of SiC ceramics with SiC-N like properties have been shown to be achievable in -SiC sintered with AlN additive. Microstructure achieved in SiC-AlN

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163 composition are generally similar than those observed in SiC-N. The toughness of -SiC-5AlN hot pressed at 2200C is just below that of SiC-N. Further annealing of this material would then achieve the toughness and hardness desired. Use of Al additive in achieving SiC-N like properties is not impossible. However, several precautions and fabrication steps need to be added to the process to achieve microstructure similar to SiC-N. The effect of Al addition in sintering -SiC is to promote growth of elongated grains, especially if hot pressing at higher temperature (2100 to 2200C). The resulting microstructure in -SiC-Al system generally contains elongated grains in a matrix of fine equiaxed grains. In the case of AlN, the resulting grains are generally equiaxed. Increased amount of AlN addition results in grain growth retardation. Heat treatments at 1300C for three hours of selected SiC specimens (SS-SiC, YAG-SiC and ABC-SiC) have shown that improvement in fracture strength and toughness in the manufactured SiC specimens are possible. Exposure in argon, air and water vapor-containing environments resulted in strength recovery in all materials studied. Decreased in crack sizes were also observed upon crack healing. Toughening in SS-SiC in water vapor containing environment due to the possible formation of secondary phases have also been observed.

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LIST OF REFERENCES 1. N. P. Padture, In-situ Toughened Silicon Carbide, J. Am. Ceram. Soc., 77 [2] 519-23 (1994). 2. J. J. Cao, W. J. Moberlychan, L. C. De Jonghe, C. J. Gilbert and R. O. Rithcie, In-situ Toughened Silicon Carbide with Al-B-C Additives, J. Am. Ceram. Soc., 79 [2] 461-69 (1996). 3. G-D Zhan, M. Mitomo and Y-W Kim. Microstructural Control for Strengthening of Silicon Carbide Ceramics, J. Am. Ceram. Soc., 82 [10] 2924-26 (1999). 4. G-D. Zhan, R-J. Zie, M. Mitomo and Y-W. Kim. Effect of -toPhase Transformation on the Microstructural Development and Mechanical Properties of Fine-Grained Silicon Carbide Ceramics, J. Am. Ceram. Soc., 84 [5] 945-50 (2001). 5. S. K. Lee and C. H. Kim. Effects of -Sic versus -SiC Starting Powders on Microstructure and Fracture Toughness of SiC Sintered with Al2O3-Y2O3 Additives, J. Am. Ceram. Soc., 77 [6] 1655-58 (1994). 6. J-Y. Kim, Y-W. Kim, M. Mitomo, G-D. Zhan and J-G. Lee. Microstructure and Mechanical Properties of -Silicon Carbide Sintered with Yttrium-Aluminum Garnet and Silica, J. Am. Ceram. Soc., 82 [2] 441-44 (1999). 7. M. Nader, F. Aldinger and M. J. Hoffmann. Influence of the /-SiC Phase Transformation on Microstructural Development and Mechanical Properties of Liquid Phase Sintered Silicon Carbide, J. Mater. Sci., 34 1197-1204 (1999). 8. K. Biswas, G. Rixecker, I. Wiedmann, M. Schweizer, G. S. Upadhyaya, and F. Aldinger. Liquid Phase Sintering and Microstructure-Property Relationships of Silicon Carbide Ceramics with Oxynitride Additives, Mat. Chem. Phys., 67 180-191 (2001). 9. S-G. Lee, Y-W. Kim and M. Mitomo. Relationship Between Microstructure and Fracture Toughness of Toughened Silicon Carbide Ceramics, J. Am. Ceram. Soc., 84 [6] 1347-53 (2001). 10. Y. Zhou, K. Hirao, Y. Yamauchi and S. Kanzaki. Tailoring the Mechanical Properties of Silicon Carbide Ceramics by Modification of the Intergranular Phase Cemistry and Microstructure, J. Eur. Ceram. Soc., 22 2689-96 (2002) 164

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165 11. G-D. Zhan, M. Mitomo, H. Tanaka and Y-W. Kim. Effect of Annealing Conditions on Microstructural Development and Phase Transformation in Silicon Carbide, J. Am. Ceram. Soc., 83 [6] 1369-74 (2000). 12. C-W. Jang, J. Kim and S-J. L. Kang. Effect of Sintering Atmosphere on Grain Shape and Grain Growth in Liquid-Phase-Sintered Silicon Carbide, J. Am. Ceram. Soc., 85 [5] 1281-84 (2002. 13. L. K. L. Faulk. Microstructural Development During Liquid Phase Sintering of Silicon Carbide Ceramics, J. Eur. Ceram. Soc., 17 983-994 (1997). 14. A. L. Ortiz, T. Bhatia, N. P. Padture and G. Pezzotti. Microstructural Evolution in Liquid-Phase-Sintered SiC: Part III, Effect of NitrogenGas Sintering Atmosphere, J. Am. Ceram. Soc., 85 [7] 1835-40 (2002). 15. N. P. Padture and B. R. Lawn, Toughness Properties of a Silicon Carbide with an In situ Induced Heterogeneous Grain Structure, J. Am. Ceram. Soc., 77 [10] 2518-22 (1994). 16. Y-W. Kim, M. Mitomo, and H. Hirotsuru, Grain Growth and Fracture Toughness of Fine-Grained Silicon Carbide Ceramics, J. Am. Ceram. Soc., 78 [11] 3145-48 (1995). 17. M. A. Mulla and V. D. Krstic, Mechanical Properties of -SiC Pressureless Sintered with Al2O3 Additions, Acta Metall. Mater., 42 [1] 303-308 (1994). 18. M. Flinders, D. Ray, A. Anderson, and R. A. Cutler. High-Toughness Silicon Carbide as Armor, J. Am. Ceram. Soc., 88 [8] 2217-2226 (2005). 19. A. Ezis. Monolithic, Fully Dense Silicon Carbide Material, Method of Manufacturing, and End Uses, U. S. Patent 5,372,978 (Dec. 13, 1994). 20. B. Matchen. Applications of Ceramics in Armor Products, Key Eng Mat., 122-124 333-342 (1996). 21. M. Mitomo, T. Nishimura, and M. Tsutsumi. Crack Healing in Silicon Nitride and Alumina Ceramics, J. Mater. Sci. Lett., 15 1976-1978 (1996). 22. H. E. Kim and A. J. Moorhead. Effects of Active Oxidation on the Flexural Strength of -Silicon Carbide, J. Am. Ceram. Soc., 73 [7] 1868-72 (1990). 23. R. P. Jensen, W. E. Luecke, N. P. Padture and S. M. Wiederhorn. High-Temperature Properties of Liquid-Phase-Sintered -SiC, Mat. Sci. Engg., A282 109-114 (2000). 24. G. Rixecker, I. Wiedmann, A. Rosinus and F. Aldinger. High-Temperature Effects in the Fracture Mechanical Behavior of Silicon Carbide Liquid-Phase Sintered with AlN-Y2O3 Additives, J. Eur. Ceram. Soc., 21 1013-19 (2001).

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BIOGRAPHICAL SKETCH Edgardo L. Pabit began grain boundary and triple junction characterization of liquid phase sintered SiC research under the supervision of Dr. Darryl Butt at the Advanced Ceramics Materials Laboratory in the summer of 2004. He previously worked on the area of high-temperature oxidation for which he received his MS degree in 2003. He entered the University of Florida materials science and engineering Ph.D. program in 2000 and in his first year worked as a teaching assistant assigned to Prof. Robert Dehoff. Prior to graduate school at UF, he worked as a teaching associate at the National Institute of Physics, University of the Philippines, at Diliman, Quezon City, Philippines, where he received his undergraduate degree, B.S. in aplied physics, in 1996. He is proud to be born and raised in a small town by the sea at Barangay Sumilang, Calauag, Quezon, Philippines. 173


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Title: Grain Boundary and Triple Junction Chemistry of Silicon Carbide Sintered with Minimum Additives for Armor Applications
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Permanent Link: http://ufdc.ufl.edu/UFE0012140/00001

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Title: Grain Boundary and Triple Junction Chemistry of Silicon Carbide Sintered with Minimum Additives for Armor Applications
Physical Description: Mixed Material
Copyright Date: 2008

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GRAIN BOUNDARY AND TRIPLE JUNCTION CHEMISTRY OF SILICON
CARBIDE SINTERED WITH MINIMUM ADDITIVES FOR ARMOR
APPLICATIONS
















By

EDGARDO L. PABIT


A DISSERTATION PRESENTED TO THE GRADUATE SCHOOL
OF THE UNIVERSITY OF FLORIDA IN PARTIAL FULFILLMENT
OF THE REQUIREMENTS FOR THE DEGREE OF
DOCTOR OF PHILOSOPHY

UNIVERSITY OF FLORIDA


2005

































Copyright 2005

by

Edgardo L. Pabit

































This document is dedicated to my loving wife.















ACKNOWLEDGMENTS

I would like to thank my adviser, Dr. Darryl Butt, for his guidance, patience and

support during the course of my studies. I would also like to thank Dr. John Mecholsky

Jr., Dr. Paul Holloway, Dr. Wolfgang Sigmund and Dr. Mark Orazem, my committee

members, for taking the time to answer my questions and for the helpful comments on

my dissertation. I would like to acknowledge Ceramatec, Inc. for the funding of this

project and for providing the mechanical property data for this dissertation. I would also

like to acknowledge Kerry Seibien of MAIC at the University of Florida and Dr. Helge

Heinrich of MCF at the University of Central Florida. Their help on performing the

HRTEM and EFTEM studies presented in this dissertation is greatly appreciated.

Finally, I would like to thank Dr. Butt's group members for their support and for

providing a semblance of sanity during the long hours in the laboratory.
















TABLE OF CONTENTS

page

A C K N O W L E D G M E N T S ................................................................................................. iv

LIST OF TABLES ................................................................ ........ viii

LIST OF FIGURES ......... ......................... ...... ........ ............ ix

A B S T R A C T .........v.................................... ....................... ................. xv

CHAPTER

1 IN T R O D U C T IO N ............................................................................. .............. ...

2 LIQUID PHASE SINTERING AND MECHANICAL PROPERTY
M ODIFICATION S IN SILICON CARBIDE ........................................ ...................6

Sintering Additives used in Densification of SiC.................................... ..................7
Effects of Processing Parameters in LPS-SiC ...................................... ...............10
Effect of Sintering Tim e and Tem perature..................................... ................ 10
Effect of Choice of Additives.................. ..................................12
Effect of the Starting SiC Pow ders................................. ........................ 15
Effect of Sintering Atmosphere.......................... ......... .................. 17
Effect of A nnealing C onditions............................ ......................... ..................19
Microstructure and Mechanical Property Modification in Sintered SiC ....................20
Fracture Toughness Enhancem ent in SiC.................................................... .... 21
Crack Healing as a Tool for Mechanical Property Improvement in SiC ...........26

3 STATEMENT OF THE OBJECTIVE ............................................ ............... 31

4 M ATERIALS AND M ETHODS ........................................ ......................... 33

M a te ria ls ........................................................................... 3 3
C characterizations ................................................... ................... ...... .... .. 36
M echanical Properties and SiC Polytypes............................... ............... 36
M ic ro stru ctu re ............................................................................................... 3 7
C hem ical A naly sis........... ........................................................ ...... .... ... ..37
Transmission Electron M icroscopy ................ .............................................. 38
TEM sam ple preparation .................................... ......................... .. ......... 39
High resolution TEM (HRTEM) ..............................40









Energy dispersive spectroscopy (ED S) ....................................................... 41
Energy filtered transmission electron microscopy (EFTEM) ....................43

5 RESULTS OF MECHANICAL PROPERTY, GRAIN BOUNDARY, AND
TRIPLE JUNCTION CHARACTERIZATION ............. .........................................46

P-SiC Sintered with 0.6 wt.% B, 2 wt.% C and 0 4 wt.% Al (ABC-SiC) ..............46
Processing and Polytypes ............................................. ............ ............... 46
Microstructure and Mechanical Properties.................... .... ............... 48
Grain Boundary, Triple Junction and Other Secondary Phases ........................51
Grain boundary width and crystallinity.....................................................51
Grain boundary, triple junction and secondary phase composition ............53
p -S iC 0 .6B -2 C ............................................ ................ 54
-SiC 0.5A l 0.6B 2C .............................. ... ................ ............... 55
P-SiC-1A1-0.6B-2C .......................................................... ...... ...... .... 59
P-SiC 1.5Al 0.6B 2C ............................ .. ..................... .....61
-SiC 4A 1 0.6B 2C ............ ........................................ ............... 65
a-SiC Sintered w ith A l A dditive ........................................ .......................... 68
Processing and Polytypes ............................................. ............ ............... 68
Microstructure and Mechanical Properties........................ .................70
Grain Boundaries, Triple Junctions and Secondary Phases .............................76
a-SiC 1.65 wt. % Al hot pressed at 19000C..........................................76
a-SiC 1.65 wt. % Al hot pressed at 21000C ..........................................80
a-SiC 1.65 wt. % Al hot pressed at 22000C.....................................83
a-SiC 3.3 wt. % Al hot pressed at 20000C ............................................85
a-SiC 3.3 wt. % Al hot pressed at 22000C ...............................90
a-SiC 1.65 wt. % Al 0.5 wt. % B4C 2 wt. % C ................................93
a-SiC Sintered w ith A 1N A dditive................................................... .....................96
Processing and Polytypes ............................................. ............ ............... 96
Microstructure and Mechanical Properties........................ .................98
Grain Boundaries, Triple Junctions and Secondary Phases ............................101
a-SiC 2.5 wt. % A1N hot pressed at 21000C........................................101
a-SiC 2.5 wt. % A1N hot pressed at 22000C........................................104
a-SiC 5 wt. % A1N hot pressed at 2000C ...........................................107
a-SiC 5 wt. % A1N hot pressed at 2200C ...........................................109
a-SiC 2.5 wt. % A N 0.5 wt. % B4C .....................................................111
a-SiC 2.5 wt. % A1N 0.5 wt. % B4C 2 wt. % C ..............................115
SiC-N a-SiC-3.1A1203, a-SiC-0.5B4C-2C ......................................... ...............117
Processing and Polytypes ......................................................... ............... 118
M icrostructure and M echanical Properties....................................................... 119
Grain Boundaries, Triple Junctions and Secondary Phases ...........................121
SiC-N ....................................................... ..... ......... 121
a-SiC 3.1 w t. % A l20 3..................................................... ....................124
a-SiC 0.5 w t. % B 4C 2 w t. % C ...................................... ............... 126









6 SUMMARY AND DISCUSSION: MECHANICAL PROPERTY, GRAIN
BOUNDARY, AND TRIPLE JUNCTION CHARACTERISTICS ........................129

P-SiC Sintered with 0.6 wt.% B, 2 wt.% C and 0 4 wt.% Al (ABC-SiC) ...........129
Influence of Residual Stress ....... ................................... ................. 130
Absence of Intergranular Films in Materials with Al Additive in Excess of
S olu b ility L im it ............................................................................... 13 3
SiC Ceramics Fabricated from a-SiC Starting Powders.................... ............... 135
Comparison of Mechanical Properties and Grain Boundary and Triple
Junction Characteristics of SiC Sintered with Equimolar Al Addition .........136
Comparison of SiC Ceramics with Equimolar Al Addition..............................136
Effect of B and C A additions ............... .......... ............. ..................... 140
Comparison of a-SiC Ceramics Sintered with Al or A1N ............ ... ..................144
Effect of Hot Pressing Temperature and Amount of Additives ........................144
Achievement of SiC-N-like Properties................................... ............... 147

7 CRACK HEALING OF SELECT SPECIMEN......................................................151

M materials an d M eth od s .................................................................. ..................... 152
R results and D discussion ..................................... .......... .... .... .. ........ .. 155
C o n clu sio n s.................................................... ................ 16 0

8 C O N C L U SIO N ......... ....................................................................... ........ .. ..... .. 16 1

L IST O F R E FE R E N C E S ........................................................................ ................... 164

BIOGRAPHICAL SKETCH ............................................................. ............... 173
















LIST OF TABLES


Table pge

4-1 Compositions and processing conditions of the materials investigated .................34

4-2 Raw materials used in powder processing. ................................... ............... 35

5-1 Density and polytypes of SiC ceramics from P-SiC starting powders ...................47

5-2 Grain size, aspect ratio, hardness, toughness and observed fracture mode in SiC
ceramics with P-SiC starting powders...... ..................... ...............49

5-3 Grain boundary width and crystallinity of P-SiC with Al, B, and C additions........53

5-4 Density, chemical analysis and phase assemblage of a-SiC sintered with Al
additive e. ..............................................................................68

5-5 Mechanical properties of a-SiC sintered with Al additive .....................................73

5-6 Density, chemical analysis and phase assemblage of a-SiC sintered with A1N
additive e. ..............................................................................97

5-7 Mechanical properties of a-SiC sintered with A1N additive ..............................100

5-8 Density, chemical analysis and phase assemblage of SiC-N, a-SiC-A1203 and a-
SiC -0 .5B 4C -2 C ............................................................................ ............... 119

5-9 Hardness, toughness, and mode of fracture observed in SiC-N, a-SiC-3.1A1203
and a-SiC-0.5B4C-2C ................................................................ ...... 120

6-1 Summary of observations for SiC ceramics sintered from P-SiC starting
pow ders. ........................................................................... 130

6-2 Mechanical properties, grain boundary and triple junction characteristics of SiC
materials with equimolar aluminum addition............................................. 136

6-3 Mechanical properties, grain boundary and triple junction characteristics of SiC
sintered with Al or A1N addition ................................................ ........ ....... 145

7-1 Composition, processing conditions and mechanical properties of materials used
for crack healing.................................................................................................. 152
















LIST OF FIGURES


Figure p

4-1 Signals generated from high-energy electron beam/thin specimen interaction.....38

4-2 Schematic overview of the three-window technique................. .................44

5-1 Transmission electron micrographs of the microstructure observed in SiC
ceramics fabricated from a P-SiC starting powder.. .............................................48

5-2 High resolution transmission electron micrographs of grain boundaries in P-
SiC with Al, B and C additives. .................................. .............................. 52

5-3 STEM of a grain boundary and a secondary phase in P-SiC-0.6B-2C..................55

5-4 STEM of a triple junction in P-SiC-0.6B-2C. ........................................ ...........56

5-5 EFTEM of a triple junction and grain boundary in P-SiC-0.6B-2C. ..................56

5-6 STEM of a triple junction in P-SiC-0.5A1-0.6B-2C. .........................................57

5-7 STEM of a grain boundary in P-SiC-0.5Al-0.6B-2C. ..........................57

5-8 EFTEM of a grain boundary in P-SiC-0.5Al-0.6B-2C .............................58

5-9 STEM of a multiple grain junction in P-SiC-0.5Al-0.6B-2C containing
metallic impurities as shown by the corresponding EDS. ....................................58

5-10 STEM of a triple junction in P-SiC-1A1-0.6B-2C ............. .............. 60

5-11 EFTEM of a triple junction in P-SiC-1A1-0.6B-2C ......... ...............60

5-12 EFTEM of a grain boundary in P-SiC-1A1-0.6B-2C ........................... 61

5-13 STEM of a triple junction in P-SiC-1.5A1-0.6B-2C. ....... .................62

5-14 EFTEM of a triple junction in P-SiC-1.5A1-0.6B-2C. .........................63

5-15 STEM of a grain boundary in P-SiC-1.5Al-0.6B-2C .......................................63

5-16 EFTEM of a grain boundary in P-SiC-1.5Al-0.6B-2C .............................64









5-17 STEM of a grain boundary and associated bulk grains in P-SiC-1.5Al-0.6B-
2 C ...................................... ..................................................... 6 5

5-18 Residual carbon in triple junction of P-SiC-4Al-0.6B-2C..................................66

5-19 EFTEM of a grain boundary in P-SiC-4A1-0.6B-2C ........................................66

5-20 EFTEM of a multiple grain junction in P-SiC-4A1-0.6B-2C.............................67

5-21 STEM and EDS of secondary phases found in P-SiC-4A1-0.6B-2C. ....................67

5-22 Transmission electron micrographs of the microstructures observed in a-SiC
sintered w ith A l additive ............................................... ............................. 71

5-23 Optical micrographs of Vickers indentation in a-SiC sintered with aluminum
addition s .............................................................................73

5-24 SEM of fracture surfaces of a-SiC sintered with aluminum additions ................74

5-25 High resolution transmission electron micrographs of several grain boundaries
studied in a-SiC-1.65Al hot pressed at 1900 C.................................. ............... 76

5-26 STEM and EDS of a grain boundary and associated grains in a-SiC-1.65A1
hot-pressed at 1900 C ........................... .................... ......................... 77

5-27 STEM and EDS of a triple junction in a-SiC-1.65Al hot-pressed at 19000C. ....78

5-28 EFTEM of a triple junction and grain boundary in a-SiC-1.65Al hot pressed at
1 9 0 0 0 C .......................................................................... 7 8

5-29 EELS taken from a triple junction in a-SiC-1.65Al hot pressed at 19000C..........79

5-30 High resolution images and EDS of triple junctions in a-SiC-1.65Al hot
pressed at 19000C .................... ................ ................................. 79

5-31 HRTEM of several grain boundaries in a-SiC-1.65Al hot pressed at 21000C......80

5-32 EFTEM of a grain boundary in a-SiC-1.65Al hot pressed at 21000C..................81

5-33 STEM and EDS of triple junction in a-SiC-1.65Al hot pressed at 21000C. .........82

5-34 STEM of triple grain junction phase with lower dihedral angle in a-SiC-
1.65A 1 hot pressed at 21000C ................................... ......... .... ............ ... 82

5-35 EFTEM of a triple junction in a-SiC-1.65Al hot pressed at 21000C ............... 83

5-36 HRTEM of grain boundaries in a-SiC-1.65Al hot pressed at 22000C. .................84









5-37 STEM of a triple junction in a-SiC-1.65Al hot pressed at 22000C....................... 84

5-38 EFTEM of a triple junction and grain boundary in a-SiC-1.65Al hot pressed at
22000C ............................................................................85

5-39 HRTEM of grain boundaries in a-SiC-3.3Al hot pressed at 20000C ..................86

5-40 STEM of a grain boundary near a triple junction in a-SiC-3.3Al hot pressed at
20000C ............................................................................86

5-41 STEM of a triple junction in a-SiC-3.3Al hot pressed at 20000C.........................87

5-42 STEM of two grains with different composition found in a-SiC-3.3Al hot
pressed at 20000C. ................................... ... ... ... .. .. ............ 88

5-43 EFTEM of a triple junction in a-SiC-3.3Al hot pressed at 20000C.....................89

5-44 HRTEM of grain boundaries in a-SiC-3.3A1 hot pressed at 22000C ..................90

5-45 STEM of a grain boundary in a-SiC-3.3A1 hot pressed at 22000C .....................91

5-46 EFTEM of a grain boundary in a-SiC-3.3A1 hot pressed at 22000C ...................91

5-47 STEM and EDS of a triple junction containing primarily Al-rich phase and
metallic impurities in a-SiC-3.3A1 hot pressed at 22000C. ..................................92

5-48 STEM and EDS of Al-rich grain found in a-SiC-3.3A1 hot pressed at 22000C....92

5-49 EFTEM of a triple junction in a-SiC-3.3A1 hot pressed at 22000C.....................93

5-50 HRTEM of grain boundaries observed in a-SiC-1.65A1-0.5B4C-2C....................94

5-51 EFTEM of a grain boundary in a-SiC-1.65A1-0.5B4C-2C................... ................94

5-52 STEM of a triple junction in a-SiC-1.65A1-0.5B4C-2C ..... ..............95

5-53 STEM and EDS of other secondary phases in a-SiC-1.65A1-0.5B4C-2C. ............96

5-54 TEM of the microstructures observed in a-SiC sintered with A1N additive. ........99

5-55 Optical micrographs of Vickers indentation on a polished surface of a-SiC
sintered w ith A 1N additive. ...................................................................... ....... 100

5-56 SEM of fracture surfaces of a-SiC sintered with A1N additive........................... 101

5-57 HRTEM of grain boundaries observed in a-SiC-2.5A1N hot pressed at
2 10 0 0C ........................................................................ 10 2









5-58 STEM and EDS taken from a-SiC-2.5A1N hot pressed at 21000C..................... 102

5-59 EFTEM of a triple junction in a-SiC-2.5A1N hot pressed at 21000C................ 103

5-60 Low and high resolution image of a triple junction in a-SiC-2.5A1N hot
pressed at 2 100 C ........................ ...................... ... .. ....... .... ...........104

5-61 HRTEM of grain boundaries in a-SiC-2.5A1N hot pressed at 2200C. ..............105

5-62 STEM and EDS of grain boundary in a-SiC-2.5A1N hot pressed at 22000C......105

5-63 STEM of a triple junction in a-SiC-2.5A1N hot pressed at 22000C. ................... 106

5-64 EFTEM of a triple junction in a-SiC-2.5A1N hot pressed at 22000C..................06

5-65 HRTEM of grain boundaries in a-SiC-5AIN hot pressed at 2000C. ................. 107

5-66 STEM of triple grain junction in a-SiC-5AIN hot pressed at 20000C................. 108

5-67 EFTEM of triple junctions in a-SiC-5AIN hot pressed at 20000C...................... 109

5-68 HRTEM of grain boundaries in a-SiC-5AIN hot pressed at 2200C. .................10

5-69 STEM of a triple junction in a-SiC-5AIN hot pressed at 22000C. ...................... 110

5-70 EFTEM of a multiple grain junction and triple grain junction in a-SiC-5A1N
hot pressed at 22000C. ............................................ ... ...... .. ........ .. .. 111

5-71 HRTEM of several grain boundaries observed in a-SiC-2.5AlN-0.5B4C...........112

5-72 STEM of triple junction observed in a-SiC-2.5A1N-0.5B4C.............................113

5-73 EFTEM of a grain boundary observed in a-SiC-2.5AlN-0.5B4C..................113

5-74 EFTEM of a triple junction observed in a-SiC-2.5AIN-0.5B4C.......................... 114

5-75 STEM of secondary phases observed in a-SiC-2.5AIN-0.5B4C. ....................114

5-76 HRTEM of several grain boundaries in a-SiC-2.5AIN-0.5B4C-2C. ...................115

5-77 EFTEM of a triple junction and grain boundary in a-SiC-2.5AIN-0.5B4C-2C...116

5-78 Another EFTEM taken from a different triple junction in this material..............16

5-79 STEM of several secondary phases observed in a-SiC-2.5AIN-0.5B4C-2C.......117

5-80 TEM of microstructure observed in SiC-N, a-SiC-0.5B4C-2C and a-SiC-
3 .1A 120 3 ...................................................................... 12 0









5-81 Optical micrographs of Vickers indent on a polished surface of SiC-N, a-SiC-
0.5B4C-2C and a-SiC-3.1A1203 ............................................ .......... ................ 120

5-82 SEM of fracture surfaces of SiC-N, a-SiC-0.5B4C-2C and a-SiC-3.1A1203 .....121

5-83 HRTEM of several grain boundaries observed in SiC-N. .............. ............... 122

5-84 STEM of a grain boundary in SiC-N ................................... .......................... 122

5-85 EFTEM of a grain boundary in SiC-N ............................. 123

5-86 STEM of a triple junction in SiC-N .... ............................. 123

5-87 EFTEM of a triple junction in SiC-N. .................................... .................124

5-88 HRTEM of several grain boundaries in a-SiC-3.1A1203............................... 125

5-89 STEM of a triple junction in a-SiC-3.1A1203 .................................................125

5-90 EFTEM of a grain boundary and a triple junction in a-SiC-3.1A1203 ...............126

5-91 HRTEM of grain boundaries observed in a-SiC-0.5B4C-2C. ............................ 127

5-92 STEM of a triple junction in a-SiC-0.5B4C-2C .................................................127

5-93 STEM of a triple junction containing residual carbon and metallic impurities
in a-SiC -0.5B 4C -2C. ...................... .. .... ........................................... 128

5-94 EFTEM of triple junction and grain boundary in a-SiC-0.5B4C-2C...................128

6-1 Calculated residual stresses due to the presence of secondary phase ................132

6-2 TEM of microstructure observed in a-SiC sintered with equimolar Al
addition ..........................................................................138

6-3 SEM of crack profiles in SiC materials with equimolar Al addition..................139

6-4 SEM of crack profiles in materials containing B4C and C. ........................... 141

6-5 TEM of microstructure observed in a-SiC sintered with further addition of
B 4C and C ........................................................................142

6-6 SEM of crack profiles in a-SiC sintered with Al additives. ........................... 147

6-7 SEM micrographs of crack path in SiC-N and a-SiC sintered with A1N
addition ..........................................................................148

7-1 Optical micrographs of the Vickers indentation using a 98 N load.....................153









7-2 Schematic diagram of the four-point bend testing fixture for the flexural
strength m easurem ent. ............................................... .............................. 155

7-3 Weight changes of the materials studied as a result of exposure in different
atm o sph ere. ...................................................................... 155

7-4 Fracture strengths of the SiC materials before and after crack healing. ............156

7-5 SEM of fracture surfaces in materials exposed in air. ................... ................158

7-6 Crack radius calculated from the measured long axis diameter and short axis
radiu s .......................................................................... .... .. .... 158

7-7 Fracture toughness calculated from the crack radius of the SiC materials..........159















Abstract of Dissertation Presented to the Graduate School
of the University of Florida in Partial Fulfillment of the
Requirements for the Degree of Doctor of Philosophy

GRAIN BOUNDARY AND TRIPLE JUNCTION CHEMISTRY OF SILICON
CARBIDE SINTERED WITH MINIMUM ADDITIVES FOR ARMOR
APPLICATIONS

By

Edgardo L. Pabit

December 2005

Chair: Darryl Butt
Major Department: Materials Science and Engineering

Silicon carbide containing a minimum amount of additives for armor application

was fabricated by hot pressing. Microstructural development, phase information,

mechanical properties, and triple junction and grain boundary chemistry were described.

Correlation between the characteristics of the secondary phases and mechanical

properties was also presented. The grain boundary and triple junction phases were

characterized using high resolution transmission electron microscopy, energy dispersive

spectroscopy and energy filtered transmission electron microscopy.

The grain boundary and triple junction phase characteristics in silicon carbide

sintered with the addition of aluminum, boron and carbon varied with the amount of

aluminum additive. Silicon carbide sintered with the addition of boron and carbon only

did not form triple junction and grain boundary phase, while addition of 0.5 and 1 wt. %

aluminum resulted in the formation of amorphous Al-O-rich triple junction phase.









Crystalline A1203 formed at the triple junction pockets upon addition of 1.5 and 4 wt. %

Al. Formation of amorphous intergranular films, with thickness of approximately Inm,

was also observed at these amounts of aluminum additive. An increase in toughness,

from 2.7 to 6.1 MPa m/2, accompanied the presence of this grain boundary phase,

which was also coincidental with the transition from transgranular to intergranular mode

of fracture. The increased toughness and the change in fracture mode from transgranular

to intergranular fracture were attributed to the residual stresses in the intergranular films.

The triple junction and grain boundary phase characteristics in silicon carbide

sintered with Al, A1N, and A1203 addition did not vary with the source of Al additive.

The triple junction phases were observed to be crystalline A1203 while the grain boundary

phase composition remained primarily Al-O-rich. The intergranular film also remained

amorphous with thicknesses of approximately 1 nm. Toughness value of 6.8 MPa*ml/2

was attained by hot pressing at 22000C and using 3.3 wt. % Al additive. The resulting

microstructure in this material consisted of elongated grains in a matrix of fine equiaxed

grains. The use of A1N additive resulted in a retarded the grain growth and equiaxed

grain morphology. Addition of B4C and C resulted in increased grain growth.

Heat treatment at 13000C for 3 hours of select SiC specimens showed that crack

healing was possible. Exposure to argon, air and water vapor-containing environments

resulted in strength recovery in all materials studied. Decreases in crack sizes were also

observed upon crack healing. Toughening in solid state sintered SiC in water vapor-

containing environment, due to the possible formation of secondary phases, was also

observed.














CHAPTER 1
INTRODUCTION

Liquid-phase-sintered silicon carbide (LPS-SiC) has attracted increasing interest

for its ability to form an in-situ toughened material and potentially superior mechanical

properties relative to solid-state-sintered SiC.1'4 Another incentive that makes LPS-SiC

technologically attractive is the ability to slightly tailor the resulting mechanical

properties through improved control of the microstructure of the sintered materials.3 1

Depending on the types and amount of additive used, types of starting powder, sintering

atmosphere, and processing conditions, the resulting grain morphology and property of

the materials varied. For instance, it is generally accepted that higher temperature and

longer sintering time result in a pronounced grain growth.11-13 The use of nitrogen

overpressure during sintering usually results in an equiaxed grain morphology regardless

of the types of starting powder used,14 whereas use of argon results in equiaxed or

elongated grain morphology depending on the starting powder.4'12 In the case of a-SiC

starting powder, the resulting grain morphology is generally equiaxed.6'7 For a P-SiC or

P-SiC containing a-SiC seeds, the resulting grains are generally elongated due to the P-to-

a transformation during sintering or annealing.5'7 It has been proposed that elongated

microstructure usually increases the fracture toughness in SiC by crack bridging or crack

deflection.15-17

The incremental increase in fracture toughness of LPS-SiC is generally promoted

by the elongated grains with relatively weak intergranular grain boundary and triple

junction phases.1's17 The intergranular phases and triple junction composition are









generally controlled by the type of additives used. A fracture toughness of 8 MPa*m1/2

has been reported in SiC with additions of yttrium aluminum garnet (YAG, Y3Al5012).

The amount of additives used in this system generally exceeds 10 wt. %. A slightly

higher fracture toughness of 9 MPa*ml/2 was also reported for SiC sintered with

aluminum, boron and carbon (ABC-SiC) additives. In this case, higher toughness was

achieved with additives of only about 5 wt. %.2

One of the applications envisioned for SiC with improved properties is in ceramic

armor.18 Currently, the state of the art ceramic armor material against heavy threats is a

silicon carbide densified with A1N, 19 produced by Cercom, Inc. under the trade name

SiC-N. In armor applications, the structural ceramic component is desired to have both

high hardness and toughness to perform its intended function. The armor is designed to

"defeat" the projectile at the interface by increasing its dwell time.20 In ballistic tests,

long dwell times are achieved for high hardness ceramics. And at the same level of

hardness, higher toughness leads to an improved resistance to penetration of the

projectile.18 Generally, high hardness results in low toughness and vice versa. The

manufacturing challenge, therefore, is to find an optimal balance of both properties.

Literature data generally show that higher fracture toughness is associated with

intergranular mode of fracture in SiC ceramics. Intergranular mode of fracture is usually

observed in the presence of intergranular films and triple junction phases in the sintered

material. The presence of these secondary phases, however, diminishes the hardness in

the final product. Investigation of the interdependence of these variables (grain

boundary, triple junction chemistry, microstructure and processing conditions) would

therefore be important in achieving a hard and toughened SiC.









In this dissertation, several approaches were made in order to understand the

interdependence of the grain boundary and triple junction chemistry, microstructure,

processing conditions and the resulting mechanical properties of SiC ceramics. Several

sets of materials were fabricated with minimum additives, and characterized with respect

to the aforementioned variables.

The ABC-SiC system has been of interest due to the high toughness available in

this material. Five specimens were made and studied to assess the effect of aluminum

concentration on the grain boundary and triple junction chemistry and the resulting

mechanical properties. The materials were fabricated from a P-SiC starting powder and

hot pressed with the addition of 0.6 wt. % B, 2 wt. % C and varying amounts of Al (0 to 4

wt. %). Results obtained from these materials were then used as a basis for comparison

to and development of succeeding test matrices.

The effect of source of Aluminum on the resulting mechanical properties and grain

boundary and triple junction characteristics was also investigated. The materials in this

case were fabricated from a-SiC starting powders and hot pressed with the addition of Al,

A1N and A1203. The amount of additive used was made higher than the amount of Al in

ABC-SiC which yielded higher toughness on the materials investigated. For the purpose

of comparison, the amount of Al was also made to correspond (or be equimolar with) to

the amount of Al added in the commercially available SiC-N.

The effect of B and C on the densification and microstructure of SiC fabricated

from P-SiC starting powders was already known. In the case of SiC fabricated from a-

SiC starting powders, however, the effect of B and C addition has not been properly









studied. To address this issue, incorporation of B4C and C on materials fabricated from

a-SiC powders sintered with Al or A1N addition was also investigated.

Tailoring of the microstructure of SiC ceramics to achieve mechanical properties

similar to that of state of the art ceramic armor material (SiC-N) was also attempted. a-

SiC with a different amounts of Al (1.65 and 3.3 wt. %) or A1N (2.5 and 5 wt.%)

additions was hot-pressed at different temperatures (19000, 21000 and 22000C).

Comparison of the triple junction and grain boundary phase characteristics and

mechanical properties was performed and methods for possible enhancement were also

presented.

Improvement of fracture strength in SiC has been reported possible through crack

healing. Crack healing can be accomplished via an additional heat treatment. In this

dissertation, crack healing of select specimen was also presented. The heat treatments

were carried out in air, argon and water-vapor containing environment.

The work presented in this dissertation was performed in cooperation with

Ceramatec, Inc. Ceramatec, Inc. fabricated all the specimens studied. The mechanical

property data presented, including the fracture mode and grain size measurements, were

based on the characterization performed by Ceramatec, Inc. Grain boundary and triple

junction characterizations were performed at the University of Florida and the University

of Central Florida. Crack healing experiments were performed at the University of

Florida.

Presentation of this dissertation will be as follows: Chapter 2 will cover the

literature related to the liquid-phase sintering of SiC including discussions on the effect

of starting powders, processing conditions, amount and type of additives and mechanical






5


property modification; Chapter 3 will present the objective of this study; Chapter 4 will

deal with the materials and methods used in this study; Chapter 5 will present the results

of mechanical property, grain boundary and triple junction chemistry characterizations;

Chapter 6 will show discussion of the results presented; Chapter 7 will deal with the

results and discussion of the crack healing experiments; and finally, a general conclusion

will be presented in Chapter 8.














CHAPTER 2
LIQUID PHASE SINTERING AND MECHANICAL PROPERTY MODIFICATIONS
IN SILICON CARBIDE

Fabrication of SiC ceramics through liquid phase sintering made possible the use of

SiC in demanding structural applications, such as armor component and high

temperature-heat exchanger materials.19,22-27 These two applications require different

sets of material properties. In the case of armor component applications, researchers are

in agreement that the ceramic material should possess relatively high hardness and

toughness, 8-20, 26-27 although these two properties in and of themselves are only part of

the desirable attributes. High hardness and toughness are achieved in part through the

judicious selection of fabrication parameters and use of liquid phase forming additives.

For high temperature applications, however, the high-temperature strength retention of

the ceramics should be increased. High temperature strength retention can be improved

through the incorporation of minimal secondary phase.22-25 Thus, the variable that gave

high toughness for ceramic armor becomes undesirable for SiC ceramics intended for

high temperature applications. This phenomenon clearly shows the need for

understanding the effect of processing variables on the microstructure that eventually

controls the mechanical properties desired for a given application.

To address this issue, discussion of the different additive system used for the

densification of SiC ceramics will be presented in this section. In addition, the effects of

several sintering parameters on the densification and resulting microstructure, and efforts









reported in the literature for improvement of the mechanical properties of SiC ceramics

will also be presented.

Sintering Additives used in Densification of SiC

The fabrication of dense SiC ceramics requires a sintering additive because of the

extremely low self diffusivities in the strongly covalently bonded a- and P-SiC

structures.28 An effective additive promotes densification and creates an environment

which inhibits the decomposition of the SiC powders at the high sintering temperature.3

In liquid phase sintering, the additive should be able to participate in liquid phase

formation and the resulting liquid phase should act as a mass transport media during

densification. To perform this function, the liquid phase generated should be of sufficient

volume to allow complete wetting of the solid phase. Further, it can be desirable for the

solid phase to have an appreciable solubility in the liquid to promote solution-

reprecipitation.13'29 These characteristics are dependent on the particular additives, the

relevant eutectic temperatures and the densification parameters such as sintering

atmosphere and temperature regime. The resulting densification rates, grain morphology

and microstructures, and mechanical properties of the sintered material are thus

controlled by the appropriate selection of fabrication parameters.

In an attempt to obtain full densification and remove processing flaws that

practically limit the strength and toughness of SiC ceramics, several additive systems

have been investigated in the literature.30-ss In most cases, the additive system is based

on aluminum, added as metallic aluminum, aluminum oxide or aluminum nitride, in

combination with other elements, such as B and C, or metallic oxides, such as Y203 and

rare-earth oxides. Aluminum has been primarily used since it has been observed to

facilitate liquid phase sintering in SiC. The presence of aluminum lowers the









densification temperature, induces the P-to-a phase transformations, promotes anisotropic

grain growth of SiC grains, forms amorphous grain boundary films, and produces various

secondary phases in triple junction pockets.53'54'58

The use of aluminum, in combination with boron and carbon, was extensively

studied by several investigators from the University of California, Berkeley.2'53-58 Their

work has contributed to what has become known as the aluminum-boron-carbon

containing SiC (ABC-SiC) system. The starting SiC powder in these studies is generally

of the P-SiC polytype and utilizes the P-to-a phase transformation to toughen the

resulting monoliths. Toughness values as high as 9.1 MPa*m1/2 and a 4-point bend

strength of around 660 MPa have been achieved.2 This toughness value is the highest

reported for SiC so far and was achieved at a relatively low sintering temperature of

19000C using a minimum amount of additives (3 wt.% Al, 0.6 wt.% B and 2 wt.% C).

Other liquid phase-forming additives used to densify SiC ceramics reported in the

literature include Al203, A1203-Y203, AIN-Y203, A1203-Y203-CaO, Al203-rare-earth

oxides and Al203-lanthanum series oxides.30-52 A2O3 is believe to react with SiO2, which

was always present on the surface of SiC particles, to form a liquid melt and, with

increasing temperature, an oxycarbide melt due to dissolution of SiC.13 The use of

A1203-Y203 system was because this compound forms a eutectic melt at a lower

temperature.5'31-33 The resulting microstructure from this additive system usually yielded

a yttrium-aluminum-garet (YAG)-containing grain boundary phase, which was believed

to enhance the toughness values of the SiC ceramics. A toughness value of about 8

MPa*ml/2 was reported with this additive system.5 The difficulty, however, of using

either Al203-Y203 or Al203 additives is controlling the partial reduction of Al203 at the









sintering temperatures usually used. A1203 is reduced at the sintering temperatures near

20000C via the reaction,

A1203 + SiC Al20(g) + SiO(g) + CO(g) (1)

Below 20000C, yttria alone does not show appreciable gas phase reaction with SiC,

nor does it form sufficient melt phase to allow complete densification.23 Thus, the use of

SiO2-rich powder bed to prevent the escape of reaction products during sintering has been

employed. This complicated process is not desirable for an industrial production and

usually results in low mechanical property-reproduction in resulting monoliths. A better

alternative was the use of AlN-Y203 additive, where the presence of aluminum was

maintained.23 The decomposition reaction of AIN, given by

A1N -2Al() + N2(g) (2)

can be effectively suppressed by a nitrogen overpressure during sintering. An added

advantage, realized in the later studies,7'36'40 was the formation of solid-solution between

A1N and SiC which produces a wurtzite 2H-SiC structure that effectively impinges on the

anisotropically growing hexagonal (6H) grains during sintering.8 The 2H-SiC structure

does not form a solid solution with 6H-SiC polytype.8 This phenomenon offers a way of

controlling the observed anisotropic grain growth and coarsening in SiC ceramics which

results in strength reduction of the manufactured ceramics.

Addition of CaO in A1203-Y203-containing SiC ceramics has been used to lower

the vapor pressure of product gases and to further lower the sintering temperature.3'50'59

Addition of CaO further decreases the grain growth rate and reduces the mass loss during

processing. The use of rare earth and lanthanum series oxides was driven by the

similarity of chemical and physical properties of these oxides with Y203.42 The









differences in the cationic field strength (z/r2; valence (z), bond length (r)), however,

resulted in the differences in properties of the grain boundary phases formed from the

different oxides. It has been reported that a decrease in the cationic radius of the rare-

earth oxides was accompanied by an increase in Young's modulus, hardness, and flexural

strength of the SiC ceramics, whereas the fracture toughness was improved by

incorporating rare-earth oxides of larger cationic radius.42 In the case of lanthanum series

oxides, specifically Lu203, formation of highly refractory crystalline rare earth disilicates

in the grain boundary and intergranular phase leads to an improvement in the high

temperature properties of the SiC ceramics produced.25

Effects of Processing Parameters in LPS-SiC

For a SiC ceramic to be useful in its desired application, be it as a high-temperature

structural material or as a ceramic armor component, it should be able to exhibit

relatively high hardness and toughness. Like traditional ceramics, SiC exhibits a

hardness/toughness trade-off. Thus, judicious selection of fabrication parameters for this

material should be made in order to have mechanical properties suitable for its intended

applications. In SiC, microstructure and grain boundary and triple junction phases

control the mechanical properties. Knowledge of the fabrication parameters that controls

these variables is therefore important.

Effect of Sintering Time and Temperature

The range of sintering temperature, as reported in the literature, varies from 1750C

to 22000C.30-58 The choice of sintering temperature generally depends on the melting

temperature of the additives used, in the case of single element additive, and the

temperatures of the eutectic melts, in multi-element additive system. For liquid-phase









sintered SiC, the need for enough liquid-phase volume to allow adequate wetting of the

solid during densification becomes the primary consideration in sintering temperature

selection.13 As an example, SiC sintered with A1203 and rare-earth oxides is usually

sintered at 18000C. 24 This temperature is higher than the range of eutectic temperatures,

17200 to 17800C, of these additives.

Another factor that determines the range of temperature used in sintering of SiC is

the formation of a desired phase useful in the toughening of SiC ceramics, although phase

formation or transformation also depends on the type of additive used. The P-to-a-SiC

phase transformation responsible for most of the toughening observed in SiC ceramics

occurs at 19500C.2 The temperature range for the stability of P-SiC, however, can be

extended through the use of appropriate additive, i.e., A1N or oxynitride-phase forming

additives.7'23 Formation of 2H-SiC from P-SiC with A1N additive that controls the

anisotropic grain growth of 6H-SiC polytype, depending on the amount of A1N additive,

occurs at 16000C. Thus, sintering temperature choices are usually higher than the above-

mentioned temperatures.

The effect of sintering temperature on the microstructure of the resulting ceramics

is manifested in the grain growth and morphology. Higher temperature generally leads to

bigger grains and, for 3-SiC starting powders, an elongated grain morphology. It has to

be noted, however, that sintering time also plays an important role in determining the

resulting microstructure and acts in tandem with the choice of temperature. Sintering

times reported for densification of SiC ceramics vary from 30 minutes to a couple of

hours. At a given temperature, sintering time controls the grain growth and the









completion of the P-to-a SiC phase transformations. Shorter sintering times generally

correspond to less grain growth and extent of phase transformation.

In addition to their effect on microstructure and grain morphology, sintering time

and temperature also affect the density of the sintered ceramics. In the time and

temperature range reported in the literature,30-s8 almost all of the SiC materials were

sintered to above 95% of the theoretical density. Near theoretical densities, however, are

usually attained at relatively higher temperatures and longer sintering times. Attainment

of fully densified SiC ceramics, however, also depends on how the sintering process was

carried out. For A1203-Y203 sintered SiC ceramics, use of powder beds with similar

composition to that of the materials being sintered is necessary to prevent the evolution of

reaction products during sintering that contributes to porosity formation in the final

material.8'23 For AIN-Y203 sintered SiC, an overpressure of nitrogen is necessary to

prevent the evolution of nitrogen gas during processing.8'23

Effect of Choice of Additives

Literature data indicate that the choice of additive used in sintering of SiC generally

controls the secondary phases (grain boundary, triple junction and residual) retained in

the material upon cooling. The secondary phases in the sintered ceramics affect the

toughening mechanisms and mode of failure operable in this material. It is desired, in the

case of ceramic armor application, that the secondary phase provides a weak interface

between the matrix grains to facilitate the toughening mechanisms operable. It is also

accepted that the weak interface due to the secondary phases ensures an intergranular

failure mode in the SiC ceramics. 18

As mentioned in the earlier section of this chapter, the additive system used in

sintering SiC includes A1203, A1203-Y203, AIN-Y203, A1203-Y203-CaO, A1203-rare-









earth oxides and Al203-lanthanum series oxides. Al, Al203 and A1N, in conjunction with

the SiO2 present in the starting SiC powders, generally form Al and O enriched secondary

phases in the sintered materials.5'38'60 The Al and O-enriched grain boundary and triple

junction phases provide a sufficiently weakened interface for intergranular mode of

failure to predominate in SiC sintered with these additives. In the case of A1203-Y203

sintered SiC, the resulting grain boundary and triple junction phases generally contains

yttrium-aluminum-garnet (YAG).5'59 In most cases, the grain boundary and triple

junction phase formed are generally weaker than the SiC grains, thus providing an easier

path for crack formation leading to intergranular mode of failure in the material.

Aside from the formation of secondary phases, several additives also affect the

resulting microstructure of the sintered SiC. In the case of SiC ceramics fabricated from

P-SiC starting powders, some of the additive used also affect the P-to-a phase

transformation responsible for the in-situ toughening in this material.2'53-58 As shown by

studies performed by the University of California, Berkeley group,2,53-58 the use of Al, B

and C has corroborating effects that promote and enhance the P-to-a SiC phase

transformation. They have observed that in terms of developing phase composition,

boron is more effective in promoting the P-to-a phase transformation than carbon.

Aluminum, on the other hand, retards the P-to-a phase transformation, but promotes 6H-

to-4H transformation. In terms of grain morphology, aluminum and carbon promote

anisotropic grain growth resulting in a plate-like or elongated microstructure in the final

ceramics. Boron, on the other hand, tends to coarsen the grains but reduces the average

aspect ratio.









Studies of the ABC-SiC system have also revealed that the combined roles often

override the individual roles of Al, B and C during processing of this material.2'53-58

Although both boron and carbon together favor the P-to-a phase transformation

associated with grain elongation, the resulting microstructure does not necessarily lead to

strongly elongated grains as boron and carbon additives have the opposite effect on

anisotropic grain growth. Thus, the final grain configuration in the resulting ceramics is

determined by the B/C ratio. The effect of aluminum, however, is such that if the B/C

ratio favors the anisotropic grain growth, Al accelerates such growth so that the aspect

ratio is further increased. It has also been observed that reduction of the Al/B and Al/C

ratio results in a weakened aluminum effect on the final microstructure due to the

decrease in liquid phases generated. Even at constant Al/B/C ratios, the resulting grain

configurations are significantly altered by a change in the total amount of additives.

Thus, in the case of ABC-SiC system, as might be true in other sintering additive

systems, literature data imply that both the ratios and total amount of sintering additives

affect the final grain morphology and microstructure of the sintered SiC ceramics.

Another sintering additive that affects not only the retained secondary phase but the

resulting microstructure is AIN.23'35-37'61-64 It has been shown that for P-SiC starting

powders, use of AIN as sintering additive results in the formation of 2H-SiC polytype,

together with the common 4H and 6H SiC structure, in the densified ceramics.3537 The

formation of 2H structure results in equiaxed grain morphology, while 6H and 4H yield

elongated grains. At appropriate temperature and amount of additive, the 3C-to-2H SiC

transformation can be made to go into completion such that the resulting grains are all

equiaxed. If completion of the 3C-to-2H transformation is not achieved, the resulting









microstructure contains elongated and equiaxed grains. Thus, different microstructure

can be achieved in this system depending on the amount of AIN and the processing

temperature.

Effect of the Starting SiC Powders

SiC crystallizes in a zinc-blend cubic structure and a number of different polytypes

with a hexagonal or rhombohedral symmetry. The cubic structure is known as the P-SiC

phase and is given the Ramsdell notation of 3C.65 The hexagonal and rhombohedral

polytypes are collectively called the a-SiC phase and the most common, in Ramsdell

notation, are 2H, 4H, 6H and 15R.65 P-SiC is a metastable phase and transform at higher

temperature to one of the a-SiC polytypes, depending on the type of additive used. P-SiC

transform to a 6H a-SiC polytype in the presence of boron, while the 4H modification is

dominant in the presence of aluminum.2 Use of aluminum and nitrogen, either in the

form of A1N or Al with N2 as sintering atmosphere, generally stabilizes the 2H

polytypes.3-37 The rate of transformation, however, in most cases is relatively slow due

to the small difference in free energy (about 2 kJ/mol) between the a- and P-SiC phases.23

The polytype of the starting powder is not important for the densification of the SiC

ceramics. It has been reported that near theoretical densities can be achieved with the

appropriate choice of sintering additives, sintering time and temperatures.5 It has also

been shown that the P-to-a phase transformation does not contribute to the densification

of the SiC ceramics.4 The starting polytype, however, severely affects the resulting

microstructure. SiC sintered from a-SiC starting powders usually shows equiaxed grain

morphology.4 Although grain growth usually occurs, the aspect ratio of the larger grains

remains near unity. This is particularly true for starting a-SiC powders of unimodal grain

size distribution. In cases where the grain size of the starting powder is intentionally









varied, i.e., the grain size exhibits a bimodal distribution; the resulting microstructure

usually shows a bimodal grain distribution distinctly different from the starting powder.

In an experiments where a-SiC seeded with large a-SiC grains are used, the resulting

microstructure shows a bimodal grain size distribution consisting of large elongated

grains in small equiaxed grain matrix.45 This was attributed to the grain growth

mechanism in liquid phase sintering, i.e., solution-reprecipitation. The large grains grow

at the expense of the smaller grains, thus, resulting in a microstructure with large

elongated grains in a matrix of small equiaxed grains.

Microstructure with small equiaxed grains can also be fabricated from a P-SiC

starting powder. This is possible if the P-SiC starting powder is sintered in a temperature

and time range where the P-to-a phase transformation was prevented.2 SiC sintered from

P-SiC starting powders processed in an overpressure of nitrogen, with A1N or AIN-Y203

additive also shows the same equiaxed grain morphology when processed at temperatures

below 20000C.8'14 This is believed to be due to the effect of the oxynitride phase that

forms during processing which extends the stability regime of the P-SiC phase. If the P-

to-a phase transformation was allowed, the resulting microstructure usually contains

plate-like or elongated grain morphology.8 If the starting P-SiC powders are seeded with

similarly sized a- or P-SiC, the resulting microstructure still shows a unimodal grain size

distribution.45 With a different grain size seeds, the resulting microstructure shows a

bimodal grain size distribution consisting of large elongated grains in a matrix of small

elongated grains.45 This is especially true for P-SiC starting powders of grain size in

nanometer range seeded with up to 5 wt. % of micron-size grains.51









Effect of Sintering Atmosphere

The sintering atmosphere can have several important effects on densification and

microstructure development during sintering. In many instances, the atmosphere can

have a decisive effect on the ability to reach a high density with controlled grain size.

The important effects of sintering atmosphere are associated with the gas solubility and

the chemical reaction with the powder system. Sintering atmosphere reactions with

powder system is especially true for SiC since volatility of reaction products in this

material, depending on the additive, is an issue during sintering.

In sintering SiC ceramics, densification to near theoretical density are readily

achieved with the use of argon and nitrogen atmosphere.12 Nitrogen atmosphere, in most

cases, is used to address the issue of escape of gaseous reaction products during

processing.30 For SiC sintered with A1203-Y203-SiO2 additive (SiO2 is always present on

the surface of SiC powders), the reaction of A1203 and SiO2 with SiC at high

temperatures generate gaseous reaction products.31 As reported in the literature,23 the use

of powder beds of composition similar to the powders being sintered is generally required

to attain a high density ceramics. A change in additive, i.e., AIN-Y203 system, and an

overpressure of N2 gas provides a way to control the gaseous reaction products resulting

in a high density final ceramics.23'30 In some cases, the use of nitrogen as sintering

atmosphere results in the absorption of nitrogen in the grain boundary and triple junction

phases. The oxynitride phase is believed to increase the stability of P-SiC such that the P-

to-a phase transformation can be prevented.8

The effects of argon and nitrogen as sintering atmosphere in the resulting

microstructures of SiC ceramics are patent. For a-SiC starting powders of uniform size

distribution, it has been observed that sintering in argon atmosphere produces equiaxed









grains in the microstructure.11'12 If the starting powder, however, is seeded with large a-

SiC grains, the resulting microstructure shows a bimodal distribution where large

elongated grains exists in a matrix of small elongated grain." For a P-SiC starting

powders sintered in argon atmosphere, the resulting microstructure usually shows

presence of elongated grains.4 This is due to the P-to-a phase transformation that easily

occurs in argon atmosphere. If the P-to-a phase transformation is made to occur to

completion, the resulting microstructure has a unimodal distribution of elongated grains.

If the phase transformation is incomplete, the resulting microstructure usually has a

bimodal grain distribution wherein the elongated a-SiC grains is contained in the matrix

of equiaxed P-SiC grains.

In a nitrogen sintering atmosphere, the resulting microstructure using a-SiC starting

powders is similar to that observed in materials processed in argon atmosphere." The

grain morphology is usually equiaxed and grain growth is minimal. In the case of P-SiC

starting powders, the resulting microstructure is different than those observed in argon.12

The use of nitrogen is believed to extend the stability of P-SiC such that the P-to-a phase

transformation is shifted to a higher temperature.14 Thus, at low temperature, below

20000C, the resulting microstructure usually shows equiaxed grain morphology. At high

temperature, above 20000C, the resulting microstructure generally shows presence of

elongated grains. If the additives used, however, can form a solid-solution with SiC,

example of which is A1N, the resulting microstructure varies.3537 If the resulting phase in

the microstructure is a combination of 6H-, 4H- and 2H-SiC structure, the resulting

morphology shows a bimodal grain size distribution. 6H- and 4H-SiC generally appear

elongated and 2H-SiC shows equiaxed grain morphology. If the phase transformation is









driven to form the 2H-SiC structure, the resulting microstructure generally shows

equiaxed grains.

Effect of Annealing Conditions

Annealing in the fabrication of SiC is an additional step in the processing wherein

further grain growth and phase transformation can be achieved. It has been demonstrated

that appropriate choice of annealing conditions can greatly enhance controllability of the

resulting microstructure.5 Lee et af fabricated SiC ceramics from a-SiC powders

sintered with SiO2 and YAG-forming additive. Upon hot pressing at 18500C for one

hour under a 25 MPa load in argon atmosphere, the resulting microstructure of the

densified material showed relatively fine equiaxed grains. Annealing at 19500C in

flowing argon for 4 hours resulted in a self-reinforced microstructure consisting of large

elongated grains and relatively fine elongated grains.

Introduction of external pressure during annealing affect the microstructure of the

sintered ceramics. Kim et a6 hot pressed SiC ceramics at 17500C for 40 minutes under

25 MPa in argon atmosphere and further annealed the resulting monolith at 18500 to

1950C with or without applied pressure. The hot pressed material showed fine,

equiaxed grains since the hot pressing temperature is low enough for the P-to-a phase

transformation to occur. Upon annealing for 4 hours, without applied pressure, the

resulting microstructure showed elongated grains due to the P-to-a phase transformation

that occurred. The densities measured are lower than those materials annealed with

pressure. The average grain diameter and aspect ratio were observed to increase with

annealing temperature, however, the grain growth comes mainly from an increase in

aspect ratio. For materials annealed with applied pressure at temperatures below 1950C,









the resulting microstructure showed equiaxed grains with relatively low aspect ratios. At

1950C, elongation of some grains occurred resulting in a duplex microstructure

consisting of large elongated grains and small equiaxed grains. The aspect ratios of the

elongated grains, however, are lower than those observed in materials annealed without

pressure. These results have shown that annealing under an applied pressure generally

inhibits grain growth and the P-to-a phase transformation, which can provide a new

strategy for control and optimization of mechanical properties in SiC ceramics.

Microstructure and Mechanical Property Modification in Sintered SiC

It is a known fact that optimization of mechanical properties of materials can be

achieved via the control of microstructure in the final product. The mechanical properties

desired depend on particular applications that the material is being considered. In the

case of SiC, this material has been used, and considered, in several structural applications

because of its superior properties in terms of wear, corrosion, high-temperature creep,

and oxidation resistance, as well as its high temperature retention characteristics. The

drawbacks that make SiC inapplicable in some of its envisioned structural applications,

as in most polycrystalline ceramics, are the low fracture toughness and extremely flaw-

sensitive strength. The latter is a manufacturing/processing issue and can generally be

improved by a post-processing treatment while the former can be addressed via the

control of the microstructure of the final product. Discussion on what has been done to

address the fracture toughness issue of SiC will be presented here. The post-processing

treatment that has been used to address the issue of flaw-sensitive strength will be

presented in latter section.









Fracture Toughness Enhancement in SiC

Cao et al2 have used P-SiC starting powders and took advantage of the P-to-a phase

transformation that produces platelike microstructure in the resulting monoliths in

producing a high toughness SiC. The P-to-a phase transformation were further enhanced

and promoted through the addition of aluminum, boron and carbon additives. The

resulting microstructure was controlled through the variations in sintering time and

temperature parameter during processing. The measured fracture toughness of the

resulting ceramics reached a value of 9.5 MPa*m1l2, the highest value reported so far.

The microstructure of the ceramics exhibiting this fracture toughness value showed

mostly elongated grains. Cao et al have also shown that grain growth observed

correspond to an increase in the fracture toughness and bend strength. The fracture

toughness of the SiC ceramics has also shown a strong dependence on the aspect ratio of

the elongated grains. The bend strength, on the other hand, shows no sensitivity towards

the aspect ratio.

The mechanisms responsible for toughening in ABC-SiC were reported to be due

to crack bridging and subsequent grain pullout, with minor contribution from crack

deflection. Cao et al observed partially debonded platelike grains, which resided just

behind the crack tip, bridging the crack, thereby reducing the effective stress intensity.

Based on SEM of crack path propagation on ABC-SiC sample, crack deflection due to

the presence of elongated grains was also apparent. These mechanisms were not

observed to be operable in materials hot pressed at lower temperature, 17000C, where the

microstructure of the resulting ceramics generally shows equiaxed grains. Thus, the

toughness measured from these materials is generally lower. For materials sintered at









higher temperatures, 19000 to 19500C, resulting microstructure shows elongated grains

with higher aspect ratio. This kind of microstructure was believed to enhance the effect

of the toughening mechanisms operable resulting in a significant improvement in the

fracture toughness measured.

In another paper, Moberlychan et a!54 compared the microstructure and mechanical

properties of several SiC materials. They have found that higher fracture toughness was

associated with intergranular mode of fracture and high aspect ratio of elongated grains in

the SiC. It was found that the intergranular mode of fracture is enhanced by the presence

of "weaker" amorphous secondary phases in the grain boundary and triple junctions of

the studied material. Thus, Moberlychan et al concluded based on the comparison made

that high toughness in SiC ceramic requires an amorphous grain boundary layer with a

chemistry to weaken the grain boundary and a grain shape with a high aspect ratio. Lack

of any of these ingredients produced inherently brittle materials.

The additives used in sintering SiC ceramics affect the resulting microstructure

which in turn can affect the fracture mode. As shown in ABC-SiC system reported by

Zhang et al,53 a change in the amount of aluminum can result in a different

microstructure. For aluminum addition of 3 to 7 wt. %, the resulting microstructure

shows equiaxed and elongated SiC grains. The elongated SiC grains show an increasing

aspect ratio as the amount of aluminum was increased. The corresponding volume of the

elongated grains in the material, however, decreased as the aluminum content was

increased such that at higher aluminum content (6-7 wt.%), the resulting microstructure

has a much higher volume of fine equiaxed grains. Because of this change in the

microstructure, a noticeable change in the fracture mode between 5 and 6 wt. %









aluminum was observed. For Al additive less than 5 wt. %, the observed fracture mode

was primarily intergranular, while starting at 6 wt. %, the fracture mode observed are

predominantly transgranular. Changes in the microstructure and fracture mode were also

observed for ABC-SiC at lower Al content. Flinders et al8 found, for a similarly

processed P-SiC powders, that a transgranular to intergranular mode of fracture occurs

between 1 and 1.5 wt. % aluminum additions. The corresponding microstructure changes

from a primarily equiaxed grains to a primarily elongated ones with a reportedly lower

aspect ratio than reported by Zhang et al.

Other attempts in improving the fracture toughness of SiC ceramics processed from

P-SiC starting have been reported in the literature.4'7'9 The achieved fracture toughness

values, however, are all lower than those measured from ABC-SiC system but still

significantly higher than those observed for solid-state sintered SiC. Hilmas and Tien35

measured a fracture toughness of 8.5 MPa*ml/2 through the use of 4.95 wt.% A1203 5.97

wt.%AIN -0.5 wt.% B additives hot pressed at 21000C for one hour. The resulting

microstructure showed mostly elongated grains, due to the 3C-to-6H transformation, and

minor fine equiaxed grains, due to the 3C-to-2H transformation. The mode of fracture

observed was intergranular. Aside from crack bridging and crack deflection mechanism

of toughening, interfacial microcracking due to substantial residual stresses at the grain

boundaries which resulted from thermal expansion anisotropies have been considered

operable in this particular material. This was further supported by the measured hardness

value which is significantly lower for SiC ceramics. The low hardness value was

attributed by the authors to the possible microcracking that occurred during cooling of the

ceramics.









Improvement in the fracture toughness of SiC ceramics were also observed by Lee

et a5l who reported a fracture toughness value of 8.5 MPa*ml2. In their case, the

sintering additive used for the densification of P-SiC was 5.7 wt.% A1203 3.3 wt.%

Y203 1 wt.% CaO. The reported toughness value was from a ceramic hot pressed at

18500C for 1 hour under 25 MPa in argon atmosphere. The hot pressed ceramic was

further annealed for 4 hours at 19500C under flowing argon. The resulting

microstructure was reported to contain mostly of elongated grains with grain size and

aspect ratio of 1.25 tm and 3.9, respectively. The mode of fracture in the reported

material is predominantly intergranular.

Hot pressing of P-SiC using A1203 rare earth oxides also proved to be useful in

improving the fracture toughness of sintered SiC, as shown by Zhou et al.10 Although the

reported toughness values are relatively lower than those reported for P-SiC starting

powder, the reported values are only from the hot-pressed specimens. The microstructure

of these materials generally showed equiaxed grains and XRD analyses revealed that

most of the grains are untransformed P-SiC. If the P-to-a transformation were made to

occur, the resulting microstructure will be different and will possibly improve the fracture

toughness, as was shown by other researchers.

The use of similar additive system on the densification of P-SiC powder seeded

with large a-SiC at different additive amount and processing temperature resulted in

distinct mechanical properties. Zhan et al have used 7 wt.% A1203 2 wt.% Y203 1

wt.% CaO as sintering additive on a 90 nm sized P-SiC powder.4 The resulting fracture

toughness values are relatively lower than those reported by Lee et al.5 The hardness

values, however, are relatively higher, even for the as hot-pressed material. The









microstructure for the as hot-pressed material generally showed equiaxed grains. Further

annealing, resulted in a microstructure with uniformly elongated grains.

Improvement in the fracture toughness of SiC can also be achieved through the use

of a-SiC starting powders. Although p-to-a phase transformation resulting in elongated

grains will not occur when using a-SiC, elongated microstructure can also be produced

through the anisotropic grain growth. Kim et al were able to achieved a fracture

toughness ranging from 5.4 to 6.2 MPa'ml/2 using a-SiC starting powder sintered with

YAG and YAG-Si02 additives.6 The resulting microstructure of the hot pressed

materials generally showed equiaxed grains with aspect ratio near unity. Upon annealing

at 1950C for 4 hours under a flowing argon atmosphere, the resulting microstructure

showed a "self-reinforced" morphology consisting of large elongated grains in a matrix

of small elongated grains.

Hot pressing of a-SiC using different additives, as reported by Flinders et al and

Nader et al,7 also resulted in a possible improvement in the mechanical properties.

Flinders et al hot pressed a-SiC using Al, A1N, and A1203 additive and showed minor

improvement in the fracture toughness.26 The values reported, however, were from as-

hot pressed specimen and possible improvement can still be achieved through annealing.

Nader et al, on the other hand, used Y203 A1N and showed that pressureless sintering of

a-SiC results in moderate toughness values, even at longer sintering times (14 hours).

This is in contrast to their result using 3-SiC starting powders where marked

improvement in the fracture toughness of the SiC ceramics, even with the use of

pressureless sintering, was achieved.









To summarize, the microstructural toughening of SiC can be achieved through the

use of P-to-a phase transformation, the use of liquid phase forming additives, seeding or

controlling the initial a-SiC, and annealing for controlled grain growth. The P-to-a phase

transformation results in a development of in-situ or self-reinforced microstructure which

enhances the toughening mechanisms operable. The use of liquid phase forming additive

and seeding generally enhances the P-to-a phase transformation resulting in a toughened

ceramics. The low amount of additive used in some of the studies reported generally

results in an improvement in the resulting toughness and hardness values. The use of

annealing, even in ceramics sintered from a-SiC powders, results in a pseudo-self

reinforced microstructure which also shows promise in improving the toughness of the

material.

Crack Healing as a Tool for Mechanical Property Improvement in SiC

The phenomenon of crack healing has been observed in a variety of different

ceramic materials including single crystal,66 polycrystalline ceramics,67-7 inorganic

glasses,72 and ceramic composites.73-92 It has also been observed in metals93 and polymer

materials.94 This phenomenon is usually described as a process where the surface crack

of a given material closes, either by crack-face rebounding or by filling up of the crack

face, after a given heat treatment. Associated with the crack closure is an observed

increase in the flexural strength of the material.

The recent increase in the number of studies, in the late 1990s,73-92 on crack-

healing is driven by the promise of recovering the mechanical properties of structural

materials that are compromised due to inherent flaws introduced during processing and

handling. Application of a crack healing treatment as a post-production process will

reduce, if not entirely eliminate, the flaws/cracks introduced during processing and









handling. Crack healing will greatly enhance the reliability and integrity of the structural

material in question. It will also decrease the machining, maintenance, and inspection

costs and increase the lifetime of the ceramic material.

Most of the crack healing studies in the literature were done in an oxidizing

environment.73-92 In most cases, complete recovery of the flexural strength of the

indented sample can be achieved within an hour of heat treatment. The mechanism for

the crack healing is usually attributed to the formation of the oxidation products in the

crack surface, bonding the crack walls and rebounding the crack faces. The increase in the

flexural strength observed is usually attributed, aside from the reduction in flaw sizes, to

the stress relaxation upon heat treatment and possible crack tip blunting.

On the belief that crack healing by heat treatment is a phenomenon which occurs

for all material that sinter, F.F. Lange, K.C. Radford and T. K. Gupta showed that crack

healing was possible for ZnO, SiC and polycrystalline A120366'68-69'95 In their

experiments, deep surface cracks were introduced to the specimens by quenching into

water at room temperature. The thermally shocked specimens were then heat treated in

air at a temperature just below the sintering temperature of the starting SiC, ZnO, and

A1203 powder. For the case of SiC and A1203, the specimens were heat treated for one

hour at 14000C and 17000C, respectively. In all cases, the flexural strength measured

after heat treatment showed strength values higher than the un-oxidized thermally

shocked specimens suggesting that crack healing occurred. For longer heat treatment

times, strengths higher than that of the as-received specimen were observed. Microscopic

examination of the crack patterns after heat treatment also revealed that the crack patterns









disappeared, and in the case of SiC, formation of an oxide layer on the crack surface was

evident.

In the study of the crack healing behavior of SiC, F. F. Lange used two types of

SiC, a low density (80 % of theoretical) and a high density form.69 For the low-density

SiC, flexural strength measured after oxidation yielded a higher value than that of the un-

oxidized quenched specimens. Flexural strengths measured after 90 hours of exposure

were even higher than the flexural strength of the as-cut specimens. Results for the high

density material, on the other hand, showed a minimal (10% of the unoxidized quenched

specimen) flexural strength increase for the same oxidation time. Rationalization for the

difference in strength recovery was provided by the difference in the oxidation behavior

of the material, i.e., the high density form has experienced very little oxidation. These

observations, coupled with the observed oxidation of the cracked surface, suggest that the

formation of the oxide film was responsible for crack healing.

Crack-healing behavior also depends on how the ceramic material was

manufactured. This was shown by the studies performed on a reaction-bonded and liquid

phase-sintered silicon carbide.95s96 Both of these materials showed crack healing ability,

however, the strengthening achieved for liquid phase-sintered material was higher. In

both materials, the crack healing ability was attributed to the flow of the preexisting glass

phase across the crack plane, and the reaction of the crack surface with the environment

(oxidation). The difference in strengthening was attributed to the different residual stress

produced by the thermal expansion mismatch between the oxidation product and the bulk

material. For the reaction bonded silicon carbide, the residual stress is lower since above









5000C, silicon becomes ductile making dislocation movement possible. Thus, the

residual stress contribution for this material will be due to cooling from this temperature.

Systematic studies of the crack healing behavior of several ceramic and ceramic

composites geared towards improving the reliability and integrity of these materials were

only accomplished in the late 1990s by several Japanese investigators.73-92 They

performed crack-healing behavior evaluations on the following materials; silicon carbide,

silicon carbide-reinforced silicon nitride composite, mullite/silicon carbide ceramics and

silicon carbide whiskers toughened alumina. Several issues including: a) the effect of

chemical composition on the crack healing ability, b) the effect of healing conditions on

the strength of the healed-zone, c) determination of the maximum crack size which can

be healed, d) knowledge of the high-temperature strength of the healed zones, e)

understanding of the crack healing mechanisms and f) assessment of the cyclic and static

fatigue strengths of crack-healed ceramic member, were also addressed in their

investigations.

Results of the investigation of the Japanese group are summarized here. In most of

the materials that they studied, they found out that surface crack sizes of up to 200 tam

can be healed, if the crack-healing was done under the optimal healing conditions. Most

of these crack-healing conditions are in the temperature range of 12000 to 15000C, in air,

for a heat treatment time of 1 hour. In the best healing conditions, the cracked-healed

materials flexural strength is either higher or comparable to the flexural strength of the

smooth samples. Failure during the flexural strength testing occurred mostly outside the

Vickers indentation zone, suggesting that the healed zone has sufficient strength.

Variations of the flexural strength of the crack-healed specimens are generally the same









as that of the smooth specimens. It was also shown that most of the crack-healed

specimens are not sensitive to static fatigue testing below an applied stress of 70-75% of

the average monotonic bending strength. Cyclic fatigue testing of Si3N4/SiC (R = 0.2,

frequency = 8 Hz, N = 2 x 106 cycles) also showed that the maximum stress at which the

crack-healed specimens did not fracture is usually about 3 3.5 times that of the crack

specimens, and higher than those for smooth specimens. It was also shown that most of

the materials tested are capable of crack-healing even in the presence of cyclic stress.

These results suggest that an improvement in structural reliability can be achieved via

crack-healing, with the primary mechanisms being the formation of oxidation products

which bonds and strengthen the crack surface. The observed flexural strength increase,

on the other hand, was generally attributed to stress relaxation and crack tip blunting.














CHAPTER 3
STATEMENT OF THE OBJECTIVE

The general objective of this investigation is to verify the correlation between grain

boundary and triple junction chemistry and mechanical properties of several SiC

ceramics sintered with minimum additives for armor applications. In order to accomplish

this objective, the chemistry of the triple junction and grain boundary phase will be

determined through the combination of the following techniques: High resolution

transmission electron microscopy to determine the presence of amorphous intergranular

films, energy dispersive spectroscopy and energy filtered transmission electron

microscopy to determine the composition of the grain boundary/intergranular films and

triple junction phase, and electron energy loss spectroscopy to determine the exact nature

of the triple junction phases.

Several sets of SiC ceramics sintered with the addition of aluminum, either in the

form of aluminum metal or aluminum compounds, will be investigated. SiC ceramics

fabricated from P-SiC starting powders with fixed boron and carbon content and varying

amounts of aluminum addition will be studied to determine the effect of amount of

aluminum on the grain boundary and triple junction characteristics. Mechanical

properties of these materials, particularly toughness, hardness, and mode of fracture, will

also be determined. Correlation between the mechanical properties and the grain

boundary and triple junction characteristics will also be explored.

SiC ceramics fabricated from a-SiC starting powders with equimolar aluminum

additions (from Al, A1N and A1203) will be investigated to determine the effect of









sources of aluminum on the grain boundary and triple junction chemistry. Effect of

boron and carbon addition on the microstructure and grain boundary and triple junction

chemistry will also be explored.

SiC fabricated from a-SiC starting powders with aluminum or aluminum nitride

addition in differing amounts and hot pressed at different temperature will also be

studied. This is to determine the effect of amount of additive and processing temperature

on the resulting microstructure and mechanical properties. Grain boundary and triple

junction characteristics will also be investigated in these materials to determine the

effects of aforementioned variables. The implications of the observations on these

materials on the microstructure and property tailoring in SiC ceramics will also be

discussed. Comparison with a commercially available state-of-the-art ceramic armor will

also be made. Strength improvement through crack healing in SiC ceramics will also be

demonstrated.

Although many investigations have been carried out on the microstructure and

property tailoring of SiC, data on grain boundary and triple junction phase characteristics

of this material are very limited. The investigations performed in this dissertation aim to

contribute to the literature in this area. SiC ceramics sintered with minimum amount of

aluminum (Al, A1N and Al203) are used to demonstrate that high toughness and hardness

are achievable at low level of additives.














CHAPTER 4
MATERIALS AND METHODS

The materials processing and characterization methods used in this investigation

are presented in this chapter. The materials used in this study were all fabricated by

Ceramatec, Inc. The mechanical properties (hardness, toughness and mode of fracture),

chemical analysis and SiC polytypes characterization were also performed by Ceramatec,

Inc. Electron microscopy such as energy dispersive spectroscopy (EDS) and high

resolution transmission electron microscopy (HRTEM) were done at MAIC at the

University of Florida. Energy filtered transmission electron microscopy was performed

at MCF at the University of Central Florida.

Materials

The materials used in this investigation are presented in Table 4-1. The materials

were grouped in five classifications according to the type of starting SiC powders, kind of

additive and intended experiments. ABC-SiC system was investigated to determine the

effect of amount of Al additive on the grain boundary and triple junction. It was also

intended to find any correlation between the grain boundary and triple junction

characteristics, and the mechanical properties and mode of fracture in these materials.

Materials fabricated from a-SiC were used to investigate the effect of different sources of

aluminum additive, hot pressing temperature and amount of additive and find an

alternative way to tailor properties similar to commercially available ceramic armor

material (SiC-N). Additions of B (in form of B4C) and C were also performed to









investigate effect of these additives on the mechanical properties and grain boundary and

triple junction characteristics.

Table 4-1. Compositions and processing conditions of the materials investigated.
Compositions (wt. %) Hot Pressing Conditions
Temperature Holding Time Atmosphere/Pressure
ABC-SiC System
P-SiC-0.6B-2C 21000C 1 hr Argon/28 MPa
P-SiC-0.5A1-0.6B-2C 21000C 1 hr Argon/28 MPa
P-SiC-1A1-0.6B-2C 21000C 1 hr Argon/28 MPa
P-SiC-1.5A1-0.6B-2C 21000C 1 hr Argon/28 MPa
P-SiC-4A1-0.6B-2C 21000C 1 hr Argon/28 MPa
a-SiC i ilth Al
a-SiC-1.65Al 19000C 1 hr Argon/28 MPa
a-SiC-1.65Al 21000C 1 hr Argon/28 MPa
a-SiC-1.65Al 22000C 1 hr Argon/28 MPa
a-SiC-3.3A1 20000C 1 hr Argon/28 MPa
a-SiC-3.3A1 22000C 1 hr Argon/28 MPa
a-SiC-1.65A1-0.5B4C-2C 21000C 1 hr Argon/28 MPa
a-SiC i ith AIN
a-SiC-2.5A1N 21000C 1 hr Argon/28 MPa
a-SiC-2.5A1N 22000C 1 hr Argon/28 MPa
a-SiC-5A1N 20000C 1 hr Argon/28 MPa
a-SiC-5A1N 22000C 1 hr Argon/28 MPa
a-SiC-2.5A1N-0.5B4C 21000C 1 hr Argon/28 MPa
a-SiC-2.5AlN-0.5B4C-2C 21000C 1 hr Argon/28 MPa
Comparison Materials
SiC-N Proprietary Proprietary Proprietary
a-SiC-3.1A1203 21000C 1 hr Argon/28 MPa
a-SiC-0.5B4C-2C 21000C 1 hr Argon/28 MPa
Crack Healing
P-SiC-0.6B-2C 21000C 1 hr Argon/28 MPa
P-SiC-1.5A1-0.6B-2C 21000C 1 hr Argon/28 MPa
P-SiC-0.43A1-0.5Y203 22000C 1 hr Argon/28 MPa


The starting powders used for processing are listed in Table 4-2 along with the data

provided by the suppliers. Ceramics fabricated from different starting powders were

processed differently. In the case of materials fabricated from p-SiC starting powders, P-

SiC powders were mixed with Al, B and C additives and dispersed with polyamine









polyester polymer by adding 1 wt. % of the polymer, based on solids, to 400 grams of

reagent toluene. The aluminum was added in 0.5, 1, 1.5 and 4 wt. % whereas the born

and carbon were fixed at 0.6 and 2 wt. %, respectively. The carbon was added as 4 wt.%

apiezon wax which would result in 2 wt. % C incorporation upon pyrolysis. The slurries

were then deagglomerated for two hours with paint shaker and rolled overnight before

drying. The powders were sieve through a 44 [tm screen before hot pressing at 28 MPa

in stagnant argon inside a graphite die. The hot pressing temperature used is 21000C and

holding time is 1 hour.

Table 4-2. Raw materials used in powder processing.
Powder Supplier Grade Information from Supplier (wt.%)
p-SiC Superior HSC-059 SA=15-17 m2/g
Graphite
a-SiC H.C Starck UF-15 0=1.0%, C=29.6%, d50=0.5 am, SA=15m2/g
Al Valimet H-3 Fe=0.12%, Al=99.98%, d5o=5.0 am
A1N Tokuyama Soda F 0=0.78%, C=0.03%, SA=3.4 m2/g
A1203 Sasol North Am. SPA-0.5 Si=14 ppm, d50=0.47 am, SA=7.6 m2/g
Y203 Molycorp 5600
B H.C Starck S-432B
B4C H.C. Starck HS 0=1.5%, B/C=3.80, d50=0.96 am, SA=18.0
m2/g
C Apiezon W 50 wt.% yield after pyrolysis
C Capital REsin CRC-720 40 wt.% yield after pyrolysis


In the case of SiC ceramics fabricated from a-SiC starting powders, all

compositions were prepared by watching 600 grams of powder in two-liter high-density

polyethylene (HDPE)jars filled with 1.6 kg solid state sintered SiC media and 700 g

reagent grade acetone. The slurries were mixed for 16 hours on a ball mill in order to

disperse the agglomerates. The powders were stir dried before screening through an 80-

mesh screen. Compositions batched with B4C and C were pyrolyzed by heating in N2 to

6000C. Billets (45 mm x 45 mm x 6 mm) were hot pressed at 28 MPa inside graphite









dies by heating to 15000C in a vacuum of 1 torr, holding at 15000C for one hour to

remove SiO and CO from the samples, and then backfilling with Ar and heating to

21000C for one hour. The vacuum hold was performed to eliminate porosity in the

materials. SiC-N is a commercially available SiC from Cercom, Inc. produced via the

pressure assisted densification method.

Characterizations

Mechanical Properties and SiC Polytypes

Toughness and hardness measurements were performed on a test bar of dimension

3 mm x 4 mm x 45 mm. The bars were pre-cracked with a commercially-available

fixture (Maruto model MBK-603C) after using a single 98 N Knoop indenter to initiate

the crack. The tests were stopped after an audible pop-in noise was detected, with loads

ranging from 4 kN to 18 kN using bridging spans of 4, 5, or 6 mm. The original crack

was marked by a dye (black ink from an inkjet printer), using vacuum infiltration after

unloading the sample. The dye was oven dried at 500C overnight and cooled before

testing. Fracture toughness measurements were conservative since the crack location

used in the calculation was prior to the small amount of stable crack growth associated

with the SEPB tests. All crack planes were parallel to the hot pressing direction. Each

data point reported is the mean of 3-7 bars tested, with error bars representing + one

standard deviation.

All microhardness data (Leco Model LM-100) were obtained using a one kilogram

load Vickers (HV1) on polished surfaces. The dwell time at load was nominally 15

seconds. Rietveld analysis9 was used to determine SiC polytypes present in the









densified samples with X-ray diffraction patterns collected from a diffraction angle of 30-

800, with a step size of 0.020/step and a counting time of 4 sec/step.

Microstructure

In the case of SiC materials fabricated from P-SiC starting powders, the grain size

and aspect ratio reported are based on measurements from the polished and etched

surface of the sample. Polished samples of materials containing aluminum (0.5, 1, 1.5

and 4 wt. %) were plasma-etched by evacuating and back-filling with 400 millitorr of

CF4-10%02 and etched for 20 to 40 minutes. The solid state sintered material (no

aluminum) was etched in molten KOH at 5500C for 10 to 15 seconds. Grains size was

determined by line-intercept method, where the multiplication constant ranged from 1.5

(for equiaxed grains) to 2.0 (for elongated, platelike grains).99 The mean grain size were

measured from 200 to 300 grains for each composition and the aspect ratio was estimated

from the five most acicular grains.

In the case of SiC materials fabricated from a-SiC starting powders, the grains size

and aspect ratio reported were based on the measurements from a low magnification

TEM images taken from the sample. Average grain sizes were determined from 20

grains and the aspect ratio was measured from at least 5 elongated grains.

Chemical Analysis

Chemical analysis for nitrogen and oxygen constituents on selected starting

compositions and sintered samples were performed for materials sintered from a-SiC

starting powders. The analyses were performed based on the specification of ASTM

method E1409 for the determination of oxygen by the inert gas fusion/thermal










conductivity detection technique with a Leco TC-136 model analyzer. The detection

range for oxygen of this analyzer is from ppm level to below 5 wt. %.

Transmission Electron Microscopy

Transmission electron microscopy is unique among materials characterization

techniques in that it enables essentially simultaneous examination of microstructural

features through high-resolution imaging and the acquisition of chemical and

crystallographic information from small (usually in submicrometer range) regions of the

specimen. The signals generated, collected, and analyzed in TEM are produced by

interactions between the electrons in the high-energy incident beam and the material in

the thin film target. Some of these signals are illustrated in Figure 4-1. The direction

shown for each signal is not necessarily the physical direction but indicates where the

signals are generally detected.


Incident high-kV
beam
Backscattered
electrons (BSE) Se dary electrons (SE)
Characteristic X-rays
Auger electrons i
SIble Ight


'Absorbed Electron-hole
electrons' pairs



S/ Bremsstrahlung X-rays


Eacay Direct beam Inastically
scattered electrons electrons
scattered electrons


Figure 4-1. Signals generated from high-energy electron beam/thin specimen
interaction.100









Electron beam/specimen interactions, or scattering events, can be divided into two

categories: elastic and inelastic events. Elastic events affect the trajectories but do not

significantly affect the velocities or kinetic energies of the incident electrons. This

process leads to the emission of forward diffracted and transmitted electrons as well as

backscattered electrons. Inelastic events, on the other hand, results in a transfer of energy

to the solid with very little change in the electron trajectory. This process leads to

generation of secondary electrons, Auger electrons, characteristic and continuum x-rays,

visible light, electron-hole pairs, lattice vibrations and electron oscillations.

A number of these signals produced by the interaction of the incident electron beam

with the thin specimen are used in TEM. Elastically and inelastically scattered electrons,

secondary electrons and backscattered electrons are used for imaging. Inelastically

scattered electrons are also used for chemical microanalysis due to the characteristic

energy loss during inelastic scattering in a technique called electron energy loss

spectroscopy (EELS). The characteristic x-rays are also used for microanalysis through a

technique called energy dispersive x-ray spectroscopy (EDS).

TEM sample preparation

Specimen used in TEM studies requires that the samples are electron transparent.

This translates to a thickness of typically less than 100 nm. To achieve such, bulk

specimens must be sectioned and eletrothinned or ion milled to produce regions that

permit transmission of the electron.

TEM specimen from the bulk SiC samples were prepared by the standard

mechanical thinning method. A 3mm x 4 mm rectangle with a thickness of 1 mm was cut

from each specimen test bar with a low speed diamond saw. The cut sample was then

crystal-bonded to an aluminum polishing stub and wet polished to a thickness of- 100









am. A 3mm diameter disc was then cut from the polished sample using an ultrasonic

drill. The 3 mm disc was further polished to a thickness of- 50 [am using a 3 [am

diamond suspension in a precision dimpling machine and a flattening tool. The specimen

center was then subsequently thinned to a thickness of 20 [am using a dimpling tool. At

this point, the specimen center typically appears translucent. To make the specimen

electron transparent, argon ion milling was then performed in a Gatan duo-mill ion

milling machine. The typical parameters used for the SiC materials are: 4 kV

accelerating voltage, 1 aA ion gun current and a beam angle of 130. Ion milling time

generally depends on the type of SiC materials and usually takes from 3 to 10 hours. Ion

milling was carried out until a small perforation on the sample was observed. Due to the

brittle nature of the thin section of the specimen (the thin section falls off during TEM

investigation), several samples are usually made.

High resolution TEM (HRTEM)

High resolution transmission electron microscopy (HRTEM) is generally used to

obtain lattice images. To produce a lattice image, the transmitted beam and one

diffracted beam are passed through the objective aperture and combined. The phase

interference between these two beams yields the periodic intensity fringes present in the

image. Under the appropriate imaging conditions, there is a one-to-one correspondence

between the intensity fringes in the image and the atomic planes from which the electrons

were diffracted. In this case, the spacing of the fringes in the image is equivalent to the

spacing of the atomic planes.

Using lattice images, it is possible to examine the detail of grain boundaries, phase

interfaces and to image edge dislocations. In this dissertation, the lattice images are









primarily used for the examination of the grain boundaries and to determine presence of

intergranular films. The lattice images needed for this analysis therefore requires that it

contains details from the two grains and the interface. To achieve this, lattice image of

one of the adjacent grain is acquired first and the sample is tilted and oriented such that

lattice image of the other grain is obtained. In the absence of secondary phase, a direct

transition from the lattice images of the two grains can be observed. If the intergranular

film is present in amorphous form, absence of lattice fringes on the interface would be

apparent. If the intergranular film is crystalline, lattice fringes of distinctly different

appearance (different from the two lattice image of surrounding grains) could be

observed.

Obtaining lattice images was a tedious and time consuming process. It also

requires that the grain boundaries being studied are "head-on". In the case of the

materials studied in this dissertation, the nature of the microstructure (interpenetrating

elongated grains) severely limited the number of grain boundary investigated. For this

reason, only three or four high resolution images from each specimen were taken. The

high resolution images presented in this dissertation was generated on a JEOL JEM-

2010F FEG with a point to point resolution of 0.19 nm. The images were acquired with

the help of Kerry Seibien of the Major Analytical Instrumentation Center (MAIC) at the

University of Florida.

Energy dispersive spectroscopy (EDS)

The EDS data presented in this dissertation was taken with a JEOL 2010F equipped

with an Oxford INCA 200 system. The EDS spectra were generated under STEM mode

and used a probe size of approximately 5 nm. The electron beam hitting the sample,

however, is probably about 8 nm (based on the burn patterns observed on investigated









samples) due to the inhomogenieties in the electron delivery system. The size of the

probe used also limits the results obtained to a merely qualitative guide, i.e., only indicate

the elements present. This is particularly true for the compositions obtained on the grain

boundary since the grain boundary width (or thickness of the intergranular film) observed

is generally ~ 1 nm.

The advantage of EDS relative to other microanalytical technique is its ability to

generate x-ray spectra from, and therefore determine the composition of, very small

volumes of the specimen. Characteristic x-rays are generated from the interaction of

energetic electron and an atom. A sufficiently energetic electron can interact with an

atom and cause the ejection of tightly bound inner-shell electron, leaving the atom in an

ionized and highly energetic state. During subsequent de-excitation, an electron

transition occurs in which an electron from an outer shell drops inward to fill the inner-

shell vacancy. The change in energy (or the energy released) due to this electron

transition will be in the form of x-ray or an ejected outer-shell electron. Since the electron

structure of each atom is unique, the x-rays generated will also be unique and a

characteristic of a given element.

The decision to use the EDS results presented in this dissertation for qualitative

purposes only lies on the limitation of this technique. Since the EDS were taken from

TEM specimens, the number of x-ray counts generated will be minimal. The x-ray

generation volume will be smaller due to the thickness of the specimen. Thus, the

elemental detectability limits will be greatly diminished. This problem is compounded

by the elemental constituents of the materials being investigated. In the case of ABC-









SiC, the amount of B in the sample is only 0.5 wt. % (0.6 wt. % in a-SiC), which was

borderline with the quoted detectability limits in JEOL 2010F.

Energy filtered transmission electron microscopy (EFTEM)

EFTEM is a technique based on electron energy loss spectroscopy (EELS). EELS

involves analysis of the energy distribution of the inelastically scattered electrons in the

transmitted beam. Information pertaining to the chemical and bonding from the atoms in

the sample can be determined through this technique. During EELS, all the inelastically

scattered electrons are detected making the signal intensity much higher than in EDS

where characteristic x-rays just comes from small portion of inelastic scattering events. It

is thus a better technique than EDS for quantitative analysis. It also offers an advantage

in detecting low Z elements.

In EFTEM, an "imaging filter" is used to generate an image consisting of electrons

with specific energy losses. The image filter only allows electrons with an energy loss of

E + AE/2, where E is the selected energy and AE is the slit width of the energy window.

It is therefore possible to make chemically selective images from a given sample. Due to

the non-element specific background in an EELS spectrum, however, EFTEM images

can not be directly interpreted as elemental maps. Background removal has to be

accomplished first.

A method for removing background EELS spectrum and achieving elemental maps

from EFTEM images used in this dissertation is a technique called three window

elemental mapping. Three window elemental mapping allows generation of quantitative

images wherein the intensity is proportional to the EELS spectrum intensity under a

certain excitation edge. The schematic overview of this technique is shown in Figure 4-2.








The technique uses three images made up of electrons coming from three different energy

regions in the spectrum. Two pre-edge images are used to estimate the non-specific

background in each pixel of the image, by estimating the parameters of the background

model AE-r for each pixel. The third image (post-edge) is obtained from the energy

region which contains the excitation edge of interest. The estimated background from

each pixel is then subtracted from the post-edge image and an excitation specific image

remains.


AE
















CN
0_
rOi


- I I I I -


Excitatin spedfic
Image nributlan







AE-4


0
-


Figure 4-2. Schematic overview of the three-window technique.'0o

The EFTEM elemental maps presented in this dissertation was taken using a Tecnai

F30 equipped with FEG and GATAN GIF. The accelerating voltage used in the TEM is


S


~sc,


--I






45


300kV. All the elemental maps were taken with the help of Dr. Helge Heinrich in

Materials Characterization Facility at the University of Central Florida.














CHAPTER 5
RESULTS OF MECHANICAL PROPERTY, GRAIN BOUNDARY, AND TRIPLE
JUNCTION CHARACTERIZATION

Due to the large number of materials investigated, the results chapter of this

manuscript will be subdivided into four main sections based on the types of additive

used. The first section will cover the results for P-SiC sintered with 0.6 wt.% B 2 wt.%

C and 0 4 wt.% Al. The second section will cover the results for a-SiC sintered with Al

additive, while the third section will deal with the results obtained for a-SiC sintered with

A1N. The last section will present the result for a-SiC sintered with 3.1 wt.% A1203, the

solid state sintered a-SiC and the commercially available material SiC-N. Discussions

pertinent to the results presented in this chapter will be relegated to the next chapter of

this dissertation.

P-SiC Sintered with 0.6 wt.% B, 2 wt.% C and 0 4 wt.% Al (ABC-SiC)

Processing and Polytypes

The SiC ceramics fabricated from P-SiC starting powders with Al, B and C

additions sintered to near-theoretical density of pure SiC. The theoretical density of SiC

was calculated to be 3.2 g/cc. As shown in Table 5-1, the material with only B and C

additions gave a value near the theoretical density of SiC. Increasing addition of

aluminum has yielded density values which are progressively lower, with the highest

amount of aluminum showing the lowest density values.

The polytypes of SiC present in the manufactured ceramics, based on the Reitveld

analysis performed on the XRD patterns obtained from the material, are also shown in









Table 5-1. Hot pressing at 21000C for one hour enabled the complete P-to-a phase

transformation as shown by the absence of the cubic 3C polytype in the sintered material.

When only B and C were used as additives, the P-to-a phase transformation was reported

to begin at temperatures greater than 1950C.2,102-104 The addition of a small amount of

metallic aluminum or aluminum compounds can lower this phase transformation

temperature. As reported by Shinozaki et al.105 and William et al.,106 the onset of the P-

to-a phase transformation at approximately 1800C is even possible. Since the

temperature used in hot pressing the P-SiC in this study is higher than 1950C, it is not

surprising that complete P-to-a phase transformation was achieved.

Table 5-1. Density and polytypes of SiC ceramics from P-SiC starting powders.
Sample (wt. %) Density (g/cc) Phase Assemblage (wt. %)
3C 4H 6H 15R
P-SiC 0.6B -2C 3.17 + 0.01 0.0 0.0 93.1 6.9
-SiC 0.5A1 0.6B -2C 3.08 + 0.01 0.0 0.0 87.5 12.5
P-SiC 1Al 0.6B -2C 3.15 + 0.01 0.0 22.4 68.2 9.4
P-SiC 1.5Al 0.6B -2C 3.13 + 0.01 0.0 34.0 58.9 7.1
P-SiC 4A1 0.6B -2C 3.12 + 0.01 0.0 81.4 14.5 4.1


The polytype present in the manufactured ceramics varies with the amount of

aluminum added. When only B and C are added, the resulting polytypes are mostly 6H,

with the remainder being that of the rhombohedral 15R structure. As the amount of

aluminum is increased, the amount of 4H polytype in the fabricated ceramics also

increases, with a corresponding decrease in the amount of 6H polytype present. This

observed trend is similar with those reported for P-SiC with Al-B-C addition processed at

19000C.2









Microstructure and Mechanical Properties

The resulting microstructure of the hot-pressed P-SiC generally shows interlocking

platelike grains. As reported in the literature,24 formation of the platelike grains in

specimens that utilized P-SiC starting powders is due to the P-to-a phase transformation

that occurred during sintering. As shown in the previous section, complete P-to-a phase

transformation was achieved in the materials investigated, thus, the microstructure is

expected to consist of mainly elongated grains. Examples of the observed microstructure

in sample with 0, 0.5 and 1 wt.% aluminum are shown in Figure 5-1.













Figure 5-1. Transmission electron micrographs of the microstructure observed in SiC
ceramics fabricated from a P-SiC starting powder. a) microstructure found in
P-SiC-0.6B-2C, b) microstructure found in P-SiC-1Al-0.6B-2C, and c)
microstructure found in P-SiC-1.5A1-0.6B-2C. The marker shown in each
micrograph is for 2 tm.

The measured grain size, aspect ratio, hardness, toughness and observed failure

mode of the SiC materials sintered from P-SiC starting powders are shown in Table 5-2.

It is observed that correlation between the grain size and the amount of aluminum added

does not exist for the materials investigated. The grain size is smallest (4.1+0.5 [am) for

the sample with 1 wt.% Al added. This value, however, is not significantly different than

those measured for P-SiC with 0.5 wt.% Al addition if one considers the standard

deviation of the measured grain size. For the material with no Al additive, the grain size









measured is similar to those of otherwise similar material with 1.5 and 4 wt. % Al

additions and is higher than those measured from P-SiC with 0.5 and 1 wt. % Al. Thus,

considering that the grain morphology in these materials is generally comprised of

elongated grains, correlation between the measured grain size and the amount of

aluminum added is really non-existent.

Table 5-2. Grain size, aspect ratio, hardness, toughness and observed fracture mode in
SiC ceramics with P-SiC starting powders.
Sample (wt. %) Grain Aspect Toughness Hardness Fracture
Size Ratio (MPa HV1 Mode
(tm) m1/2) (GPA)
P-SiC 0.6B 2C 5.41.1 5.41.1 2.6 + 0.2 25.5 + 0.7 Transgranular
P-SiC 0.5Al 0.6B 4.60.2 5.91.3 2.6 + 0.1 20.9 0.9 Transgranular
2C
P-SiC 1Al 0.6B 4.10.5 5.40.9 2.7 + 0.1 21.5 + 0.9 Transgranular
2C
P-SiC 1.5Al 0.6B 5.20.4 3.90.6 6.1 0.3 15.6+ 1.2 Intergranular
2C
P-SiC 4A1 0.6B 5.30.5 5.11.0 6.1 0.2 20.8 0.9 Intergranular
2C


The aspect ratios of the elongated grains are deemed important in determining the

toughness of a given ceramic. Literature data suggests that the higher the aspect ratio of

the elongated grains in a microstructure, the higher the toughness achievable.9 This is of

course dependent on the toughening mechanisms operating in the material system. In the

case of SiC with a microstructure consisting of elongated grains, the toughening is

usually proposed to be associated with crack bridging, microcracking and crack

deflection mechanisms. And in these mechanisms, the aspect ratio of elongated grains

plays an important role. As shown in Table 5-2, the aspect ratio of the elongated grains

in the materials investigated are almost similar, with the exception of those measured in

P-SiC with 1.5 wt. % Al. The toughness value, however, does not show the same trend.









Additions of 0 to lwt.% Al yielded an SEPB toughness value of about 2.6 MPa m/2,

while additions of 1.5 and 4 wt. % Al showed a higher toughness value of 6.1 MPa m1/2

This trend in toughness value has a better correlation with the observed fracture mode of

the materials investigated. At low toughness values, the materials exhibited a

transgranular mode of failure while the higher toughness value materials showed an

intergranular fracture mode.

Hardness/toughness trade-off common in liquid phase sintered SiC ceramics was

also observed in the materials investigated. As shown in Table 5-2, high toughness

material generally shows a slightly lower hardness value while the low toughness

material shows a higher hardness. The hardness/toughness trade-off is readily apparent

when one considers the case of P-SiC with no Al addition and P-SiC with 1.5 wt. % Al

additive. In P-SiC with no aluminum addition, a Vickers hardness of -25.5 GPa and

SEPB toughness of -2.6 MPa-ml/2 were measured. In P-SiC with 1.5 wt. % Al, Vickers

hardness of -15.6 GPa and the SEPB toughness of -6.1 MPa-ml/2 were reported. These

two materials exhibits the extreme in the hardness and toughness measured from the

materials investigated. The low hardness value of P-SiC with 1.5 wt.% Al additive might

be due to the presence of amorphous grain boundary and the intergranular mode of failure

observed in the material which results in a possible grain boundary sliding during

indentation.

Achievement of high toughness at the same level of hardness was also

demonstrated in the P-SiC ceramics investigated. The Vickers hardness measured in the

P-SiC with 0.5, 1 and 4 wt. % Al are statistically similar while the toughness values are

markedly different. The SEPB toughness for P-SiC with 0.5 and 1 wt.% Al are 2.6 and









2.7 MPa-ml/2, respectively, while that for p-SiC with 4 wt.% Al addition is -6.1

MPa-ml2. This observation suggests the possibility of manufacturing a toughened hard

SiC ceramics.

Grain Boundary, Triple Junction and Other Secondary Phases

To ascertain the correlation between the grain boundary and triple junction

chemistry and crystallinity and mechanical properties of the SiC ceramics fabricated from

a P-SiC starting powder with Al, B and C additions, high resolution transmission electron

microscopy (HRTEM), energy dispersive spectroscopy (EDS), and energy filtered

transmission electron microscopy (EFTEM) were performed. High resolution imaging

was performed to determine the presence and crystallinity of intergranular films in the

SiC ceramics, while EDS and EFTEM were done to determine the chemical constituents

of the grain boundary, triple junction and other secondary phases found in the SiC

ceramics.

Grain boundary width and crystallinity

The results of HRTEM performed on the materials investigated are shown in

Figure 5-2 and summarized in Table 5-3. The grain boundary width and crystallinity was

observed to vary with the amount of aluminum added. At 0 to 1 wt.% Al additions, the

resulting grain boundary shows no intergranular phase and direct transitions of the fringe

pattern in-between grains are observed. Increasing aluminum content, from 1.5 to 4 wt.

%, leads to the formation of amorphous intergranular film with thickness of

approximately 1 nm. This is shown in Figure 5-2d and Figure 5-2e where the fringe

patterns of the two grains shown are delineated by a region where a fringe pattern was

non-existent.









Formation of secondary phases is expected in the presence of sintering additives in

concentrations well above their solubility limits in SiC. The solubility limit reported for

aluminum in SiC is 0.26 and 0.50 wt. % at 18000 and 20000C, respectively. The

solubility of boron in SiC is 0.1 wt. % at 2500C.107 Thus, based on the solubility limit

of aluminum in SiC, formation of an intergranular film at the grain boundary of the P-SiC

with 1 wt. % Al addition should occur. Since this is not the observed phenomenon, the

presence or lack thereof of the intergranular film cannot be explained in terms of the

solubility limit alone. Discussion along these lines will be presented in the next chapter

of this dissertation.


Figure 5-2. High resolution transmission electron micrographs of grain boundaries in 3-
SiC with Al, B and C additives, a) clean grain boundary in P-SiC-0.6B-2C. b)
clean grain boundary in P-SiC-0.5A1-0.6B-2C. c) clean grain boundary in P-
SiC-1Al-0.6B-2C. d) amorphous grain boundary in P-SiC-1.5A1-0.6B-2C and
e) amorphous grain boundary in P-SiC-4A1-0.6B-2C. The markers in a), b)
and c) are 2 nm while the markers in d) and e) are 5 nm.









Determination of the chemical constituents in the grain boundary, triple junction

and other secondary phases found in P-SiC materials investigated is presented in the next

section. As will be shown, the intergranular amorphous film was found to contain Al and

0.

Table 5-3. Grain boundary width and crystallinity of p-SiC with Al, B, and C additions.
Sample (wt. %) Grain Boundary Width Intergranular Film
P-SiC 0.6B -2C 0 none
-SiC 0.5Al 0.6B -2C 0 none
P-SiC 1Al 0.6B -2C 0 none
-SiC 1.5Al 0.6B -2C 1 nm Yes, Amorphous
p-SiC 4A1 0.6B -2C 1 nm Yes, Amorphous


Grain boundary, triple junction and secondary phase composition

It is known that the use of sintering additives may lead to the formation of

secondary phases in a given ceramic.13 The sintering additives attract impurities in the

starting powder, react with the native oxide on the particle surface and form a mass

transport medium during densification.13 Upon cooling, some of the mass transport

medium remain and become a crystalline or glassy phase in the resulting microstructure.

In the case of ABC-SiC, it has been reported that B and C does not form secondary

phases in the absence of aluminum.2 Carbon reacts with the native oxide on the surface

of the starting powder and form SiO and CO gas.2 This reaction reduces the amount of

oxygen available in the system. Thus, in the case of P-SiC-0.6B-2C, formation of

secondary phases is not expected. However, if the amount of carbon additive used is

more than the amount of native oxide available for the formation of SiO and CO, then,

presence of residual carbon is expected. This phenomenon is shown by the results obtain

for P-SiC-0.6B-2C.









In the presence of aluminum above its solubility limit in SiC, formation of

secondary phases is expected. Due to its low melting point (-6600C), the aluminum

metal will melt at the processing temperature used. B, C and O will then be transported

into the aluminum melt and form secondary phases. The liquid phase will then flow and

fill the porosity between SiC particles and engulf many grains. Thus, upon cooling,

secondary phases can be found in triple and multiple grain junctions. Due to the low

vapor pressure of aluminum, it is also possible that aluminum vapor will coat the SiC

particle surfaces and react with the native oxide on the powder surface as well as the

added boron and carbon and form a liquid grain boundary phase. Based on the above

discussion, the amount of secondary phases available in the final ceramics should vary

with the amount of aluminum added when the amount of B and C are fixed.

P-SiC 0.6B -2C

Based on the above discussion, presence of secondary phases in this material is not

expected, except for possible residual carbon in the system. Formation of intergranular

phases are also expected to be non-existent since boron and carbon does not form

intergranular phase in the absence of aluminum and will just get incorporated in the SiC

grains. As shown in Figures 5-3, 5-4 and 5-5, variation in the grain boundary and triple

junction composition was not observed. However, as shown in Figure 5-3, presence of

residual carbon, indicating that excess carbon additive was used in this material, was

detected. Figure 5-3 and 5-4 also suggests that the grain boundary and triple junction are

relatively "clean" in this material except for several residual carbon phase usually found

in the triple junction. Although not shown here, several metallic impurities were also










observed to segregate in the triple junction. This segregation of metallic impurities,

however, was observed for almost all of the materials investigated in this dissertation.

Figure 5-5 is an EFTEM of grain boundaries and associated triple junction and as

shown, variation in the elemental map suggesting possible elemental segregation in grain

boundary and triple junction was not observed. The absence of grain boundary and triple

junction phases might have contributed to the low toughness value of this material. The

presence of residual carbon might also act as a failure origins which may result in low

fracture strength achievable in this material.


3
'^^Bs
v^pB^


\I k


Bulk GuIam ia
- GIJIUIn Bouilan


05 1 15 2
ull Scale 113 ds Cursor 0 004 keV (963 ds)ke

Figure 5-3. STEM of a grain boundary and a secondary phase in P-SiC-0.6B-2C. The
accompanying EDS identify the secondary phase as residual carbon.

P-SiC 0.5A1 0.6B 2C

Results of the EDS and EFTEM studies performed in this material are shown in

Figures 5-6 to 5-9. The triple junction in this material was found to contain Al and 0-

rich phases, as shown in Figure 5-6. The grain boundaries observed in this material

generally do not contain secondary phase. However, the region shown in Figure 5-7

show some indication that secondary phase in the grain boundary might be present. Still,

the amount of aluminum and oxygen are too small to be believable. In addition, the EDS


Spectrum
/Residual Carbon



F1 ^ .


L^,









of this particular grain boundary was taken near a triple junction, thus, it is not

unexpected that a greater amount of aluminum and oxygen will be seen.


Figure 5-4. STEM of a triple junction in P-SiC-0.6B-2C. Variation in the triple junction,
bulk grain and grain boundary composition is non-existent.


Figure 5-5. EFTEM of a triple junction and grain boundary in P-SiC-0.6B-2C. Variation
in the elemental map is non-existent.

The assertion that no grain boundary phase is present can be clearly seen in Figure

5-8. As shown, the elemental maps generated by EFTEM do not indicate any variation in









composition across the grain boundary. This is particularly true for the map of

aluminum, oxygen and boron. As observed in P-SiC-0.6B-2C and most of the materials

studied, metallic impurities are also present in this material as shown in Figure 5-9. In

addition, the presence of B and C-rich secondary phase was also observed.


Figure 5-6. STEM of a triple junction in P-SiC-0.5A1-0.6B-2C. The accompanying EDS
were taken from the triple junction and bulk grains.


Figure 5-7. STEM of a grain boundary in P-SiC-0.5A1-0.6B-2C. The accompanying EDS
were taken from grain boundary and bulk grains.



















Grain Boundary
100 nm m















Figure 5-8. EFTEM of a grain boundary in P-SiC-0.5A1-0.6B-2C. Variation in
composition across the grain boundary is non-existent.


Spectrum 1
Metallic Impurities





,, ,-. .. ,,


B-C rich ,irins






0 05 1 15 2
ull Scale 69 cts Cursor -0 031 keV (107 cts) keV


Figure 5-9. STEM of a multiple grain junction in P-SiC-0.5A1-0.6B-2C containing
metallic impurities as shown by the corresponding EDS. B-C rich grains are
also shown.

Based on the value of solid solubility of aluminum in SiC reported by Tajima and

Kingery,107 formation of secondary phases either in the triple junction or grain boundary









of this material should be non-existent since the amount of aluminum used is close to the

solubility limit. However, smaller solubility value has been reported for aluminum in

SiC. Kinoshita et al.108 have reported that solubility of aluminum in SiC is -0.2 wt. % Al

based on their work in a-SiC with Al203 addition. The solubility limit was calculated

based on the change in d-spacing in SiC which was assumed to be proportional to the

amount of Al. Once the solubility limit was reached, then the d-spacing based on the x-

ray diffraction results should be constant. Constant d-spacing values was reached at -0.4

wt. % Al203 and the corresponding Al amount is -0.2 wt. % Al. Thus, based on this

value of solubility limit, which was deemed more appropriate for the present case, P-SiC-

0.5A1-0.6B-2C can have secondary phases in the triple junction and grain boundary. The

formation of secondary phase in the triple junction was observed but a secondary phase in

grain boundary was not. This suggests, as mentioned earlier, that solubility limit alone

cannot explain the formation of intergranular film in the grain boundary.

P-SiC-1AI-0.6B-2C

This material has an aluminum additive of amount well above the solubility limits

of aluminum in SiC and is expected to have secondary phase in the triple junction and

grain boundary. As shown in Figures 5-10 and 5-11, the triple junction of this material

contains Al-O rich composition. Based on the elemental map shown in Figure 5-11,

aluminum and oxygen segregates in the triple junction. Corresponding depletion of

silicon and carbon are also shown. The appearance of the triple junctions in this material

is different than those observed for P-SiC-0.5A1-0.6B-2C. In addition, the Al and O

content of the triple junction is greater. This observation supports the earlier assertion

that the amount of secondary phases in the triple junction and grain boundary will vary

with the amount of aluminum additive.


































0 05
ull Scale 619 cts Cursor -0 006 keV (970 cts


Figure 5-10. STEM of a triple junction in P-SiC-1A1-0.6B-2C. Accompanying EDS was
taken from bulk grains, grain boundary and triple junction.


Junction as-n Map


'trk


'.: ,".*-<'."> ".<
., -., :


1 i ... ,.7
.., ****. ... ..e


-V,


Aluminum Map


Figure 5-11. EFTEM of a triple junction in P-SiC-1Al-0.6B-2C. Segregation of Al and O
in the triple junction is observed.


:4I









The grain boundary of this material, however, does not show the expected

formation of intergranular phase. As shown in Figure 5-10, the grain boundary EDS

shows the presence of Al and 0, however, the amount present is comparable to the

amount observed in the bulk. Further evidence of non-formation of secondary phase in

the grain boundary is shown in Figure 5-12. The EFTEM elemental map does not show

any variation in composition across the grain boundary that may suggests possible

segregation of secondary phase.
























Figure 5-12. EFTEM of a grain boundary in P-SiC-1Al-0.6B-2C. Compositional variation
across the grain boundary is non-existent.

P-SiC 1.5A1 0.6B 2C

As discussed in the grain boundary width and crystallinity section of this chapter,

this material shows the presence of amorphous intergranular film in the grain boundary.

It was also shown previously that an increase in toughness and change in fracture mode

from transgranular to intergranular occurred in this specimen. Since the amount of










aluminum in this specimen is well in excess of the solubility limit of aluminum in SiC, it

is expected that formation of secondary phases in grain boundary and triple junction

should occur. As shown in Figures 5-13 to 5-16, formation of secondary phases in triple

junction and grain boundary was observed in this material. Figure 5-13 and Figure 5-14

shows that the triple junction is filled with Al and O rich composition. In Figure 5-14,

the elemental maps of aluminum and oxygen show definite segregation of these elements

in the triple junction. Corresponding depletion of carbon and silicon was also observed.

The boron map, however, only indicates that boron is almost undetectable in the grain

and the triple junction region. The amount of Al and O in the triple junction and in the

grain boundary in this material is considerably greater than those observed in the P-SiC

specimen discussed previously. The Al and O-rich composition in the triple junction is

believed to be in the form of Al203, as will be shown later.


0 05 1 15 2
uScale 01 Curr -0 017 keV (360 cs) keV












2II0


SI Spectrum 2

SGrain Boundary

c



0 o0 1 15 2
ull Scale 601 cts Curor -0 017 keV (471 cts) keV


Figure 5-13. STEM of a triple junction in P-SiC-1.5A1-0.6B-2C. The accompanying
EDS was taken from the triple junction, grain boundary and the bulk grain.


































Figure 5-14. EFTEM of a triple junction in P-SiC-1.5A1-0.6B-2C. Elemental maps
indicate segregation of aluminum and oxygen and depletion of silicon and
carbon at the triple junction.


Figure 5-15. STEM of a grain boundary in P-SiC-1.5A1-0.6B-2C. The accompanying
EDS was taken from the grain boundary and bulk grains. Definite aluminum
and oxygen segregation in the grain boundary is shown.


U. rn a

































Figure 5-16. EFTEM of a grain boundary in P-SiC-1.5A1-0.6B-2C. Elemental maps
shows segregation of aluminum and oxygen in the grain boundary.

As shown in Figure 5-15 and 5-16, the grain boundary of this material definitely

contains Al and O. Segregation of these elements in the grain boundary is readily

apparent in the elemental maps presented in Figure 5-16. Thus, based on the results of

EDS and EFTEM performed on the grain boundary, the amorphous intergranular film

shown in Figure 5-2d and 5-2e is due to the presence of Al and O, which is probably in

the form of A1203.

Other secondary phases were also observed in this material. Aside from the

common residual carbon and metallic impurities that segregate primarily to triple

junctions, presence of Al-O rich grains was also observed. This is shown in Figure 5-17

where the EDS of the lower grain indicate the presence of Al and O in that particular

grain. This phenomenon was not observed in the previous sample studied.


Aluminum Map























Al-O rich grain Al


C



o05 1 15
ull Scale 355 cts Cursor -0 017 keV (298 cts) keV

Figure 5-17. STEM of a grain boundary and associated bulk grains in P-SiC-1.5A1-0.6B-
2C. EDS taken at the lower grain shows presence of Al-O rich grain not
found in previous samples.

p-SiC 4A1 0.6B 2C

The triple junction and grain boundary of this material also contains secondary

phases rich in Al and 0. This observation is shown in Figures 5-18 to 5-20. Aside from

the usual metallic impurities and residual carbon, other secondary phases present in triple

junction and multiple grain junctions were also observed. Some of these secondary

phases are shown in Figure 5-18, 5-20 and 5-21. Residual carbon, which was common in

all the materials investigated, is shown in Figure 5-18. Figure 5-20 shows a multiple

grain junction rich in Al, B and C, whereas Figure 5-21 shows Al-C and B-C rich grains

found in the material. The presence of large number of secondary phases in this material

is not unexpected. The amount of aluminum in this material is quite high, 4 wt. % Al,

and as discussed in the previous section, the amount of secondary phases will be

proportional to the amount of aluminum additive.






























Figure 5-18. Residual carbon in triple junction of B-SiC-4A1-0.6B-2C.


Figure 5-19. EFTEM of a grain boundary in P-SiC-4A1-0.6B-2C. Elemental maps
confirms segregation of Al and O in grain boundary.

































Figure 5-20. EFTEM of a multiple grain junction in P-SiC-4A1-0.6B-2C. Box shows
region rich in Al and C.


\I-C' Iich 1 In B-C rich grain







1 .. 1 4
,i i 1 1 1














Figure 5-21. STEM and EDS of secondary phases found in P-SiC-4A1-0.6B-2C. The top
two EDS identifies the Al-C and B-C rich grain.


SiC Grains and^^H
ScondajTM'j~ry Phaes^Bf









a-SiC Sintered with Al Additive

Processing and Polytypes

Table 5-4 summarizes the density, chemical analysis and phase assemblage of the

a-SiC sintered with aluminum addition. a-SiC-1.65A1-0.5B4C-2C was included in this

group to determine the effect of B4C and C. As shown, the fabricated ceramics sintered

to a high density. It is apparent that an increase in the Al additive, (comparison of a-SiC-

1.65A1/2200C and a-SiC-3.3Al/2200C), leads to a decrease in the achievable density.

Increase in the hot pressing temperature, at the same aluminum additive level, also results

in a decrease of the density of the final ceramics. Addition of B4C and C results in a

much lower density as compared to those materials with Al addition only. The density

measured, however, is still about 98% of the theoretical density of SiC.

Table 5-4. Density, chemical analysis and phase assemblage of a-SiC sintered with Al
additive.
Sample (wt.%) Density Chemical Analysis Phase Assemblage (wt. %)
(g/cc) (wt. %)
Oxygen Nitrogen 3C 2H 4H 6H 15R
a-SiC-1.65Al, 3.22 1.01 0.006 0 0 8.1 83.5 8.5
19000C
a-SiC-1.65Al, 3.21 0.76 0.008 0 0 61.6 33.0 5.4
2100C
a-SiC-1.65Al, 3.21 0.46 0.006
22000C
a-SiC-3.3A1, 3.19 0.80 0.01 0 0 12.7 79.7 5.9
20000C
a-SiC-3.3A1, 3.18 0.27 0.007
22000C
a-SiC-1.65Al- 3.16 0.98 0.008 0 0 21.2 71.7 7.1
0.5B4C-2C,
21000C


The results of the chemical analysis performed on the fabricated ceramics may not

be an exact representative of the bulk due to the small sample volume used, less than 1









gram of material. The results, however, are still valid considering that much smaller

volume are being used for grain boundary and triple junction characterization. In order to

perform a meaningful comparison, the amount of oxygen on the starting powders used in

processing these materials should be considered. The amount of oxygen in the starting a-

SiC was analyzed to have 1.2 wt. % oxygen, whereas the Al starting powder was

measured at 0.76 wt. % oxygen. As shown in Table 5-4, the amount of oxygen retained

in the fabricated specimens varies as a function of amount of aluminum added and the hot

pressing temperature used. The retained oxygen decreases as a function of hot pressing

temperature, as observed for a-SiC with 1.65 wt. % Al additive. Doubling the amount of

aluminum additive at the same temperature, as shown by comparing a-SiC-

1.65A1/2200C and a-SiC-3.3A1/22000C, results in a decreased retained oxygen in the

system. In the case of a-SiC-1.65Al-0.5B4C-2C, the amount of oxygen retained is higher

than what was retained in a material with the same aluminum content, a-SiC 1.65A1. As

will be shown later, addition of B4C and C effectively removes the oxygen in the system.

However, whenever Al is present, either in the form of Al or A1N, considerable amount

of oxygen is retained.

The amount of retained oxygen in the fabricated ceramics may have an important

implication on the densification of these materials. The decreasing amount of retained

oxygen with increasing hot pressing temperature suggests that closed porosity was not

obtained until the hold temperature was reached. The presence of porosity provides an

escape route for the oxygen vapor generated during processing. Since heating at a higher

temperature will take more time than heating at lower temperature, oxygen losses will be

higher for materials hot pressed at a higher temperature. Higher retained oxygen content









when B4C and C were added with aluminum also implies a change in the densification

kinetics in these materials. It seems that the addition of B4C and C speed up the

densification rate and the material achieved a state of closed porosity earlier.

The starting a-SiC powders used in sintering the materials studied are primarily of

6H polytype. The use of aluminum additive in sintering SiC is known to promote the 6H

to 4H polytype transformation.2 As shown in Table 5-4, the 6H to 4H transformation

occurred in the materials fabricated. Comparison of the phase assemblage of a-SiC-

1.65A1 hot pressed at 19000C and 21000C also indicates an increase in the amount of

phase transformation as a function of temperature.

Microstructure and Mechanical Properties

The resulting microstructure of the hot-pressed a-SiC with aluminum additions are

shown in Figure 5-22. In the case of a-SiC with 1.65 wt. % Al additive, an increase in

the hot pressing temperature results in a change in grain morphology and grain size

distribution. Hot pressing at 19000C results in a near-equiaxed grains with an average

grain size of~ 1 m and an aspect ratio range of 1 to 2. Since the average grain size of

the starting powder is about 0.5 am, it is apparent that a grain growth had occurred during

processing. The small aspect ratio also indicates that anisotropic grain growth, as

illustrated by the grain elongation, occurred in a very small degree. Hot pressing at

21000C resulted in distinctly different grain morphology. Elongation of some of the

grains occurred while some of the grains remained equiaxed. In effect, a bimodal grain

morphology consisting of large elongated grains in a matrix of small equiaxed grains was

achieved at this hot pressing temperature. The small equiaxed grains average size is

about 1 am and the aspect ratio of the elongated grains at this temperature ranges from 2









to 5. Hot pressing at 22000C did not dramatically change the microstructure observed.

Large elongated grains in a matrix of small equiaxed grains are still apparent, however,

the aspect ratio of the elongated grains increased. The small equiaxed grains average size

is still about 1 [m but the aspect ratio range increases to about 3 to 6.

























Figure 5-22. Transmission electron micrographs of the microstructures observed in a-SiC
sintered with Al additive. The magnifications used are not the same due to the
differences in grain sizes observed.

Increasing the amount of Al additive in the material did not drastically change the

resulting microstructure. As shown in Figure 5-22, the microstructure of a-SiC with 3.3

wt. % Al additive hot pressed at 20000C is similar to the microstructure observed in a-

SiC-1.65A1 hot pressed at 19000C. The only difference is the slightly bigger grain sizes

observed in the former material. The bigger grain size, however, is more suitably

correlated with the higher temperature. A more suitable comparison to determine the

effect of increased aluminum additive would be that of a-SiC-1.65Al and a-SiC-3.3A1









hot pressed at the same temperature of 22000C. In both materials, the microstructure

shows large elongated grains in a matrix of fine equiaxed grains. Although the particular

TEM of the microstructure shown for a-SiC-3.3A1 hot pressed at 22000C consisted of

larger volume of elongated grains, on the average, the volume of elongated grains in both

materials is about the same. It is evident, however, that the aspect ratio of some of the

elongated grains in a-SiC-3.3A1 is larger than the aspect ratio observed for the elongated

grains in a-SiC-1.65Al.

The effect of B4C and C addition in the resulting microstructure of a-SiC sintered

with Al is readily apparent. As shown in Figure 5-22, the microstructure of a-SiC-

1.65A1-0.5B4C-2C exhibits large grains indicating a high grain growth in the material.

The marker bar for the transmission electron micrograph of the microstructure of this

material is set at 5 am. Average grain size is about 5 am and the aspect ratio of the

elongated grains is about 1 to 2. As anticipated, the addition of boron and carbon

promoted normal grain growth in a-SiC ceramics.

The measured hardness, toughness and observed fracture mode from the materials

sintered with a-SiC with aluminum additions are shown in Table 5-5. The fracture

modes of the materials studied are gleaned from the appearance of the Vickers

indentation and fracture surfaces which was shown in Figure 5-23 and 5-24, respectively.

The observed fracture mode in a-SiC sintered with Al addition remains mixed

independent of hot pressing temperature and the amount of aluminum added.

Direct correlation between the observed microstructure and measured hardness and

toughness exists for a-SiC sintered with aluminum additives. In the case of a-SiC with

1.65 Al additions, hot pressing at 19000C resulted in the highest hardness measured and







73


the lowest toughness value. The high hardness can be attributed to the grain size and

morphology observed in this material, i.e., equiaxed grains with an average grain size of

about 1 [m which was the smallest grain observed for a-SiC sintered with Al additive.

Increasing the hot pressing temperature resulted in a decreased hardness and increased

toughness in the resulting material. As discussed above, increasing the hot pressing

temperature in a-SiC-1.65Al also resulted in a change in microstructure from small

equiaxed morphology to a self-reinforced microstructure consisting of large elongated

grains in a matrix of small equiaxed grains.

Table 5-5. Mechanical pro erties of a-SiC sintered with Al additive.
Sample (wt. %) Hardness Toughness Fracture Mode
(HV1, GPa) (MPa m1/2
a-SiC-1.65Al, 19000C 25.30.7 4.00.2 Mixed
a-SiC-1.65Al, 21000C 22.1+0.7 4.70.4 Mixed
a-SiC-1.65Al, 22000C 20.8+0.3 5.70.1 Mixed
a-SiC-3.3A1, 20000C 22.1+0.6 4.20.1 Mixed
a-SiC-3.3A1, 22000C 20.50.5 6.80.1 Mixed
a-SiC-1.65A1-0.5B4C-2C 20.30.3 3.1+0.1 Transgranular


&d


1.65AI. 1900C 65A5AI. 21000C i" 1.65AI.








3.3AI. 200(C 3.3AI. 2200C 1.G5AI-C
A *I
; *. .







^r :. ;"",, *, /
.4 .. .. *



Figure 5-23. Optical micrographs of Vickers indentation in a-SiC sintered with
aluminum additions. Marker bars shown are 10 km.


2200C
,a, .


l....



























Figure 5-24. SEM of fracture surfaces of a-SiC sintered with aluminum additions. The
fracture mode remains mixed for all materials.

The decreasing hardness as a function of increasing hot pressing temperature in a-

SiC-1.65A1 seems to correlate with the anisotropic grain growths which have resulted in a

larger average grain size at higher processing temperature. This observation is in stark

contrast with the ABC-SiC system presented earlier, SiC-YAG and solid-state SiC which

has shown little or no hardness dependence on grain size.9 The differences in hardness

dependence in grain size may be due to the fracture mode, which remains mixed in all

processing temperature, observed in a-SiC-1.65A1. As suggested by Rice et al,o

hardness dependence on grain size is possible for ceramics fracturing intergranularly

when the grain size is sufficiently smaller than the Vickers indent size. As shown in

Figure 5-23, this condition may be true for a-SiC-1.65Al ceramics at the hot pressing

temperatures used. Another possibility on why the hardness decreases with increased in

grain size may be due to the trend in ceramic materials wherein the inherent flaw sizes

present scales up with the grain size. Since the mechanical properties in ceramics are

highly dependent on flaw sizes in the material, it is not surprising that hardness decreases









with increasing grain size since increasing grain size would imply an increase in the

inherent flaw size present.

As shown in Table 5-5, the toughness measured in a-SiC-1.65Al also correlates

with the hot pressing temperature and the resulting microstructure. Increasing toughness

values were measured at increasing hot pressing temperature. As discussed previously,

an increase in the aspect ratio of the elongated grains in the materials processed at higher

temperature was also observed. Literature report indicates that the presence of key

elongated grains in the microstructure greatly affects the toughness of toughened

ceramics.11-13 As observed by Lee et a13 an increase in aspect ratio and volume of key

elongated grain results in an increase toughness in silicon carbide possessing a "self-

reinforced" microstructure. Although increased in volume of elongated grains in the a-

SiC-1.65A1 was not observed, increased in the aspect ratio of elongated grains were.

Thus, the resulting increased in aspect ratio of the elongated grains at higher processing

temperature is deemed primarily responsible for the toughening observed for a-SiC

sintered with Al additive. This also applies for the observed toughness value in a-SiC-

3.3A1 processed at 22000C since the resulting microstructure in this material also showed

elongated grains of aspect ratio higher than those observed for a-SiC-1.65Al processed at

the same temperature.

As shown in Table 5-5, B4C and C additions in a-SiC-1.65Al resulted in low

hardness and toughness values. This observation was primarily due to the microstructure

achieved in this material. As shown, the grain size is relatively large and the grains are

mostly equiaxed. Presence of elongated grains with high aspect ratio in this material was









not observed thus the toughening mechanisms responsible for high toughness in a-SiC

with pure aluminum additions are not operable in this material.

Grain Boundaries, Triple Junctions and Secondary Phases

a-SiC 1.65 wt. % Al hot pressed at 19000C

Figure 5-25 shows several grain boundaries studied in a-SiC-1.65Al hot pressed at

19000C. As shown, presence of amorphous intergranular film in the grain boundary of

this material is apparent. However, some of the grain boundaries observed, as shown in

Figure 5-25d, also show absence of intergranular film. The amount of aluminum additive

used in this material is well in excess of the solubility limit of aluminum in SiC. The

amount is even higher than the amount of aluminum used in 3-SiC where intergranular

amorphous film in the grain boundaries were observed. Thus, it is expected that

formation of intergranular film in the grain boundary of this material should occur.




















Figure 5-25. High resolution transmission electron micrographs of several grain
boundaries studied in a-SiC-1.65Al hot pressed at 19000C. a), b), and c)
Shows amorphous grain boundary with a width of about 1 nm. d) Grain
boundary with no amorphous intergranular film.







77


The composition of grain boundary and triple junction phases in this material are

shown in Figures 5-26, 5-27 and 5-28. The grain boundary, as typified by Figure 5-26, is

generally filled with Al and O rich phase. The Al and O concentration in the grain

boundary, as shown by the EDS data, are generally higher than those found in the nearby

grains. This indicates that definite segregation of Al and O in the grain boundary really

occurred. As shown in Figure 5-27 and 5-28, segregation of Al and O in the triple

junctions of this material was also observed. Electron energy loss spectra (EELS) were

taken from the triple junction and an example of which is shown in Figure 5-19. The Al

and O peaks were then compared with literature data for crystalline A1203 and found to

be consistent. The crystallinity of the triple junction is further ascertained via high

resolution lattice imaging of the triple junction shown in Figure 5-30. And as shown,

lattice fringes consistent with crystalline structure were observed.


Spectrum 4

Bulk Grain




S\ Al
0 05 1 15 2
ull Scale 131 cts Cursor -0021 ke\ (257 ct ke


-


Spectrum 1

Grain Boundary




Al
I ,
05 1 15 2
.:- ts Curor -0012 eV(539 cs) keV
Spectrum 5

Grain Edge





0 05 1 5 2
l Scale 113 cts Cursor r 001 keV(589 cts) keV


Figure 5-26. STEM and EDS of a grain boundary and associated grains in a-SiC-1.65Al
hot-pressed at 19000C.










Trip e Junction


0
'r \


Spectrum 1

Ai


0 05 1 15 2
ull Scale 130 cts Cursor -0 005 k V (5100 cts) ke\


0 05 1 15 2
-ull Scale 149 cts Cursor -0 005 keV (5450 cis) keV
S Spectrum 3
S--`lB Ik Grain



c
0 05 1 15 2
ull Scale 210 cts Cursor -0 016 keV (404 cts) keV


Figure 5-27. STEM and EDS of a triple junction in a-SiC-1.65A1 hot-pressed at 19000C.
The triple junction phases observed are generally Al and O rich.


Grain Boundary and
Triple Junction Ak


ilm:


;616
c,
-"" ;" -?5


Figure 5-28. EFTEM of a triple junction and grain boundary in a-SiC-1.65A1 hot pressed
at 19000C. Al and 0, in the form of Al203, segregates at the triple junction.


S Spectrum 2


Triple Junction


Oy


I

















Figure 5-29. EELS taken from a triple junction in a-SiC-1.65Al hot pressed at 19000C.
The Al and O peaks are consistent with crystalline A1203.




Al C)







Figure 5-30. High resolution images and EDS of triple junctions in a-SiC-1.65Al hot
pressed at 19000C. a) and b) are high resolution images, c) is an EDS taken
from the triple junction.

The appearance of the triple junction phase in the manufactured ceramics may yield

information about the process in which the liquid phase distributes itself to cover the

surfaces of the particulate solids. By measuring the dihedral angle, based on STEM

images and EFTEM elemental maps, one can determine if complete penetration of the

triple junction and grain boundary by the liquid phase occurred. In the case of a-SiC-

1.65A1 hot pressed at 19000C, the measured dihedral angles range from 15 to 600. This

implies that a continuous liquid phase penetration of all triple grain junctions and partial

penetration of the grain boundary occurred in this material. This also implies that the

composition of the triple junction phase and the grain boundary phase (or the

intergranular amorphous film) is the same, i.e., A1203.


Al K-peak in A1203


S0 K-peak in A1203









a-SiC 1.65 wt. % Al hot pressed at 21000C

Figure 5-31 shows high resolution micrographs of several grain boundaries

observed in a-SiC-1.65Al hot pressed at 21000C. As shown, grain boundaries with

distinct grain transitions which indicates absence of intergranular phase, and grain

boundary with amorphous intergranular film are also present in this material. For grain

boundaries with amorphous intergranular film, the width of the film observed is typically

about 1 nm. The composition of the amorphous intergranular film or the secondary phase

in the grain boundary is generally Al and 0. As shown in Figure 5-32, segregation of Al

and O in the grain boundary is generally observed. It thus seem conflicting that the high

resolution images of the grain boundary shows absence of intergranular films while the

EDS and EFTEM elemental maps suggests that segregation of Al and O occurs which

implies that it should be present as an intergranular film.


Figure 5-31. HRTEM of several grain boundaries in a-SiC-1.65Al hot pressed at 21000C.
a), b), and c) are grain boundaries with no intergranular phase. d) shows a
grain boundary with amorphous intergranular film and a crystalline triple
junction.
































Figure 5-32. EFTEM of a grain boundary in a-SiC-1.65Al hot pressed at 21000C.
Elemental maps shows segregation of Al and O in the grain boundary.

The triple junction phase composition observed in a-SiC-1.65Al hot pressed at

21000C is shown in Figures 5-33, 5-34 and 5-35. In all these figures, the composition of

the triple junction phase was shown to be Al and 0. EELS of the triple junction also

revealed that the phase present is in the form of crystalline A1203. The three figures also

show different appearances of the triple junction phases. Measured dihedral angle based

on the STEM images and EFTEM of triple junction phases in a-SiC-1.65Al hot pressed

at 21000C ranges from 15 to 900. Based on this range of dihedral angles, presence of

liquid phase penetrating all three grain junctions, and liquid phase that only partially

penetrates the three grain junctions, are possible. In addition, partial liquid phase

penetration of the grain boundary would still occur. This explains the observed absence

of secondary phase in some of the grain boundaries investigated.






































Figure 5-33. STEM and EDS


of triple junction in a-SiC-1.65A1 hot pressed at 21000C.


SSpectrum 5

Bulk Grain



SII



Grain Boundary





0 05 1 1 5
ull Scale 517 cts Cursor -0 015 keV (3973 cts) keV


Figure 5-34. STEM of triple grain junction phase with lower dihedral angle in a-SiC-
1.65A1 hot pressed at 21000C.


S Spectrum 1





J XA


Triple Junction

































Figure 5-35. EFTEM of a triple junction in a-SiC-1.65Al hot pressed at 21000C.

a-SiC 1.65 wt. % Al hot pressed at 22000C

The high resolution images of the grain boundary in this material, as shown in

Figure 5-36, show presence of clean grain boundary and grain boundary with amorphous

intergranular films. The thickness of amorphous intergranular film (IGF) in this material

is about 1 nm, similar to the IGF observed in specimen with the same composition hot

pressed at lower temperatures. EDS and EFTEM elemental maps taken from the grain

boundary of a-SiC-1.65Al hot pressed at 22000C showed that the grain boundary phase is

Al and O rich and is similar to the previous materials studied. The triple junction phase

composition, as shown in Figures 5-37 and 5-38, is also consistent with that of Al203.

The triple junction phase, based on high resolution imaging, is also crystalline. The grain









boundary and triple junction phases are deemed similar, i.e., both are A1203, based on the

dihedral angles measured which ranges from 40 to 900.


Figure 5-36. HRTEM of grain boundaries in a-SiC-1.65Al hot pressed at 22000C.


Figure 5-37. STEM of a triple junction in a-SiC-1.65Al hot pressed at 22000C. EDS
shown were taken from triple junction, grain boundary and bulk grains.