The Pennsylvania State University

The Graduate School

College of Earth and Mineral Sciences

MICROSTRUCTURE, CONNECTIVITY, AND MECHANICAL BEHAVIOR IN ALUMINA

AND SIALON MICROSTRUCTURE COMPOSITES

A Thesis in Materials Science and Engineering by Brandy A. Soublet

© 2011 Brandy A. Soublet Submitted in Partial Fulfillment of the Requirements for the Degree of

Master of Science

August 2011

The thesis of Brandy A. Soublet was reviewed and approved* by the following:

Gary L. Messing Distinguished Professor of Ceramic Science and Engineering Head, Department of Materials Science and Engineering Thesis Advisor

David J. Green Professor of Ceramic Science and Engineering

James H. Adair Professor of Materials Science and Engineering

*Signatures are on file in the Graduate School.

ii

ABSTRACT

Composite designs are useful in developing new materials with unique property sets tailored for any application. Mitigating the processing difficulties associated with processing dissimilar materials, however, can be expensive and limit their application. In this work, a new class of materials recently proposed called microstructure composites were fabricated and mechanically tested to gain insight into the mechanical response of these composites. Microstructure composites are composites in which the constituent components are not dissimilar materials, but different microstructures of a material with the same composition. The material used in this work is α-Al2O3 and the constituent microstructures are a fine-grained, equiaxed Al2O3 microstructure and a textured Al2O3 microstructure. Textured alumina grains exhibit crack deflective properties due to the low fracture energy of the basal plane, and can therefore be used to toughen alumina. Crack deflection increases the amount of energy absorbed by a ceramic during fracture and can be quantified by work of fracture. The textured microstructure is achieved by templated grain growth (TGG); template α-Al2O3 particles added to a colloidal slurry are aligned during forming. The addition of 0.14 wt% CaO + SiO2 liquid- forming dopant induces anisotropic growth of the template particles during densification.

Microstructure composites can be fabricated with a variety of complex architectures according to the concept of connectivity, as used for piezoelectric applications. The connectivity of a composite describes the dimensions in which each component is connected to itself. A novel processing method, slurry co-casting, which involves the simultaneous tape casting of two ceramic slurries, was utilized to cast composite tapes. These were subsequently cut and stacked in specific sequences to create microstructure composites of 1-1, 2-2, and 3-3 connectivities. Microstructural control between the equiaxed and textured layers was achieved by sintering at 1550°C. A sharp interface between the constituent components ensured that distinct connectivities were achieved. Combining constituent components in such specific configurations increases the control a researcher has over the properties of the material.

iii

The alumina microstructure composites were mechanically tested in equibiaxial flexure using the ring-on-ring method. Load/deflection curves, fracture surfaces, and failure behavior were analyzed. Crack deflection and step-wise fracture were observed for all connectivities tested. Extended crack deflection (≥ 1 mm) was observed in only the 1-1 and 3-3 connectivities. So-called “graceful failure” was observed in all connectivities and work of fracture reached values in excess of 7 kJ/m2. The mechanical response of the 2-2 connectivity was variable; it was the only connectivity to exhibit catastrophic, instead of graceful, failure in some cases. The 2-2 connectivity, however, is unique from the 1-1 and 3-3, in that due to the nature of the architecture and the chosen loading geometry, it is possible that a crack traveling through the specimen would never encounter a reinforcing, textured layer. For this reason, the 2-2 connectivity studied in this work is not recommended for structural applications. The 1-1 and 3-3 connectivities, however, demonstrated crack deflective and energy absorbing properties, which increase the toughness of the material. The fabrication and mechanical testing of equiaxed/textured microstructure composites in this work offered new insights into the effect of connectivity (namely the 1-1, 2-2, and 3-3) on the mechanical response of alumina.

Microstructure composites of SiAlON were also fabricated in which the constituent microstructures were the α- and β-SiAlON phases. α-SiAlON was used for its high hardness and β-SiAlON was used for its high toughness, due to the elongated nature of its grain structure. Microstructure composites of these phases could exploit the advantageous mechanical properties of both. Many processing issues, however, including failure in both the green state and during densification (by hot-pressing), were encountered. Initial mechanical testing of the 3-3 connectivity in SiAlON microstructure composites reveals catastrophic failure and low biaxial strength. This behavior is most likely the result of one of the processing issues encountered in the fabrication of these composites, that is, the “pre-cracking” of the green sample in the hot- press prior to densification.

iv

Table of Contents

LIST OF FIGURES

LIST OF TABLES

1 Introduction...... 1 1.1 Statement of Problem...... 1 1.2 Scientific Approach ...... 5 1.3 Organization of Thesis...... 6 1.4 References...... 7 2 Literature Survey ...... 8 2.1 Toughening Mechanisms ...... 8 2.2 Materials ...... 13 2.3 Mechanical Testing...... 14 2.4 References...... 15 3 Fabrication and Mechanical Properties of 1-1, 2-2, and 3-3 Alumina Microstructure Composites...... 18 3.1 Introduction...... 18 3.2 Experimental Procedure...... 20 3.2.1 Sample Fabrication...... 20 3.2.2 Mechanical Testing ...... 24 3.3 Results and Discussion ...... 25 3.3.1 Co-casting Observations ...... 26 3.3.2 Microstructural Control...... 27 3.3.3 Work of Fracture Calculation...... 29 3.3.4 Fracture Behavior and Work of Fracture ...... 31 3.3.5 Equibiaxial Strength...... 50 3.4 Conclusion ...... 52 3.5 References...... 53

v

4 Fabrication of α-/β-SiAlON Microstructure Composites: Processing Issues and Initial Results...... 55 4.1 Introduction...... 55 4.2 Design and Testing of the α- and β- SiAlON Phases ...... 56 4.2.1 Design Considerations...... 56 4.2.2 Experimental Procedure ...... 56 4.2.3 Results ...... 57 4.3 SiAlON Microstructure Composite Fabrication and Processing Issues ...... 61 4.3.1 Microstructure Composites ...... 61 4.3.2 Experimental Procedure ...... 62 4.3.3 Processing Issues...... 63 4.3.4 Initial Mechanical Testing Results...... 67 4.4 Conclusion ...... 68 4.5 References...... 68 5 Future Work...... 70

vi

List of Figures

Figure 1.1. The ten possible connectivities in a two-phase composite system. (Picture adapted by M. Fleck, Pennsylvania State University.)1 ...... 2

Figure 1.2 Fracture behavior in 2-2 alumina microstructure composite bend bars.2 ...... 3

Figure 1.3 Schematic (top view) of co-casting apparatus (a) in which separate slurries are poured into individual reservoirs for simultaneous tape casting. The arrow denotes the direction of the cast. The reservoirs are divided by separator fins (b); the tapered ends of the separator fins allow the slurries to merge just prior to passing under the doctor blade.2 ... 5

Figure 2.1. Particulate (SiC whiskers, left, and alumina agglomerates, right) reinforcement in alumina.6,7 ...... 9

Figure 2.2. Crack deflection in laminate SiC with weak graphite interlayers and load/deflection curve showing multiple “pop-in” features characteristic of graceful failure.9...... 10

Figure 2.3. Crack arrest by weak monazite interface layer in alumina matrix.11 ...... 10

Figure 2.4. Crack deflection in laminate alumina with varying amounts of porosity in porous alumina interlayers.12 ...... 11

Figure 2.5. Textured/equiaxed alumina microstructure laminate fracture specimen showing crack deflection, bifurcation, and arrest.18...... 12

Figure 3.1. Stacking configuration for 2-2 (tape-cast tapes) and 1-3 (co-cast tapes) connectivities.5 ...... 18

Figure 3.2. Connectivities, their respective tape stacking configurations, and sample schematics for the 1-1, 2-2, and 3-3 connectivities, respectively. The dark stripes represent the textured material and the light stripes represent the equiaxed material. The direction of casting, and therefore the orientation of the template particles, is parallel to the stripes. 19

Figure 3.3. Tape stacking sequences for the 2-2 series (a) and 2-2 parallel (b) connectivities. The nomenclature refers to the manner is which the microstructural components see the load. In the series, the equiaxed and textured layers see the load individually while in the parallel, the equiaxed and textured layers see the load simultaneously (in the same layer)...... 20

Figure 3.4. Micrographs of as-received AKP-50 (a) and alumina platelets (b)...... 21

vii

Figure 3.5. Flow chart for processing templated and equiaxed alumina microstructure composites...... 23

Figure 3.6. Schematic of ring-on-ring fixture for equibiaxial strength testing. Adapted.7 ...... 24

Figure 3.7. A closer look of the co-casting apparatus (top view) in which separate slurries are poured into individual reservoirs for simultaneous tape casting (a). Separator fins are used to divide the slurry reservoirs. Tapered edges (b) at the end of the separator fins allow for the component slurries to merge just before passing under the doctor blade (c).5 ...... 27

Figure 3.8. Composite tape design, stacking sequence, and schematic of final textured/equiaxed

Al2O3 3-3 microstructure composite. The schematic represents the “ideal” tape, with straight component stripes, no bleeding of slurries into adjacent components, and uniform thickness across the tape; when these factors are achieved, one can obtain a specimen of accurate connectivity...... 27

Figure 3.9. Micrographs demonstrating microstructural control in 3-3 fine-grained/textured

Al2O3 microstructure composites. Relative layer thicknesses (related to original tape thicknesses) are shown in (a); these thicknesses are typical of all composites made. Micrograph (b) shows a closer look of the clear interface (marked by the dashed line) between the distinct equiaxed and textured layers...... 28

Figure 3.10. Polished surfaces (a and b) and fracture surfaces (c and d) of the fine-grained, equiaxed and textured microstructures in a 3-3 microstructure composite, respectively. Microstructural control was achieved through doping (MgO for equiaxed components

and SiO2+CaO for textured components) of the ceramic slurries...... 29

Figure 3.11. Typical load/deflection curve...... 30

Figure 3.12. Schematic of the 2-2 parallel microstructure composite. The inset also shows the textured (left) and equiaxed (right) microstructural components.8 ...... 32

Figure 3.13. Step-wise fracture in a 2-2 parallel alumina microstructure composite. The cracks deflect ~100-200 µm in the textured layers before intersecting the equiaxed layers...... 33

Figure 3.14. The nature of crack propagation differs between the equiaxed and textured layers. In the equiaxed layers, the propagation is transverse (direction indicated by the arrow). In the textured layers, the propagation is longitudinal, along the basal faces of the textured grains...... 34

Figure 3.15. Micrograph of a 2-2 parallel alumina microstructure composite (a) and the respective load/deflection curve (b). While the composite shows extensive step-wise fracture and ~100 µm deflection along the textured layers, the load deflection curve

viii

indicates catastrophic failure. Still, the “work of fracture” of this specimen, 7.540 kJ/m2, is the highest seen in this work...... 36

Figure 3.16. Load/displacement curves of 2-2 parallel alumina microstructure composites which failed catastrophically. These samples represent average and high “work of fracture” values for this connectivity. The specimen on the left had a biaxial flexure strength of 167 MPa, and the specimen on the right, 247 MPa...... 37

Figure 3.17. Load/displacement curves of 2-2 parallel alumina microstructure composites which failed non-catastrophically. Each pop-in feature represents failure of a layer and loading of the next; this fracture behavior is known as “graceful failure.” The specimen on the left has a biaxial flexure strength of 154 MPa, and the specimen on the right, 194 MPa. 37

Figure 3.18. Schematic of a 1-1 microstructure composite...... 39

Figure 3.19. Micrograph of a 1-1 alumina microstructure composite in which crack deflection begins within the textured layered. That is, the crack propagates transversely through the interface between the equiaxed and textured regions (denoted by the dotted line) and deflects ~30 µm into the textured layer...... 40

Figure 3.20. Micrograph of a 1-1 alumina microstructure composite demonstrating both deflection through a textured layer (denoted by the black arrows) and no deflection (transgranular fracture) through a textured layer (denoted by the dotted lines)...... 41

Figure 3.21. The interface between the equiaxed and textured layers in a 1-1 alumina microstructure composite in which no deflection is observed through the textured layer. Fracture through the equiaxed layer (above dotted line) is intergranular, and fracture through the textured layer (below dotted line) is transgranular...... 42

Figure 3.22. Micrograph of a 1-1 alumina microstructure composite showing extended crack deflection (~1 mm). The arrows indicate two directions of crack deflection, a result of the biaxial loading. This composite had a biaxial flexure strength of 128 MPa and a “work of fracture” of 2,740 kJ/m2...... 43

Figure 3.23. Load/deflection curves of 1-1 alumina microstructure composites, all showing “pop- in” characteristics indicative of graceful failure. The respective workof fracture values are shown...... 44

Figure 3.24. Schematic of a 3-3 microstructure composite...... 45

Figure 3.25. Micrographs of 3-3 alumina microstructure composites showing two different mechanisms of deflection through a texture layer: a) deflection begins at the interface between the equiaxed and textured regions, and b) deflection begins within the textured

ix

layer, after propagating transversely through the interface between the equiaxed and textured layers, indicated by the dotted lines...... 46

Figure 3.26. Micrograph of a 3-3 alumina microstructure composite that exhibits extended crack deflection (~1.5 mm), denoted by the arrow. This behavior can drastically change the “work of fracture” of a specimen. The biaxial flexure strength of this specimen is 145 MPa and the “work of fracture” is 3.153 kJ/m2...... 47

Figure 3.27. Micrograph of a 3-3 microstructure composite showing deflection along textured layers. The ends of the textured layers (or “stripes” in the green tape) are denoted with the arrows. Had the propagating crack traveled just left of the end of these layers, it would not have been deflected...... 48

Figure 3.28. Load/deflection curves of 3-3 alumina microstructure composites, all showing “pop- in” characteristics indicative of graceful failure. The respective “work of fracture” values are shown...... 49

Figure 3.29. Two-parameter Weibull plots for 1-1, 2-2 parallel, and 3-3 connectivity samples tested in ring-on-ring equibiaxial flexure...... 51

Figure 3.30. Fracture patterns for 2-2 connectivity samples; (a) and (b) are samples from the lower flaw segment and (c) and (d) are samples from the upper flaw segment. Cracks are highlighted for better viewing. No distinct differences in the crack patterns can be discerned between the upper and lower flaw segments...... 52

Figure 4.1 α-SiAlON a) fracture surface and b) polished surface etched for 15 min in KOH..... 58

Figure 4.2 4 wt% yttria-doped β-SiAlON a) fracture surface and b) polished surface etched for 15 min in KOH...... 58

Figure 4.3. Polished surfaces of β-SiAlON doped with 7 wt% a) yttria and b) yttrium nitrate. Samples etched in molten KOH for 15 min each. Arrows indicate pockets of etched liquid phase...... 59

Figure 4.4. X-ray diffraction patterns for the α- and β-SiAlON phases, respectively...... 61

Figure 4.5. Connectivities and their respective tape stacking configurations. The dark stripes represent α-SiAlON and the light stripes represent β-SiAlON...... 62

Figure 4.6. α-SiAlON green body which failed during the burnout cycle...... 64

Figure 4.7. β-SiAlON samples which failed during hot pressing: a) complete SiAlON degradation, and b) partial degradation resulting in a core-shell structure...... 65

x

Figure 4.8. SiAlON microstructure composite showing cracks originated during the hot-pressing cycle...... 66

Figure 4.9. Typical load/deflection curves of the α-/β-SiAlON 3-3 microstructure composites tested. No “pop-in” features are observed, which is indicative of catastrophic failure... 67

xi

List of Tables

Table 3.1. Mean “work of fracture”, standard deviations, and number of specimens tested for connectivity samples...... 50

Table 3.2. Weibull moduli, characteristic strengths, and number of specimens tested for connectivity samples...... 50

xii

1 Introduction

1.1 Statement of Problem

As noted by Newnham, material science has matured in recent decades from a science to an engineering discipline in which finding the best single-phase materials for a certain application is no longer the focus, but instead is a search for the best combination of materials and determining how to process such composites.1 Researchers have followed this general trend and sought new ways of combining materials to access unique, tailored property sets. By exploiting the properties of multiple materials, researchers can design multi-phase composites optimized for any application.

Difficulties arise, however, in the processing of composite designs. In ex situ fabrication, that is, the macroscopic joining of constituent materials, problems arise in bonding. Interfaces are the weakest part of the overall structure and pose the risk of delamination if not designed precisely, which can be costly. In situ fabrication, also presents potential problems microscopically with thermal expansion mismatch and chemical stability, both of which could potentially introduce residual stresses in the material.2 These processing difficulties and the cost associated with mitigating them have prevented some designs from being implemented.

A new class of ceramic composites has been proposed that mitigates some of the problems associated with multi-phase composite structures. Microstructure composites are materials which combine distinct microstructures of a same composition ceramic to access unique property sets.2 These distinct microstructures are combined in very specific patterns, described by the concept of connectivity proposed by Newnham et al.1 Novel forming methods are used to interconnect two distinct phases in situ into a certain connectivity, a description of how many dimensions each phase is self-connected. In a 1-3 composite structure, for example, one phase is connected to itself in one dimension and the other is connected to itself in three dimensions. In a 2-2 composite structure, both constituents are self-connected in two dimensions, resulting in a laminate-type structure. The ten possible connectivities in a two-phase system are shown in

1

Figure 1.1. By combining constituent components in such specific patterns, the researcher is allowed to design and control the composite properties.

Figure 1.1. The ten possible connectivities in a two-phase composite system. (Picture adapted by M. Fleck, Pennsylvania State University.)1

The concept of microstructure composites has been demonstrated in alumina, in which the constituent microstructures were a fine-grained, equiaxed microstructure and a textured microstructure.2 Slurries of these microstructures were tape cast separately and the resultant tapes stacked to form samples of 2-2 connectivity. These laminate alumina structures demonstrated properties not seen in monolithic alumina, such as macroscopic crack deflection.

The textured alumina microstructure was achieved by templated grain growth, or TGG.3 In this process, template particles are added to the ceramic slurry and aligned in the forming stage of processing. When sintered, grain growth occurs anisotropically in the direction of the templates. This is possible in alumina due to the anisotropic nature of its rhombohedral unit cell, a critical criterion for TGG. In the microstructure composites already mentioned, tape casting was used to align the alumina template particles and liquid-phase sintering was used to

2

encourage anisotropic growth. Dopants added to the slurry encourage anisotropic grain growth by altering the sintering kinetics in the system. The dopants added to the template slurry (SiO2 + CaO) formed the liquid phase for liquid-phase sintering and the dopant added to the non- templated slurry (MgO) prevented abnormal grain growth.2,3,4 To control microstructure development and prevent dopant diffusion between the template and non-templated layers, very low levels of dopants were used.

The mechanical response of the 2-2 connectivity in alumina has been tested and observed at length.2 As seen in Figure 1.2, this material demonstrated macroscopic crack deflection, crack bifurcation, and crack arrest resulting in highly non-catastrophic failure.

Figure 1.2 Fracture behavior in 2-2 alumina microstructure composite bend bars.2

Crack deflection, not seen in monolithic alumina, occurred along the basal plane of the templated particles. It was determined that the anisotropic fracture energy of the textured layer as well as residual compressive stresses developed from thermal expansion mismatch with the equiaxed layer resulted in this crack deflection. These properties improve the toughness of the ceramic.

Fabrication of other connectivities has been demonstrated in alumina and their mechanical response observed, but no data was collected regarding their mechanical properties.2 This dissertation explored the effects of composite connectivity on the mechanical properties of alumina microstructure composites and reports preliminary data about the processing of

3

SiAlON-based microstructure composites. While textured alumina has demonstrated significant 1/2 2 crack deflection and bifurcation, the fracture toughness (KIC, text = 4.5 MPa⋅m ) does not differ 1/2 5 greatly from that of monolithic, equiaxed alumina (KIC, eq = 3.5 MPa⋅m ) . Therefore, it was desirable to study another system in which the microstructure components showed a greater difference in toughness in order to further illuminate the effect of connectivity on mechanical behavior. SiAlON was chosen as the second system because the α- and β-SiAlON phases have significantly different mechanical properties. We believe that through design and controlled processing, these microstructure composites will exhibit the same or enhanced properties of crack deflection, energy dissipation, and non-catastrophic failure as that of the 2-2 connectivity in alumina already studied by Pavlacka.2

The same constituent microstructures as reported by Pavlacka2, i.e. a fine-grained, equiaxed microstructure and a textured microstructure, was used in this study to fabricate alumina composites. The textured microstructure lends itself to crack deflection and energy dissipation due to the anisotropic fracture energy of textured alumina. The component microstructures used in the SiAlON microstructure composites were α- and β-SiAlON, two phases known for their exceptional hardness and toughness, respectively. The β-SiAlON microstructure induces crack deflection and energy dissipation by way of its elongated grain structure. It was hypothesized that these alumina and SiAlON microstructure composites would exhibit unique mechanical behavior by combining the effects of composite connectivity and toughened microstructures.

Tape casting was utilized to fabricate the microstructure composites. Colloidal processing is a highly controllable and reproducible method that produces very homogeneous microstructures. Furthermore, tape casting lends itself easily to fabricating different connectivities, most obviously, the laminate structure. A new method of tape casting called co- casting has been developed that allows more complex connectivities to be fabricated.6 Co- casting allows the simultaneous tape casting of multiple slurries, resulting in a “striped” composite tape, as shown in Figure 1.3.

4

Figure 1.3 Schematic (top view) of co-casting apparatus (a) in which separate slurries are poured into individual reservoirs for simultaneous tape casting. The arrow denotes the direction of the cast. The reservoirs are divided by separator fins (b); the tapered ends of the separator fins allow the slurries to merge just prior to passing under the doctor blade.2

The composite tape is subsequently cut, stacked, and sintered to form the composite in situ. This novel forming method was exploited to fabricate various connectivities in alumina by simultaneously casting non-templated and templated alumina slurries and then rotating and stacking these tapes to obtain alumina microstructure composites with 1-1, 2-2, and 3-3 connectivities; the same was done with α- and β-SiAlON slurries.

1.2 Scientific Approach

The objective of this thesis is to design and fabricate alumina and SiAlON microstructure composites of various connectivities and to determine how the mechanical properties are affected. Thus far, mechanical testing of complex alumina architectures has been limited to the 2-2 connectivity. No work has been done to test the effect of connectivity on SiAlON mechanical behavior. A more comprehensive set of data that describes the mechanical response of more complex connectivities could reveal a unique property set with potential for use in structural applications. This work stems from the earlier work in which the effect of alumina slurry formulations, dopant concentrations, and sintering conditions on texture quality and microstructure control were extensively studied.2 Therefore, the appropriate formulations and

5

conditions were chosen a priori based on that work. The formulations and densification conditions to fabricate α-/β-SiAlON microstructure composites were explored in this work. Some of the issues addressed:

• Fabrication of 1-1, 2-2, and 3-3 alumina microstructures composites including the structure of tapes, rotation, and stacking sequence to obtain the various connectivities • Application of co-casting to the fabrication of various connectivities • Determination of the mechanical response of those composites • Study of the effect of connectivity on fracture behavior of 1-1, 2-2, and 3-3 textured/fine-grained alumina microstructure composites • Preliminary efforts to fabricate 1-1, 2-2, and 3-3 SiAlON microstructure composites and determination of the mechanical response of the same

1.3 Organization of Thesis

Chapter 2 is a literature review of previous work in toughening mechanisms in ceramics and mechanical testing methods. Toughening mechanisms such as microstructural reinforcements and laminate designs are summarized. The materials used in this work, alumina and SiAlON, are introduced, and possible mechanical testing methods that can be used to characterize the microstructure composites fabricated are presented.

Chapter 3 discusses the processing and mechanical behavior of 1-1, 2-2, and 3-3 alumina microstructure composites. Composite architectures, including the design and stacking sequences of co-cast tapes, are presented. Equibiaxial strength and work of fracture of samples of each connectivity is reported. The work of fracture is comparable to previous work on both toughened ceramics and microstructure composites. Mechanical behavior is discussed in terms of crack path, load/deflection behavior, and previous work by other researchers. Step-wise crack deflection along the anisotropic textured grains is observed by SEM and the resultant load/deflection curves show characteristics of non-catastrophic failure.

Chapter 4 discusses the design of α-/β-SiAlON microstructure composites in terms of chemical

6

formula, tape-casting formulation, and densification. Processing difficulties involved in the fabrication of these composites, including failure both in the green state and during densification, are discussed. Initial mechanical testing results are reported.

1.4 References

1. R.E. Newnham, D.P. Skinner, and L.E. Cross, “Connectivity and Piezoelectric- Pyroelectric Composites,” Materials Research Bulletin 13(5), 525-536 (1978).

2. R.J. Pavlacka, "Processing and Mechanical Behavior of Al2O3 Microstructure Composites," Ph. D. Thesis (The Pennsylvania State University, University Park, PA, 2009). 3. M. M. Seabaugh, I. H. Kerscht, G.L. Messing, “Texture Development by Template Grain Growth in Liquid-Phase-Sintered Alpha-Alumina,” Journal of American Ceramic Society 80[5], 1181-88 (1997). 4. R.J. Pavlacka and G.L. Messing, “Processing and Mechanical Response of Highly

Textured Al2O3,” Journal of the European Ceramic Society 30, 2917-25 (2010). 5. R.G. Munro, “Evaluated Material Properties for Sintered α-Alumina,” Journal of the American Ceramic Society 80[8], 1919-28 (1997). 6. E.R. Kupp, G.L. Messing, J.M. Anderson, and V. Gopalan, “Co-Casting And Optical Characteristics of Transparent Segmented Composite Er:YAG Laser Ceramics,” Journal of Material Research 25[3],476-483 (2009).

7

2 Literature Survey

2.1 Toughening Mechanisms

Extensive research has been conducted on processing techniques that increase the fracture toughness of ceramics. Researchers have followed two approaches to this end: “flaw elimination” processing and “flaw tolerant” processing.1 Flaw elimination processing is motivated by the idea that fracture is initiated and promulgated by intrinsic defects, like microcracks and pores, and that eliminating these flaws makes the material more resistant to fracture and decreases the strength variability.1,2 Conversely, the flaw tolerant approach aims at designing and creating materials that are tolerant of crack propagation and therefore, absorb more energy prior to fracture. The flaw tolerant approach usually involves tailoring the microstructure or overall architecture to promote controlled crack growth (through deflection and bifurcation) and crack arrest, as well as delamination in laminate structures. Cracks will deflect, for example, along weak interfaces or in areas of residual stress.3 The flaw tolerant approach employs toughening mechanisms that lead to “graceful failure;” that is, the failure mode in the material is pre-determined, and the material’s capability to absorb energy is greatly enhanced. This thesis focuses on the flaw tolerant approach by controlling crack propagation through microstructural tailoring.

Crack bridging is one mechanism by which researchers have attempted to toughen ceramics. Fibers, whiskers, platelets, and large particles added to the ceramic matrix during processing act as microstructural reinforcements that disrupt crack propagation and increase the fracture toughness of the matrix.4 Alumina reinforced with silicon carbide whiskers, for example, has exhibited significant increases in fracture toughness; the use of whiskers, however, has since been avoided due to health risks involved.5,6 Khan et al. used alumina agglomerates in an alumina matrix to induce crack bridging; however, only a moderate amount of toughening was measured. Microstructures of these particulate-reinforced alumina materials are shown below in Figure 2.1.7 Using the nomenclature in this thesis, the alumina microstructure by Khan would be 0-3.

8

Figure 2.1. Particulate (SiC whiskers, left, and alumina agglomerates, right) reinforcement in alumina.6,7

While microstructural reinforcements have shown moderate toughening in ceramics, they usually have an adverse effect on the strength of the ceramic. Toughening by elongated particles like whiskers and platelets is based on crack deflection at the interfaces of these particles and the matrix; however, the same particles could act as flaws and weaken the material.8 By employing these techniques in layers rather than in particulates, the ceramic can be toughened without any compromise in strength. For this reason, many researchers have explored the effect of laminate designs on crack propagation. Among these designs is an alternating strong layer/weak layer ceramic. Clegg et al. coated green silicon carbide layers with graphite, which acted as a weak interlayer.9 Crack deflection was observed at the weak interfaces when tested in bending. When one layer failed, the crack would “pop-in” to the next layer and deflect again, a phenomenon that greatly increased its resistance to fracture. Compared to monolithic SiC, toughness was increased four-fold and work of fracture was increased hundred-fold. Figure 2.2 shows the SiC/graphite laminate structure and load/deflection behavior.

9

Figure 2.2. Crack deflection in laminate SiC with weak graphite interlayers and load/deflection curve showing multiple “pop-in” features characteristic of graceful failure.9

This curve is characteristic of “graceful failure,” in that the first drop-off in load does not represent failure of the specimen, but rather failure and delamination of the first layer.10 The specimen supports further loading by subsequent layers due to crack deflection at each weak interlayer, increasing the overall work of fracture of the specimen. Morgan and Marshall applied this idea to alumina, which was dip-coated in monazite (LaPO4) to create weak interlayers. Figure 2.3 is an example of the crack arrest that is characteristic of these weak layers.11

Figure 2.3. Crack arrest by weak monazite interface layer in alumina matrix.11

Another laminate design used to toughen ceramics involves deliberately building porous interlayers, containing 35-40% porosity, into a material. Davis et al. demonstrated exaggerated crack growth along porous interlayers in alumina in Figure 2.4.12 Alumina slurries with and without starch were tape cast to produce dense and porous layers, respectively, which were

10

alternately stacked to form the laminate. Their load/deflection plots exhibit the same “pop-in” characteristic mentioned above, demonstrating that porous layers can indeed toughen ceramics through extended crack deflection and increased work of fracture.

Figure 2.4. Crack deflection in laminate alumina with varying amounts of porosity in porous alumina interlayers.12

While crack deflection can occur at weak interfaces, it can also occur at strong interfaces where residual stresses are the driving force for deflection.10 Residual stresses arise from differences in thermal expansion coefficients between layered dissimilar materials, uneven sintering rates, and martensitic phase transformations that result in volume changes. All of these mechanisms can be used to increase fracture resistance and decrease strength variability in a material if they are controlled during processing.13 Green et al. demonstrated that residual stresses processed into silicate glass by ion exchange caused crack deflection, which 14 subsequently increased strength. This idea was applied to alumina/mullite (3Al2O3·2SiO2) laminates by Rao et al., where the source of residual stress was thermal expansion mismatch. These specimens exhibited a “threshold strength” characteristic, that is, a strength below which the system would not fail despite the existence of intrinsic flaws in the ceramic.15

In some cases, the incorporation of weak interfaces, porosity, or residual stresses into a material may not be desired. Texturing is one method that can disrupt crack propagation while avoiding the adverse effects of those toughening mechanisms. Texturing is achieved using the templated grain growth (TGG) process, whereby templates are aligned in the forming process and grow into anisotropic grains.16 Carisey et al. demonstrated that the low-energy basal face of anisotropic alumina grains can inhibit crack propagation in monolithic textured alumina.17

11

Pavlacka created a laminate (2-2 connectivity) of textured alumina layers alternated with equiaxed alumina layers which exhibited exaggerated crack deflection and bifurcation when tested in bending; significant increases in work of fracture as compared to monolithic textured alumina samples was also observed.18 A fractured specimen that demonstrates crack deflection, bifurcation, and arrests along textured layers is shown in Figure 2.5. It is interesting to note that Pavlacka observed no crack deflection in monolith samples of textured alumina; he concluded that co-sintering of the equiaxed and textured materials in a composite architecture fundamentally changes the fracture behavior of the material.

Figure 2.5. Textured/equiaxed alumina microstructure laminate fracture specimen showing crack deflection, bifurcation, and arrest.18

Layered structures offer an easy way to tailor microstructural and macrostructural properties to various applications. They are often simple to fabricate and offer a high degree of microstructural control, especially using colloidal processing. It is easy to imagine more complex structures that can be fabricated using any one of the toughening mechanisms described above. Apart from the laminate design, structures may also be designed so that two phases, such as the textured and equiaxed alumina microstructures mentioned above, may be interconnected in any number and combination of dimensions. Newnham et al. presented all the configurations possible for a two-phase system to be interconnected, including the laminate structure.19 Pavlacka demonstrated the fabrication of several connectivities in addition to laminate, or 2-2 connectivity, design with textured and equiaxed alumina microstructures.18 Specimens of the 1-3, 0-3, and 3-3 connectivities were achieved using novel green forming methods such as screen- printing and co-casting.

12

2.2 Materials

Two material systems were chosen for this study. Both systems can be tailored to have multiple, distinct microstructures, which could be fabricated together in situ to create microstructural composites for mechanical testing. The first is alumina, which can be processed both in a fine-grained, equiaxed microstructure and a textured microstructure. As mentioned above, textured alumina has shown toughening characteristics. Furthermore, a wealth of scientific literature is readily available on alumina, which can be referenced in order to tailor the processing of samples. Being such an extensively studied material, the effects of dopants and sintering trajectories on microstructure development, especially textured grain growth (which will be discussed later), is already known. Also, property sets for different microstructures are available for comparison. The availability of fine alumina powders makes it an ideal research material. The general properties of alumina are well known and will not be discussed here.

The second material selected for study is SiAlON. Extensive research has been done on this class of materials, due to its superior mechanical properties and relatively low density (3.0- 3.2 g/cm3).20 Furthermore, there exists a wide range of compositions and grain morphologies in the SiAlON family, allowing for its properties to be tailored for specific applications, in particular, the α-SiAlON phase, important for its high hardness, and the β-SiAlON phase, important for its high toughness due to an anisotropic, elongated grain structure. These two 21 phases are isostructural with and share many of the same properties as α-Si3N4 and β-Si3N4. They will serve as the distinct microstructures of SiAlON that will be fabricated into microstructural composites.

α-SiAlON is created by partially substituting both Al-N and Al-O bonds for the Si-N bonds in α-Si3N4. The resultant valence imbalance is compensated by the addition of a metal cation, yielding the chemical formula MxSi12-(m-n)Al(m+n)OnN(16-n). The parameters m and n describe the number of Si-N bonds replaced by Al-N and Al-O bonds, respectively. The parameter x is related to the valence of the metal cation, M; commonly used metal cations

13

include Yb, Y, or Nd.21 β-SiAlON is achieved by the partial substitution of Al-O bonds for Si-N 22 bonds, described by the chemical formula Si6-zAzOzN8-z, where 0

There is a wide range of compositions for both phases, and their microstructure and properties can therefore be tailored according to the composition of the starting powders. Generally, however, there is a distinct difference in the microstructure of the two phases. α- SiAlON typically exhibits an equiaxed microstructure while β-SiAlON exhibits more anisotropic, elongated grain morphology. α-SiAlON is known for its high hardness, 18-21 GPa, akin to its parent material, α-Si3N4. β-SiAlON is not as hard (~16 GPa), but is known for its high fracture toughness, 4-7 MPa·m1/2, due to the presence of elongated grains, which act as crack bridges for propagating cracks.22,23 The elongated grain structure of the β phase is best achieved by the aid of a liquid phase (usually formed with Y2O3) during densification; the resultant residual glassy phase acts as an additional toughening mechanism as it promotes crack deflection in the final material.24

The two phases of SiAlON are completely compatible and easily co-processed using the same colloidal approach utilized for the alumina composites, allowing for a wide range of possibilities for the fabrication of composite materials by tape casting.23 Research has shown 25,26,27 improvement in the toughness of α-SiAlON reinforced with both elongated β-Si3N4 and β- SiAlON grains.23 In some of these cases, though not all, the composite showed diminished hardness as compared to single phase α-SiAlON.27 In all cases, the reinforcement was seen in the bulk; no research has been done in which the mechanical properties of α- and β-SiAlON were combined in a macro-composite, such as a laminate or other composite architecture. The present work aims to explore the mechanical properties of an α/β-SiAlON composite, as well an equiaxed/templated alumina composite, in which the separate phases are interconnected according to the concept of connectivity.

2.3 Mechanical Testing

14

When testing the mechanical effects of connectivity on a material, an appropriate method must be chosen which tests specimens that are representative of the connectivity, which is macrostructural by nature. This can be difficult due to the large sample sizes required. Bend bars, for example, are too thin to represent most connectivities that are processed on the millimeter scale (with the notable exception of the 2-2 connectivity).18 Any method that requires the use of bend bars is therefore inappropriate. The single-edge notched beam (SENB) test is the most common method for testing fracture toughness in ceramics. This method, however, requires the use of bend bars for testing in three- or four-point bending, and therefore would not be appropriate for this study. Interfacial fracture resistance, another useful parameter that examines the toughness of laminates, also utilizes bending tests.18,28

Biaxial flexure testing is a method that is used to test samples of larger volume. Biaxial flexure tests thin-plate samples; Ritter et al. determined that their shape, usually circular or square, does not affect the results.29 Biaxial flexure is performed in a number ways; the three most common methods are ring-on-ring (ROR), piston-on-3 ball, and ball-on-ring. It has been determined that the ring-on-ring method gives the most accurate measure of biaxial flexure strength, as this method places the most volume of the specimen under the maximum stress (the other methods only place stress at the points where the balls touch the specimens).29 While biaxial flexure testing can be used to measure fracture toughness, the procedure is not well documented, nor does it appear in the ASTM standard. The test, however, does yield both flexural strength values and load/deflection curves, which may be used to determine work of fracture.

2.4 References

1. R. Bermejo Moratinos, "Structural Integrity of Alumina-Zirconia Multilayered Ceramics," Ph. D. Thesis (Universitat Politècnica de Catalunya, Barcelona, 2006). 2. A.G. Evans, “Perspective on the Development of High-Toughness Ceramics,” Journal of the American Ceramic Society 73[2], 187-206 (1990). 3. M. Wei, D. Zhi, and D. Brandon, “Oxide Ceramic laminates with Highly Textured α- Alumina Interlayers: I. Texture Control and Laminate Formation,” Journal of Material Science 41:7425-36 (2006).

15

4. D.J. Green, An Introduction to the Mechanical Properties of Ceramics (Cambridge University Press, 1998). 5. M. Farkash and D.G. Brandon, “Whisker Alignment by Slip Extrusion,” Materials Science and Engineering A177, 269-275 (1994).

6. T.N. Tiegs and P.F. Becher, “Sintered Al2O3-SiC-Whisker Composites,” American Ceramic Society Bulletin 66[2], 339-342 (1987). 7. A . Khan, H.M. Chan, and M.P. Harmer, “Toughening of an Alumina-Mullite Composite by Unbroken Bridging Elements,” Journal of the American Ceramic Society 83[4], 833- 840 (2000). 8. P.F. Becher, “Microstructural Design of Toughened Ceramics,” Journal of the American Ceramic Society 74[2], 255-269 (1991). 9. W.J. Clegg, K. Kendall, N. McN. Alford, T.W. Button, and J.D. Birchall, “A Simple Way to Make Tough Ceramics,” Nature 347, 455-457 (1990). 10. H.M. Chan, “Layered Ceramics: Processing and Mechanical Behavior,” Annual Review of Materials Science 27, 249-282 (1997). 11. P.E. Morgan and D.B. Marshall, “Ceramic Composites of Monazite and Alumina,” Journal of the American Ceramic Society 78[6], 1553-63 (1995). 12. J.B. Davis, A. Kristoffersson, E. Carlstrom, and W.J. Clegg, “Fabrication and Crack Deflection in Ceramic Laminates with Porous Interlayers,” Journal of the American Ceramic Society 83[10], 2369-74 (2000). 13. V.M. Sglavo, M. Paternoster, and M. Bertoldi, “Tailored Residual Stresses in High Reliability Alumina-Mullite Ceramic Laminates,” Journal of the American Ceramic Society 88[10], 2826-32 (2005). 14. D.J. Green, R. Tandon, and V.M. Sglavo, “Crack Arrest and Multiple Cracking in Glass Through the Use of Designed Residual Stress Profiles,” Science 283, 1295-97 (1999). 15. M.P. Rao, A.J. Sánchez-Herencia, G.E. Beltz, R.M. McMeeking, and F.F. Lange, “Laminar Ceramics That Exhibit a Threshold Strength,” Science 286, 102-105 (1999). 16. M.M. Seabaugh, I.H. Kerscht, and G.L. Messing, “Texture Development by Template Grain Growth in Liquid-phase-sintered Alpha-Alumina,” Journal of American Ceramic Society 80[5], 1181-88 (1997).

16

17. T. Carisey, I. Levin, and D.G. Brandon, “Microstructure and Mechanical Properties of

Textured Al2O3,” Journal of the European Ceramic Society 15, 283-289 (1995).

18. R.J. Pavlacka, "Processing and Mechanical Behavior of Al2O3 Microstructure Composites," Ph. D. Thesis (The Pennsylvania State University, University Park, PA, 2009). 19. R.E. Newnham, D.P. Skinner, and L.E. Cross, “Connectivity and Piezoelectric- Pyroelectric Composites,” Materials Research Bulletin 13(5), 525-536 (1978). 20. T. Ekström, P.O. Käll, M. Nygren, P.O. Olsson, “Dense Single-Phase β-SiAlON Ceramics by Glass-Encapsulated Hot Isostatic Pressing,” Journal of Materials Science 24, 1853-61 (1989). 21. P. Sajgalik, Z. Lences, and M. Hnatko, “Nitrides,” in Ceramics Science and Technology, Volume 2, edited by R. Riedel and I-W. Chen, 59-89 (2010). 22. G. Z. Cao and R. Metselaar, “α’-SiAlON Ceramics: A Review,” Chemistry of Materials 3, 242-252 (1991). 23. M.I. Jones, H. Hyuga, and K. Hirao, “Optical and Mechanical Properties of α/β Composite SiAlONs,” Journal of the American Ceramic Society 86[3], 520-22 (2003). 24. T. Ekström and P-O Olsson, “β-SiAlON Ceramics Prepared at 1700˚C by Hot Isostatic Pressing,” Journal of the American Ceramic Society 72[9], 1722-24 (1989).

25. I-W Chen and A. Rosenflanz, “A Tough SiAlON Ceramic Based on α-Si3N4 With a Whisker-Like Microstructure,” Nature 389, 701-704 (1997).

26. S-L Hwang, H-T Lin, P.F. Becher, “Mechanical Properties of Β-Si3N4 Whisker Reinforced α’-SiAlON Ceramics,” Journal of Materials Science 30, 6023-27 (1995).

27. N. C. Acikbas, R. Kumar, F. Kara, H. Mandal, B. Basu, “Influence of β-Si3N4 Particle Size and Heat Treatment on Microstructural Evolution of α:β-SiAlON Ceramics,” Journal of the European Ceramic Society 31, 629-635 (2011). 28. P. Charalambides, J. Lund, A. Evans, R. McMeeking, “A Test Specimen for Determining the Fracture Resistance of Bimaterial Interfaces,” Journal of Applied Mechanics 56[1],77-82 (1989). 29. J. E. Ritter Jr., K. Jakus, A. Batakis, and N. Bandyopadhyay, “Appraisal of Biaxial Strength Testing,” Journal of Non-Crystalline Solids 38-39, 419-424 (1980).

17

3 Fabrication and Mechanical Properties of 1-1, 2-2, and 3-3 Alumina Microstructure Composites

3.1 Introduction

The anisotropic morphology of textured grains has been shown to toughen ceramic materials by deflecting and arresting cracks. To date, this behavior has only been demonstrated in materials reinforced with individual anisotropic grains in laminate materials with textured layers.1,2,3,4 By applying the concept of connectivity, one can test the mechanical properties of unique ceramic architectures in which channels of textured material, systematically self- connected, deflect and arrest propagating cracks. Preliminary work on this concept has been done in alumina; these materials were termed “microstructure composites,” due to the composite nature of distinct microstructures of a single phase of material.5 The following work describes the design, fabrication, and mechanical testing of 1-1, 2-2, and 3-3 microstructure composites of alumina.

Achieving the level of microstructural control to create connectivities in structural ceramics requires the use of a colloidal tape-casting process and, more specifically, the co- casting process. Tape casting allows the fabrication of laminate materials by lay up of alternating, distinct ceramic tapes. Co-casting, in which two different ceramic slurries are simultaneously cast, allows the fabrication of more complicated architectures by the lay-up of ceramic composite tapes. Examples of stacking sequences for both are shown in Figure 3.1.

Figure 3.1. Stacking configuration for 2-2 (tape-cast tapes) and 1-3 (co-cast tapes) connectivities.5

18

The composite tape designs and stacking sequences of the 1-1, 2-2, and 3-3 connectivities fabricated in this work are shown in Figure 3.2.

Figure 3.2. Connectivities, their respective tape stacking configurations, and sample schematics for the 1-1, 2-2, and 3-3 connectivities, respectively. The dark stripes represent the textured material and the light stripes represent the equiaxed material. The direction of casting, and therefore the orientation of the template particles, is parallel to the stripes.

To create the 3-3 connectivity, for example, composite tapes were cast with alternating equiaxed and templated “stripes,” which were cut into 5 cm square pieces. These pieces were then stacked, rotating each 90°. This produced a part in which the both the equiaxed and templated microstructures were self-connected in all three dimensions. Rotation of each tape also prevented warpage of the specimens that might be caused by residual stresses in the tape from casting.

19

A distinction must be made between the 2-2 composites fabricated by Pavlacka and the 2- 2 composites fabricated in this work. Pavlacka fabricated what will be referred to as “2-2 (series),” which is a true laminate made of alternating equiaxed and templated layers. The 2-2 composites fabricated in this work will be referred to as “2-2 (parallel),” in which subsequent tapes with equiaxed and templated stripes are stacked. The tape stacking sequences for both are shown in Figure 3.3.

Figure 3.3. Tape stacking sequences for the 2-2 series (a) and 2-2 parallel (b) connectivities. The nomenclature refers to the manner is which the microstructural components see the load. In the series, the equiaxed and textured layers see the load individually while in the parallel, the equiaxed and textured layers see the load simultaneously (in the same layer).

The “series” and “parallel” nomenclature refers to the manner in which the microstructural components see the load, which is always normal to the textured material. In the 2-2 (series) samples, the equiaxed and textured layers see the loading separately, while in the 2-2 (parallel) samples, the equiaxed and textured stripes see the loading together in the same layer. This nomenclature will be used for the remainder of this work.

3.2 Experimental Procedure

3.2.1 Sample Fabrication

The starting powders were high-purity α-Al2O3 (AKP-50, Sumitomo Chemical CO, Ltd., Tokyo, Japan) with a particle size of 0.1-0.3 µm and single crystal alumina platelets (Advance Nanotechnology Limited, Welshpool, Australia) with a thickness of ~100 nm and a diameter of 5-10 μm, both shown in Figure 3.4.

20

Figure 3.4. Micrographs of as-received AKP-50 (a) and alumina platelets (b).

Magnesium nitrate, calcium nitrate, and tetraethylorthosilicate (TEOS) were used as dopant precursors. The dispersant and solvent systems used were Menhaden fish oil (MFO) and 50:50 (wt ratio) ethanol/xylenes, respectively. The alumina powder and appropriate precursors (magnesium nitrate for the equiaxed slurries and calcium nitrate/TEOS for the templated slurries) were dispersed in the dispersant/solvent solution and ball-milled for 24 h using high-purity alumina milling media. The precursors were precipitated with several drops of dilute NH4OH prior to milling. Binder and plasticizers (polyvinyl butyral, polyalkylene glycol, and benzyl butyl phthalate) were then added, followed by an additional 24 hours of milling. For the templated alumina slurries, the alumina templates (5 vol %) were ultrasonicated in the appropriate amount of the same dispersant/solvent mixture. The template suspension was then added to the slurry before the last 30 min of milling to prevent breakage during mixing. The milling media were then sieved from the slurries, which were subsequently de-aired while stirring. Equiaxed slurries were also poured through a 55 µm mesh screen to remove agglomerated powders. The slurries were stirred for ~4-5 h prior to tape casting.

The equiaxed and templated slurries were co-cast using a multi-reservoir doctor blade fixture at a rate of ~70 cm/min with a gap height of 0.25-0.36 mm onto silicon-coated Mylar (G10JRM; The Tape Casting Warehouse, Morrisville, PA) using a table top caster (TTC-1200; The Tape Casting Warehouse, Morrisville, PA). The resultant tapes were composite alumina tapes with equiaxed and templated “stripes.” The full tape was ~5 cm wide and the stripes were ~1 cm wide. Dry tape thicknesses ranged from 220 µm to 320 µm. These tapes were then cut

21

and stacked (~30-40 tapes per specimen) according to the desired connectivity into ~5x5x0.6 cm green samples. Alignment of the different microstructural compositions was achieved by careful measurement and cutting of the tapes.

The samples were isostatically laminated at 75°C and 20 MPa for 40 min. The organics were removed using a burnout cycle in air at 600°C for 12 h with a heating rate of 0.2- 0.3°C/min. All samples were pressureless sintered in air at 1550°C for 90 min with a heating rate of 5°C/min. Cross-sections were polished to a 1µm finish using appropriate diamond slurries and thermally etched in air at 1450°C for 30 min to reveal the grain boundaries. The fabrication process detailed above is summarized in the flow chart in Figure 3.5.

22

Figure 3.5. Flow chart for processing templated and equiaxed alumina microstructure composites.

23

All samples were fabricated and machined to meet the overhang and thickness requirements defined in the ASTM Standard for Equibiaxial Flexural Strength of Advanced Ceramics.6 All were ground flat using a surface grinder to ensure the load and support rings lay flat across the entire surface of the specimens; flatness was confirmed by measuring the thickness of various areas across the plate with a micrometer. Final sample dimensions were 40±1 mm x 40±1 mm square plates of thickness 2.5±0.5 mm. The tensile sides were beveled and polished to a 1 µm finish.

3.2.2 Mechanical Testing

The ring-on-ring equibiaxial flexure test configuration shown in Figure 3.6 was used.

Figure 3.6. Schematic of ring-on-ring fixture for equibiaxial strength testing. Adapted.7

The ring-on-ring fixture consists of a support ring of diameter (Ds) 28 mm and a concentric loading ring of diameter (Dl) 20 mm. An Instron (Model 5866, Instron Corp., Canton, MA) load frame was used to drive the loading ram in testing, applying the load to the samples at a constant displacement rate of 0.1 mm/min. The load was oriented normal to the direction of casting, and thus, normal to the basal face of the oriented grains in the textured regions. GrafoilTM (UCAR International, Inc., Nashville, TN) sheets (0.26 mm thick) and manila paper acted as compliant

24

layers, reducing the friction between the load/support rings and the test specimen which could lead to crack initiation at stresses other than the equibiaxial strength. Adhesive tape was also applied to the compressive face of the specimen to retain the fracture fragments for fractographic analysis. Steps were taken to ensure consistent alignment of the loading and support rings with the specimens for each test. The stiffness of the test apparatus was measured by running the test with the only the compliant layers (i.e. no specimen). Fracture fragments were cut for post- failure characterization and analysis. Cut samples were etched 100°C below the sintering temperature to reveal grain boundaries. These and fracture surfaces were gold-coated prior to examination in a scanning electron microscope (Philips XL20, Philips Electronic Instruments Co., Mahweh, NJ.)

Equibiaxial strength (MPa) was calculated using the following equation:

(1)

where F is the breaking load in Newtons, h is the test specimen thickness, υ is the Poisson’s ratio of the material (taken as 0.23 for Al2O3), and D is the characteristic diameter of a circle that expresses the characteristic size of a plate, defined as

(2)

where l =0.5 (l1 +l2), an average of the x and y lengths of the specimen, which were not always perfectly square. Work of fracture (J/m2) was calculated as the area under the load-deflection curve divided by twice the cross-sectional area of the test specimen.2,5 Percent theoretical density, based on a theoretical density of 3.986 g/cm3, was determined by the Archimedes method.

3.3 Results and Discussion

25

3.3.1 Co-casting Observations

To create microstructure composites of accurate connectivities and microstructural control, care must be taken to ensure the casting of tapes of both uniform thickness and distinct component features, or “stripes”. Final tape thickness is affected by four factors: the casting rate, the blade gap (or height of the doctor blade), reservoir height, and slurry viscosity.5 Previous co- casting work showed that the slurry viscosity was the most important factor, given that the casting rate and blade gap were constant across the composite tape, and reservoir height was never more than 1 cm when casting by hand, which would have a negligible effect on final thickness.5 Mismatched slurry viscosities result in tapes with differential thickness, which could induce cracking upon drying at the interface of the microstructural components (the stripes), rendering the tape useless. Paying careful attention to slurry viscosities and evaporating solvent when needed can prevent this.

It is also important to have distinct component features, or straight stripes, in the composite tapes to ensure accurate connectivity in the final part. This is achieved by careful assembly of the multi-reservoir co-casting apparatus and prevention of slurry mixing in the apparatus. A closer look of the co-casting apparatus is shown in Figure 3.7.

26

Figure 3.7. A closer look of the co-casting apparatus (top view) in which separate slurries are poured into individual reservoirs for simultaneous tape casting (a). Separator fins are used to divide the slurry reservoirs. Tapered edges (b) at the end of the separator fins allow for the component slurries to merge just before passing under the doctor blade (c).5

If the separator fins are not parallel and flush with the tabletop, slurry mixing can occur under the fins within the reservoir before passing under the doctor blade. Furthermore, slurry mixing can occur if the reservoir height is too high, inducing spillage over the tapered component fins. Figure 3.8 shows the tape design, stacking sequence, and final component schematic of a 3-3 alumina microstructure composite.

Figure 3.8. Composite tape design, stacking sequence, and schematic of final textured/equiaxed

Al2O3 3-3 microstructure composite. The schematic represents the “ideal” tape, with straight component stripes, no bleeding of slurries into adjacent components, and uniform thickness across the tape; when these factors are achieved, one can obtain a specimen of accurate connectivity.

The figure shows an ideal tape, in which the component stripes are straight and have not bled into adjacent slurries. Careful measurement, cutting, and subsequent stacking produce a specimen of accurate connectivity with distinct microstructural components.

3.3.2 Microstructural Control

27

Microstructural control between the equiaxed and textured layers was achieved in the microstructure composites. Figure 3.9a shows the spatial control between the layers of a 3-3 microstructure composite and the relative thicknesses of the textured and equiaxed layers, 208 µm and 138 µm, respectively. These thicknesses are typical of all specimens and are controlled by original tape thicknesses, which ranged from 220 µm to 320 µm. Original tape thicknesses were kept in this range to ensure full textured grain growth in the textured layer; previous work shows that textured layers less than 200 µm have regions of equiaxed grains due to dopant diffusion in thin layers.5 Figure 3.9b shows the interface (denoted by the dotted line) between the fine-grained, equiaxed microstructure and the textured microstructure.

Figure 3.9. Micrographs demonstrating microstructural control in 3-3 fine-grained/textured

Al2O3 microstructure composites. Relative layer thicknesses (related to original tape thicknesses) are shown in (a); these thicknesses are typical of all composites made. Micrograph (b) shows a closer look of the clear interface (marked by the dashed line) between the distinct equiaxed and textured layers.

Polished surfaces as well as fracture surfaces of the equiaxed and textured microstructures in the 3-3 fine-grained/textured microstructure composite are shown in Figure 3.10. Again, the fine- grained (G.S. = 1.5-2.5 µm), equiaxed microstructure was achieved with the addition of 1000 ppm MgO dopant and the textured microstructure (G.S. = ~2.0 µm thick, ~16 µm long) was achieved with the addition of 5 vol% α-Al2O3 template particles and 0.14 wt% CaO + SiO2 dopants. Sample theoretical densities were in the range of 94 – 98%, assuming a theoretical 3 density of Al2O3 of 3.986 g/cm .

28

Figure 3.10. Polished surfaces (a and b) and fracture surfaces (c and d) of the fine-grained, equiaxed and textured microstructures in a 3-3 microstructure composite, respectively. Microstructural control was achieved through doping (MgO for equiaxed components and

SiO2+CaO for textured components) of the ceramic slurries.

3.3.3 Work of Fracture Calculation

Work of fracture is measured by determining the area under the load/deflection curve and dividing by twice the cross-sectional area of the specimen.2,5 In this work, the area under the curve was determined from the raw data using the trapezoidal rule. A typical load/deflection curve is shown in Figure 3.11. The curves exhibit non-linear loading, followed by a linear rise, an initial drop in load, and successive reloading of the sample. The multiple reloading indicates multiple crack arrests before failure.

29

Figure 3.11. Typical load/deflection curve.

Non-linear loading is common in mechanical testing due to the alignment and stiffening of the compliant layers and the sample as well stiffening of the test machine. This presents some difficulty, however, in quantitative fracture analysis as the curves represent both the alignment of the fixture as well as fracture of the specimen. In order to accurately measure the work of fracture of the specimens tested in this work, the alignment/stiffening portion of the load/displacement curves must be separated from the fracture energy associated with the specimen as the linear portion of the loading curve represents both the elastic displacement of the machine and of the specimen. The stiffness of the machine can be measured by testing only the compliant layers, that is, running a test without a specimen (see Fig 3.6). This test was performed and the machine stiffness was found to be ~100 N/mm.

A more accurate quantitative analysis of the work of fracture could be achieved by making corrections for the non-linearity in the graphs, as these do not accurately represent the energy for fracture. Structurally improving the mechanical testing apparatus as well as accounting for the stiffness of the machine could be used to provide a more accurate data analysis of the load/deflection curves. These corrections could be validated by evaluating the expected specimen deflection values from the corrected curves. The resultant displacements of the specimens shown on the corrected loading curves should be consistent with the plate deflection equation listed in the ASTM standard6:

30

3F 1− v 2 D2 ⎛ D2 ⎡ (1 v)(D2 D2 ⎤ ⎛ D ⎞⎞ ( ) L S − S − L S (3) δ = 3 ⎜ 2 ⎢1 + 2 ⎥ −⎜1 + ln ⎟⎟ 8πEh ⎝ DL ⎣ 2(1+ v)D ⎦ ⎝ DL ⎠⎠ where δ is the deflection and other parameters defined previously. While these corrections were beyond the scope€ of this work, the work of fracture results are still presented to give the reader an idea of the relative fracture behavior of the various connectivities tested. Another consideration is the kinetic effect of the testing machine on the fracture specimen. The potential energy stored in the testing apparatus manifests itself into kinetic energy during fracture, leading to a misrepresentation of the work of fracture. The reader is therefore cautioned that the work of fracture results presented are not corrected and can be assumed to be overestimated due to both the non-linearity and the kinetic effect; these results will be referred to as “WOF.”

3.3.4 Fracture Behavior and Work of Fracture

For this discussion, the “transverse” direction will refer to the direction normal to the templated material (the direction of the loading) and the “longitudinal” direction will refer to the direction of the textured layer (the direction of casting).

2-2 Parallel Connectivity

A schematic of the 2-2 parallel microstructure composite, including a closer look at the microstructural components, is shown in Figure 3.12.

31

Figure 3.12. Schematic of the 2-2 parallel microstructure composite. The inset also shows the textured (left) and equiaxed (right) microstructural components.8

To date, the mechanical behavior of only the 2-2 series alumina microstructure composite has been evaluated. The microstructure composites presented in this work (2-2 parallel, 1-1, and 3- 3) are identical in chemistry and processing parameters as the 2-2 series; they differ only in connectivity and method of testing. While Pavlacka studied a number of different template loadings, he showed that 5 vol% templates yielded the greatest “work of fracture” in the 2-2 series alumina microstructure composites; therefore, this was the template loading used in this work.5 The 2-2 series samples were tested in bending, as their composite architecture (laminate) is readily represented in bend bars.5 The 2-2 parallel composites fabricated and tested in this work were tested in biaxial flexure. While crack deflection, bifurcation, and arrest are easily observed in bend bars in which the stress is applied in one direction, these fracture features are somewhat more difficult to observe in biaxial flexure specimens in which the stresses are applied in two directions. Despite this, comparisons can still be made between these microstructure composites.

As previously mentioned, Pavlacka observed extended crack deflection, step-wise fracture, and bifurcation in the 2-2 series alumina microstructure composite. Crack deflection in the 2-2 parallel microstructure composite is limited to ~100-200 µm, whereas Pavlacka observed

32

deflections up to 1 mm in the 2-2 series. (Deflections of up to 1 mm were observed in the 1-1 and 3-3 connectivities discussed later). Extensive step-wise fracture, however, is seen readily in the 2-2 parallel as seen in Figure 3.13.

Figure 3.13. Step-wise fracture in a 2-2 parallel alumina microstructure composite. The cracks deflect ~100-200 µm in the textured layers before intersecting the equiaxed layers.

Figure 3.14 shows a closer look of this step-wise fracture and demonstrates the nature of crack propagation through the individual layers. Cracks propagate transversely through the equiaxed layers; the nature of crack propagation is intergranular. When the crack reaches the textured layers, it begins to deflect, or propagate in the longitudinal direction, along the basal face of the textured grains. The crack propagates intergranularly along the interface between the textured grains, down the edge, and to the next face.

33

Figure 3.14. The nature of crack propagation differs between the equiaxed and textured layers. In the equiaxed layers, the propagation is transverse (direction indicated by the arrow). In the textured layers, the propagation is longitudinal, along the basal faces of the textured grains.

Pavlacka suggested that the reason for the intergranular propagation in the textured layer is that the textured microstructure exhibits low fracture energy along the basal face, and therefore, the long interface between the textured grains creates a weak path for deflection. (Pavlacka confirmed that the fracture resistance is lower in the texture plane than in the through-thickness direction of templated material).

It is interesting to the note where deflection begins within the textured layer. Figure 3.14 shows crack deflection beginning at the onset of the textured layer; that is, as soon as the crack reaches the textured layer, it begins to deflect and propagate longitudinally. This behavior is typical of laminate materials toughened by weak interfaces, such as the SiC/graphite materials studied by Clegg.2 However, as Pavlacka noted, the deflection morphology observed in equiaxed/textured composites is slightly different. Deflection in materials with weak interfaces occurs at the bottom of the weak layer; that is, the crack kinks through the entire weak layer, and deflects at the interface with the next strong layer. This is not observed in this work (or in Pavlacka’s 2-2 series composites). Rather, deflection begins at the top interface and continues

34

steadily through the entire textured region before kinking into the equiaxed region. This suggests that textured material does not act as a weak interface. Pavlacka suggested that crack deflection may be induced by residual stresses developed during densification as a result of thermal expansion mismatch between textured and equiaxed alumina. This behavior, however, is not always observed in equiaxed/textured alumina microstructure composites. In the 2-2 series composite, Pavlacka also observed initial transverse crack propagation through the textured layer before the crack “kinked” out along the basal plane and deflected. This is seen in the 1-1 and 3-3 composites and will be discussed further in the next section.

Crack bifurcation, seen in the 2-2 series composites, was not observed in the 2-2 parallel (or any of the composites fabricated in this work) connectivity. Crack bifurcation is most readily observed, however, in bend bars, and may have been missed in the specimens tested in biaxial flexure.

While the crack deflection behavior observed was similar to that seen in the 2-2 series composites, the nature of failure sometimes differed. Non-catastrophic failure is observed in the weak-interface materials (SiC/graphite) fabricated by Clegg, as well as in the 2-2 series microstructure composites (5 vol% template loading) fabricated by Pavlacka. Both catastrophic and non-catastrophic failures were seen in the 2-2 parallel microstructure composites. Unexpectedly, specimens with extensive crack deflection and step-wise fracture did not show the “pop-in” features in their load/deflection graphs that indicate graceful failure. Figure 3.15a, for example, shows a 2-2 parallel specimen exhibiting extensive step-wise fracture. The load/deflection graph for this specimen, shown in Figure 3.15b, is characteristic of catastrophic failure, as it shows no “pop-in” features. Despite this, however, this specimen shows the largest “work of fracture” measured of any specimen in this work, 7.540 kJ/m2, and the largest biaxial flexure strength measured, 313 MPa. (Pavlacka measured ~800 J/m2 “work of fracture” and ~180 MPa flexure strength in his 2-2 series alumina microstructure composites, and Clegg measured ~4.600-6.700 kJ/m2 and ~500 MPa in his weak-interface SiC/graphite laminates.)

35

Figure 3.15. Micrograph of a 2-2 parallel alumina microstructure composite (a) and the respective load/deflection curve (b). While the composite shows extensive step-wise fracture and ~100 µm deflection along the textured layers, the load deflection curve indicates catastrophic failure. Still, the “work of fracture” of this specimen, 7.540 kJ/m2, is the highest seen in this work.

It should be noted that the deflection axes of these plots (and all subsequent load/deflection curves in this work) have been modified to begin at 0; the original data reflected higher deflection with no load, which represented the lowering of the loading ram to the specimen and alignment of the compliant layers. Deflections fell in the range of 0.60 to 1.0 mm, typical of the SiC/graphite laminates (deflection ~1.0 mm).2

The specimens shown in Figure 3.16 also show no “pop-in” characteristic in their load/deflection curve, indicating catastrophic failure. These specimens represent a high flexure strength/”WOF” specimen and an average flexure strength/”WOF” specimen.

36

Figure 3.16. Load/displacement curves of 2-2 parallel alumina microstructure composites which failed catastrophically. These samples represent average and high “work of fracture” values for this connectivity. The specimen on the left had a biaxial flexure strength of 167 MPa, and the specimen on the right, 247 MPa.

Conversely, some 2-2 parallel composites did show signs of graceful failure in their load/deflection curves. Each pop-in feature represents the failure of an initial layer and the subsequent loading of the next layer. Like the specimens described above, these were variable in strength and “work of fracture”.

Figure 3.17. Load/displacement curves of 2-2 parallel alumina microstructure composites which failed non-catastrophically. Each pop-in feature represents failure of a layer and loading of the next; this fracture behavior is known as “graceful failure.” The specimen on the left has a biaxial flexure strength of 154 MPa, and the specimen on the right, 194 MPa.

37

It is apparent that the mechanical behavior of the 2-2 parallel composite is variable. This may be attributed to the architecture of this connectivity. Recall that the 2-2 parallel connectivity is fabricated by stacking composite tapes of textured and equiaxed stripes, in which the textured stripes are stacked above each other and the equiaxed stripes are stacked above each other. This creates a composite in which the textured and equiaxed regions are self-connected in two- dimensions, in the direction of the cast and in the direction of loading. This connectivity is unique from the 1-1 and 3-3 connectivity in that a crack may travel through the entire specimen without reaching a textured, reinforcing layer. In this case, the specimen would behave as a bulk alumina specimen, with no crack deflection or step-wise fracture. While qualitative comparison of the connectivities will be discussed later, the fracture behavior of the 2-2 parallel connectivity would indicate that it is not a reliable architecture for applications requiring high toughness due to its variable mechanical response.

1-1 Connectivity

To date, the 1-1 connectivity has been fabricated but not mechanically tested. Again, the 1-1 connectivity consists of the regions of equiaxed and textured material connected in only one direction, the casting direction. One may picture this as a series of edge-connected horizontal columns; a schematic is shown in Figure 3.18. This is also a parallel connectivity, as each layer contains both textured and equiaxed regions. Unlike the 2-2 parallel connectivity, any crack propagating through this architecture will encounter a reinforcing, textured region.

38

Figure 3.18. Schematic of a 1-1 microstructure composite.

Like the 2-2 parallel connectivity, crack deflection and step-wise fracture are observed in the 1-1 connectivity; however, the behavior of deflection in the 1-1 connectivity is more like that observed in the 2-2 series samples. Similar to the 2-2 series connectivity, deflection in the 1-1 begins within the textured layer. Figure 3.19 shows a 1-1 microstructure composite in which the crack propagated transversely through the equiaxed layer, continued to propagate transversely at the onset of the textured layer (indicated by the dotted line), and “kinked” out well into the textured layer to begin propagating longitudinally. The initial transverse cracking through the textured layer is transgranular. In this case, residual stresses would not explain the cause of deflection. Pavlacka suggested that in these cases, deflection is not caused by a weak interface or by residual stresses, but rather by a flaw found in the textured layer, initiating crack propagation along the basal surface of the textured grains.5

39

Figure 3.19. Micrograph of a 1-1 alumina microstructure composite in which crack deflection begins within the textured layered. That is, the crack propagates transversely through the interface between the equiaxed and textured regions (denoted by the dotted line) and deflects ~30 µm into the textured layer.

This transgranular propagation through the textured layers is seen through entire textured layers as well, in which no deflection occurs. This is similar to the fracture behavior Pavlacka observed in monolithic textured alumina. Figure 3.20 shows a 1-1 microstructure composite which exhibits both deflection and transgranular fracture through textured regions.

40

Figure 3.20. Micrograph of a 1-1 alumina microstructure composite demonstrating both deflection through a textured layer (denoted by the black arrows) and no deflection (transgranular fracture) through a textured layer (denoted by the dotted lines).

A possible explanation for this behavior is found in the architecture of the 1-1 connectivity. Because the textured regions are only connected in one dimension, it is possible that the crack “found” the edge of the textured column (within the dotted line) and could not be deflected. The crack would then propagate transversely until it encountered the textured column in the next layer. A closer look at the interface between the equiaxed and textured regions, seen in Figure 3.21, shows the nature of fracture through each layer.

41

Figure 3.21. The interface between the equiaxed and textured layers in a 1-1 alumina microstructure composite in which no deflection is observed through the textured layer. Fracture through the equiaxed layer (above dotted line) is intergranular, and fracture through the textured layer (below dotted line) is transgranular.

Extended crack deflection (~1 mm), similar to that seen in Pavlacka’s 2-2 series microstructure composites, is observed in the 1-1 microstructure composites. Pavlacka observed that crack deflection of this magnitude could increase the “work of fracture” of a specimen 100-800 J/m2.

42

Figure 3.22. Micrograph of a 1-1 alumina microstructure composite showing extended crack deflection (~1 mm). The arrows indicate two directions of crack deflection, a result of the biaxial loading. This composite had a biaxial flexure strength of 128 MPa and a “work of fracture” of 2,740 kJ/m2.

It should be noted that this micrograph shows crack deflectionin two directions (as indicated by the arrows) due to the biaxial nature of the loading. The biaxial nature of the loading, and therefore of crack deflection, in the microstructures composites in this work would account for the drastic increase in “work of fracture” over that seen in Pavlacka’s work (~800 J/m2).

With few exceptions, 1-1 connectivity microstructure composites exhibited graceful failure, as indicated by the load/deflection curves, shown in Figure 3.23.

43

Figure 3.23. Load/deflection curves of 1-1 alumina microstructure composites, all showing “pop- in” characteristics indicative of graceful failure. The respective workof fracture values are shown.

The less variable behavior of composites with 1-1 connectivity relative to samples with2- 2 parallel connectivity is attributed to the fact that a crack cannot propagate through a 1-1 connectivity microstructure composite without encountering a reinforcing, textured layer. Biaxial crack deflection, step-wise fracture, and extended crack deflection (~1 mm) exhibited by this connectivity yielded specimens with unique fracture behavior and non-catastrophic failure.

3-3 Connectivity

Like the 1-1 connectivity, 3-3 connectivity microstructure composites have, to date, been fabricated but not mechanically tested. Recall that the 3-3 connectivity wass fabricated by the rotating each successive co-cast tape 90°, resulting in a composite in which the texture regions run in both longitudinal directions; a schematic is shown in Figure 3.24. It was thought that the biaxial nature of the reinforcing, textured material would have distinct fracture behavior when compared to the samples with 2-2 parallel and 1-1 connectivities. The fracture behavior,

44

however, is very similar to that of the 1-1 composite and the characteristic strength (discussed later) is only slightly improved.

Figure 3.24. Schematic of a 3-3 microstructure composite.

Both mechanisms of deflection through a textured layer already observed (beginning at onset and within the textured layer) are seen in the 3-3 connectivity, as shown in Figure 3.25. As discussed earlier, deflection at the onset of the textured layer (or, at the interface between the equiaxed and textured regions) is most likely caused by residual stress development from thermal expansion mismatch. Deflection within the textured layer indicates that deflection was induced by a flaw found within the layer.

45

Figure 3.25. Micrographs of 3-3 alumina microstructure composites showing two different mechanisms of deflection through a texture layer: a) deflection begins at the interface between the equiaxed and textured regions, and b) deflection begins within the textured layer, after propagating transversely through the interface between the equiaxed and textured layers, indicated by the dotted lines.

It is interesting to note that the specimen in Figure 3.25b had a biaxial flexure strength of 94 MPa, the lowest observed in this work; the “work of fracture”, however, was 5.349 kJ/m2, further demonstrating the variability in mechanical response that these microstructure composites exhibit. The specimen in Figure 3.25a had a biaxial strength of 145 MPa and a “work of fracture” of 3.152 kJ/m2.

Extended crack deflection is also seen in the 3-3 connectivity, as shown in Figure 3.26.

46

Figure 3.26. Micrograph of a 3-3 alumina microstructure composite that exhibits extended crack deflection (~1.5 mm), denoted by the arrow. This behavior can drastically change the “work of fracture” of a specimen. The biaxial flexure strength of this specimen is 145 MPa and the “work of fracture” is 3.153 kJ/m2.

Figure 3.27 demonstrates a fundamental idea regarding crack deflection in these microstructure composites. The micrograph shows the outer edge of two textured regions, denoted by the arrows. Had the propagating crack traveled just left of the end of these regions, it would not have encountered the reinforcing layer and would not have been deflected. This demonstrates the unique effect that complex architectures can have on composite materials.

47

Figure 3.27. Micrograph of a 3-3 microstructure composite showing deflection along textured layers. The ends of the textured layers (or “stripes” in the green tape) are denoted with the arrows. Had the propagating crack traveled just left of the end of these layers, it would not have been deflected.

Like the 1-1 connectivity, most of the 3-3 connectivity microstructure composites exhibited graceful failure, as seen in Figure 3.28.

48

Figure 3.28. Load/deflection curves of 3-3 alumina microstructure composites, all showing “pop- in” characteristics indicative of graceful failure. The respective “work of fracture” values are shown.

Again, the less variable behavior exhibited by the 3-3 connectivity as compared to the 2-2 parallel connectivity is attributed to the fact that a crack cannot propagate through a 3-3 connectivity microstructure composite without encountering a reinforcing, textured layer. While the 3-3 connectivity is the only connectivity in which the textured regions ran in two directions (due to the 90° rotation of the composite green tapes), this did not have an appreciable effect on the fracture behavior, or as will be shown in the next section, on the biaxial flexure strength. It should be noted that the mean “work of fracture”, 4.280 kJ/m2 (S.D. = 1.550 kJ/m2), of the 3-3 was higher than both the 1-1 (3.520 kJ/m2, S.D. = 970 J/m2) and the 2-2 parallel (3.950 kJ/m2, S.D. = 1.850 kJ/m2) connectivities; however, these findings are not statistically different according to the Student’s t-test. This could indicate that crack deflection and step-wise fracture are more prevalent in the 3-3 connectivity or that the architecture has another toughening effect on fracture behavior that was unobserved. Nevertheless, the crack deflection, step-wise fracture,

49

and extended crack deflection (~1.5 mm) exhibited by the 3-3 connectivity yielded unique fracture behavior in these microstructure composites with a positive effect on the “work of fracture” and nature of failure. A summary of the “work of fracture” results is shown in Table 3.1.

Table 3.1. Mean “work of fracture”, standard deviations, and number of specimens tested for connectivity samples.

Connectivity Mean “WOF” Std. Dev. N (kJ/m2) (kJ/m2) 1-1 3.520 0.970 10 2-2 parallel 3.950 1.850 10 3-3 4.280 1.550 12

3.3.5 Equibiaxial Strength

The results of the ring-on-ring flexure testing are shown in Table 3.2. A two-parameter Weibull analysis was used to determine the Weibull modulus, m, and characteristic equibiaxial strength, σo, of each connectivity. It should be noted, however, that these results were not found to be statistically different according to the Student’s t-test.

Table 3.2. Weibull moduli, characteristic strengths, and number of specimens tested for connectivity samples.

Connectivity m σo (MPa) N 1-1 7.7 154.5 10 2-2 parallel 20.7/1.6 165.3/160.2 7/3 3-3 5.3 166.6 12

The 2-2 connectivity samples showed a bimodal flaw distribution, as seen in the Weibull plots in Figure 3.29. Bimodal flaw distributions are characteristic of multiaxial stress states, and therefore, are not unusual for a biaxial strength test; values of m and σo for both flaw segments are reported.

50

Figure 3.29. Two-parameter Weibull plots for 1-1, 2-2 parallel, and 3-3 connectivity samples tested in ring-on-ring equibiaxial flexure.

The 3-3 connectivity samples showed the highest characteristic strength (166.6 MPa), but also a slightly lower m value (5.33) than the 1-1 connectivity (7.67), indicating an increased strength variability over the 1-1. The distribution of the first flaw population of the 2-2 parallel connectivity had very low strength variability (m=20.7), and a characteristic strength (165.3 MPa) very similar to the 3-3 connectivity. The second flaw population, however, had high strength variability (m=1.6), which is to be expected from the small number of specimens it includes (3 samples only). The characteristic strength (160.2 MPa) is lower than that of its 2-2 counterparts. The cause of the bimodal distribution is unknown. While fractographic evidence is commonly very useful in determining the cause for a bimodal flaw distribution, such as a difference in flaw origin, fractographic analysis does not reveal a distinct change in fracture pattern between the samples of the upper and lower flaw segments. Examples of macroscopic fracture patterns of the test specimens for both segments are shown in Figure 3.30.

51

Figure 3.30. Fracture patterns for 2-2 connectivity samples; (a) and (b) are samples from the lower flaw segment and (c) and (d) are samples from the upper flaw segment. Cracks are highlighted for better viewing. No distinct differences in the crack patterns can be discerned between the upper and lower flaw segments.

The only distinction in the patterns is the circumferential nature of the fracture pattern in sample D, which is only characteristic of high strength failure (sample D had the highest equibiaxial strength, 313 MPa, of the 2-2 samples) and not a distinct flaw. There is no apparent difference in the fracture patterns between the two flaw distributions and the cause of the bimodal flaw distribution remains unclear.

3.4 Conclusion

While the mechanical response of the microstructure composites studied in this work were variable, unique and interesting fracture behavior was observed. Crack deflection, step- wise fracture was readily observed in composites with each of the connectivities tested (1-1, 2-2 parallel, and 3-3). While the cause of deflection remains unclear due to the variable nature of where it began, either at the interface between the equiaxed and textured regions or within the textured regions, two possible explanations are presented. One explanation is that residual stresses are developed due to thermal expansion mismatch between equiaxed and textured

52

alumina. Another explanation is that deflection is flaw-induced. In both cases, however, deflection is intergranular between the basal faces of the textured grains, signifying that there is an anisotropic fracture energy in textured alumina. The interface between the basal faces, therefore, acts as a weak path for deflection.

Extended crack deflection was observed in the 1-1 and 3-3 connectivities; however, this behavior did not always correlate with a higher “work of fracture”. No crack bifurcation was observed in this work.

Graceful failure, indicated by load/deflection curves from flexure testing, was observed in all connectivities. 2-2 parallel composites, however, exhibited both catastrophic and graceful failure. Even in the cases of catastrophic failure, however, high strength and “work of fracture” was calculated. The lack of consistency in the fracture behavior and failure mechanism of the 2- 2 parallel connectivity, as well as the bimodal flaw distribution, indicates that this connectivity is highly variable in mechanical response and may not be reliable for structural applications.

The similarity in fracture behavior and failure mechanism (graceful failure) between the 1-1 and 3-3 connectivities indicates that there is little difference in the effect of these connectivities on mechanical response. This may be due to the fact that unlike the 2-2 parallel, a crack propagating through a 1-1 or 3-3 specimen will always encounter a reinforcing, textured layer. This makes these connectivities good candidates for architectures used in structural composite materials.

While the reinforcement of textured alumina material in these connectivities demonstrated toughening behavior, the low strength of these samples is undesirable for structural applications. Applying the concept of connectivity to another material that exhibits higher strength than alumina could further illuminate the benefits of connectivity and microstructure composites to the toughening of ceramics. Furthermore, these equiaxed/textured alumina microstructure composites were designed to take advantage of the toughening properties of the textured material; microstructure composites of other materials can and should be designed to exploit advantageous properties of both microstructural components.

3.5 References

53

1. T.N. Tiegs and P.F. Becher, “Sintered Al2O3-SiC-Whisker Composites,” American Ceramic Society Bulletin 66[2], 339-342 (1987). 2. W.J. Clegg, K. Kendall, N. McN. Alford, T.W. Button, and J.D. Birchall, “A Simple Way to Make Tough Ceramics,” Nature 347, 455-457 (1990). 3. P.E. Morgan and D.B. Marshall, “Ceramic Composites of Monazite and Alumina,” Journal of the American Ceramic Society 78[6], 1553-63 (1995). 4. J.B. Davis, A. Kristoffersson, E. Carlstrom, and W.J. Clegg, “Fabrication and Crack Deflection in Ceramic Laminates with Porous Interlayers,” Journal of the American Ceramic Society 83[10], 2369-74 (2000).

5. R.J. Pavlacka, "Processing and Mechanical Behavior of Al2O3 Microstructure Composites," Ph. D. Thesis (The Pennsylvania State University, University Park, PA, 2009). 6. ASTM C 1499-09 (2001). 7. M.H. Krohn, J.R. Hellmann, D.L. Shelleman, C.G. Pantano, and G.E. Sakoske, “Biaxial Flexure Strength and Dynamic Fatigue of Soda-Lime-Silica Float Glass,” Journal of the American Ceramic Society 85[7], 1777-82 (2002). 8. Illustration by M. Fleck, Pennsylvania State University.

54

4 Fabrication of α-/β-SiAlON Microstructure Composites: Processing Issues and Initial Results

4.1 Introduction

SiAlON microstructure composites were designed and fabricated in the 1-1, 2-2, and 3-3 connectivities, as was done with the alumina microstructure composites, to exploit the high hardness α phase and high toughness β phase. As previously stated, the chemical formula for α-

SiAlON is MxSi12-(m-n)Al(m+n)OnN(16-n) and for β-SiAlON is Si6-zAlzOzN8-z, where 0

The parameters m and n in the α phase formula describe the number of Si-N bonds replaced by Al-N and Al-O bonds, respectively. The parameter x is related to the valence of the metal cation, M; commonly used metal cations include Yb, Y, or Nd.2 β-SiAlON is achieved by the partial substitution of Al-O bond for Si-N bonds, described by the chemical formula Si6- 3 zAzOzN8-z, where 0

55

4.2 Design and Testing of the α- and β- SiAlON Phases

4.2.1 Design Considerations

The α-SiAlON formula is Y2/3Si9Al3ON15 (or MxSi12-(m-n)Al(m+n)OnN(16-n) where x = 2/3, m = 2, and n = 1). According to the literature, these parameters yield a α phase microstructure of equiaxed grains of 0.5-1.0 µm with a hardness value of 18-21GPa.2,4 Optimization of the β-SiAlON formula was performed to ascertain a β phase with high fracture toughness (4.7 MPa⋅m1/2).5 The liquid phase former, yttrium, was added at various weight percentages (4 wt% and 7 wt%) to optimize fracture toughness. Furthermore, samples were fabricated by adding yttria both as a powder or as a nitrate (Y(NO3)36H2O). Fracture toughness measurements were made on all variations to determine the compositions for the optimized composite.

4.2.2 Experimental Procedure

The starting powders were α-Si3N4 (SN-E10, UBE Industries, Yamaguchi, Japan), AlN

(Type F, Tokuyama Corp., Tokyo, Japan), α-Al2O3 (AKP-50, Sumitomo Chemical CO, Ltd.,

Tokyo, Japan), and Y2O3 (UU, Shin-Etsu Chemical Co., Ltd., Tokyo, Japan, 350 nm, 99.99% purity); in some cases, (Y(NO3)36H2O) (Alfa Aesar) was used instead of yttria powder. When yttrium nitrate was used, it was first dissolved in an ethanol/xylenes mixture before being added to the slurry. Appropriate amounts of the powders were mixed according to the above formulas for the α- and β-SiAlON phases and were ball-milled in ethanol for 24 h. The powders were then dried, sieved, and hot pressed (into monolithic α- and β-SiAlON samples) in a 50.8 mm (2”) graphite die at 1800˚C/20 MPa for 1 hour in a vacuum atmosphere. The powders were not pressed as pellets first, as the green bodies would crack on the onset of the applied pressure, creating a “pre-cracked” sintered body; this will be further discussed later. The powders were embedded in boron nitride powder, which acted both as a lubricant and prevented carbon contamination from the die. Initially, powders were pressed in a 50.8 mm (2”) diameter hot press die that was 101 mm (4”) in height; however, this die proved to be too “short” to fabricate samples of sufficient thickness (~3 mm) for mechanical testing. Fabrication of 3 mm thick specimens required a die

56

152 mm (6”) in height, which allowed for more powder to be charged, and thicker samples to be made. Samples were hot pressed into approximately 50.8 mm diameter discs, 3-4 mm in height. Following hot pressing, the samples were sand-blasted to remove the embedding powder and ground flat on both sides using a surface grinder. The tensile sides of the specimens were then polished to a 1 µm finish before being machined into bend bars. Bend bars were cut from the hot pressed discs for fracture toughness testing. The fracture toughness was measured using the single-edge notched-beam method in four-point bending, with a notch depth approximately 30% of the specimen thickness. Final bar dimensions were approximately 40 mm x 4 mm x 3 mm with ~1 mm deep notches on the tensile sides. A SiC four-point bend apparatus was used with inner and outer spans of 10 mm and 20 mm, respectively. An Instron (Model 5866, Instron Corp., Canton, MA) load frame was used to drive the loading ram in testing, applying the load to the samples at a constant displacement rate of 0.1 mm/min. A minimum of 8 specimens was tested for each variation of the phases. Hardness was determined by the Vickers method using a micro-hardness indenter (LECO V-100-C1 Hardness Tester, LECO Corporation, St. Joseph, MI) at a load of 300 g. Both α- and β- SiAlON phases were analyzed using an x-ray diffraction (Scintag x2 Diffractometer, Scintag, Inc.; Cupertino, CA). Samples densities were determined by the Archimedes method. Because the density of SiAlON (3.0-3.2 g/cm3) varies based on phase, chemical formula, and dopants, it is not reported as a theoretical density.

4.2.3 Results

The α-SiAlON phase yielded expected results in both fracture toughness and hardness. 1/2 The α phase had a KIC of 2.5 MPa⋅m and a hardness of 20 GPa. This is ideal for the α phase as the composite is meant to take advantage of its high hardness, while taking advantage of the high toughness values of the β phase. Microstructures of the α phase are shown in Figure 4.1.

57

Figure 4.1 α-SiAlON a) fracture surface and b) polished surface etched for 15 min in KOH.

The β-SiAlON phase, however, did not yield the expected fracture toughness. According to the literature, the fracture toughness of β-SiAlON doped with yttria should be 4.7 MPa⋅m1/2.5 In this study, the β phase doped with 4 wt% yttria yielded a fracture toughness of ~3 MPa⋅m1/2. This indicates that 4 wt% yttria liquid phase was not sufficient in promoting the growth of anisotropic grains in the β phase. This is confirmed by the microstructure, shown in Figure 4.2, showing few elongated grains that would promote crack deflection and increase toughness.

Figure 4.2 4 wt% yttria-doped β-SiAlON a) fracture surface and b) polished surface etched for 15 min in KOH.

The amount of yttria liquid phase was increased to 7 wt% to increase the toughness of the material. This was done with both yttria powder and yttrium nitrate. The addition of 7 wt%

58

yttria powder unexpectedly lowered the fracture toughness of the β phase to ~2.8 MPa⋅m1/2. 1/2 However, the addition of 7 wt% yttrium nitrate raised the KIC value to 3.8 MPa⋅m . Figure 4.3 compares the microstructures of the β-SiAlON samples doped with 7 wt% yttria powder and yttrium nitrate.

Figure 4.3. Polished surfaces of β-SiAlON doped with 7 wt% a) yttria and b) yttrium nitrate. Samples etched in molten KOH for 15 min each. Arrows indicate pockets of etched liquid phase.

The sample doped with the yttrium nitrate had a more uniformly distributed liquid phase and a more defined elongated grain structure whereas the sample doped with the yttria powder shows only small pockets of liquid phase (indicated by arrows in the figure) and many more equiaxed grains. The unexpected fracture toughness of both the β-SiAlON materials processed with the yttria powder (4 wt% and 7 wt%) indicate that yttria powder is not an ideal form of adding the liquid phase dopant in a SiAlON tape casting slurry. The yttrium nitrate, which is first dissolved in ethanol/xylenes mixture before being added to the slurry, is a more effective way of obtaining a more homogenous mixture resulting in a continuous liquid phase in a tape cast SiAlON specimen.

59

In summary, the α- and β-SiAlON phases to be used in the microstructure composites are

α-phase: Y2/3Si9Al3ON15 and β-phase: Si5AlON7 doped with 7 wt% yttrium nitrate. X-ray diffraction patterns for each phase are shown in Figure 4.4.

60

Figure 4.4. X-ray diffraction patterns for the α- and β-SiAlON phases, respectively.

The α-SiAlON phase is not phase pure, as it contains Si3Al3O3N5 and Si4Al2O2N6 elements as 3 well as some β-SiAlON phase, Si5AlON7. The density of the sample is 3.13 g/cm . The β-phase 3 contains some Si3N4; the density of the sample is 3.11 g/cm .

4.3 SiAlON Microstructure Composite Fabrication and Processing Issues

4.3.1 Microstructure Composites

61

Fabrication of 1-1, 2-2 parallel, and 3-3 SiAlON microstructure composites followed the same co-casting/lay-up methods as those used with alumina. The stacking configurations are shown again in Figure 4.5.

Figure 4.5. Connectivities and their respective tape stacking configurations. The dark stripes represent α-SiAlON and the light stripes represent β-SiAlON.

Instead of equiaxed/textured alumina composite, the SiAlON microstructure composites consisted of α- and β-SiAlON phases, optimized for high hardness and high fracture toughness values, respectively.

4.3.2 Experimental Procedure

The same starting powders used above for the monolithic phase samples were used for the composites. Tape casting slurries for α- and β-SiAlON were prepared and co-casted in the same fashion as discussed for the alumina slurries. The same dispersant, solvent, and organics system was used, with the formulation slightly modified for the different powders. Composite α/β tapes were consequently cut and stacked according to the desired connectivity. Samples of the 1-1, 2- 2, and 3-3 connectivities were fabricated in the same method as their alumina counterparts.

The samples were isostatically laminated at 75°C and 20 MPa for 40 min. The organics were removed using a burnout cycle in air at 600°C for 12 h with a heating rate of 0.1°C/min to prevent cracking of the green parts. All samples were hot pressed in a 50.8 mm (2”) graphite die

62

at 1800˚C/20 MPa for 1 hour. Samples were embedded in BN powder, which both acted as a lubricant and prevented carbon contamination from the die.

Hot pressed samples, approximately 50.8 mm in diameter and 3 mm in height, were sandblasted to remove the embedding powder and ground flat using a surface grinder; flatness was confirmed by measuring the thickness of various areas across the plate with a micrometer. The tensile side was polished to a 1 µm finish and the edges beveled to remove edge cracks. Equibiaxial strength testing (ROR method) followed the same procedure as that used on the alumina microstructure composites.

4.3.3 Processing Issues

Many processing issues were encountered in the fabrication of SiAlON microstructure composites. While most dealt with the difficulties of pressure-assisted sintering, others arose during each of the steps of the tape casting process. Some composite tapes showed furrows between the α and β phase “stripes” due to mismatched slurry viscosity which resulted in cracking upon drying. To mitigate this problem, care was taken to ensure that the slurries were similar in viscosity prior to casting by adjusting the addition or evaporation of solvent. Many samples were also lost in the green state. Thick (7-10 mm) stacked samples tend to fail during lamination due to leaks in the vacuum bags. Furthermore, samples failed in the burnout cycle, as seen below in Figure 4.6.

63

Figure 4.6. α-SiAlON green body which failed during the burnout cycle.

To resolve this issue, the tape casting formulation was altered to reduce the weight percentage of binder and plasticizers (PVB, PEG, and PAG), first by ½, to 2.4, 2.0, and 2.2 wt%, respectively. This reduction proved to be too much, as the subsequent tapes were far too brittle and cracked on drying. The original organic formulation was then reduced by ⅔ (to 3.2, 2.7, and 3.0 wt%) resulting in tapes with good plasticity and strength. The reduced amount of organics prevented any subsequent green bodies from cracking during burnout.

Several samples also failed during the hot pressing cycles. Before boron nitride was used as embedding powders, samples were protected from the die walls using layers of Grafoil. Samples out of the hot press, however, appeared degraded, as seen below in Figure 4.7. Some had a core shell structure, with a densified SiAlON core in a degraded shell, also shown in Figure 4.7.

64

Figure 4.7. β-SiAlON samples which failed during hot pressing: a) complete SiAlON degradation, and b) partial degradation resulting in a core-shell structure.

Grafoil layers were not sufficient to protect the samples from degradation during the hot press cycle. Boron nitride powder was instead used to embed the samples, separating them from the die walls (this also aided in ejecting the sample at the completion of the hot press cycle). This was much more successful is obtaining dense samples and pyrolysis was no longer observed. While boron nitride powder is ideal for this system as it is non-reactive with SiAlON, another problem arose from its use. All dense tape cast samples out of the hot press were “pre- cracked,” as seen in Figure 4.8.

65

Figure 4.8. SiAlON microstructure composite showing cracks originated during the hot-pressing cycle.

While it is unknown at which point in the hot press cycle the sample was broken, the possible scenarios are: the part is cracked on a) loading of the hot press (while aligning the die with the rams) before the cycle is even started, b) the onset of the application of pressure following ramp up to 1800˚C, or c) during the dwell time at pressure (20 MPa). A pressureless sinter cycle was attempted to avoid the problem altogether, but densification was not achieved. Furthermore, several attempts to change the temperature and pressure ramp rates were unsuccessful in preventing cracking. At the basis of each of these scenarios, however, is the same issue: the fine boron nitride powder does not effectively transmit the pressure from the pressing rams due to low particle fluidity. An attempt was made to mitigate this problem by using 1 mm diameter zirconia milling media as the embedding “powder,” given its greater fluidity compared to the fine BN powder. This attempt, however, was unsuccessful as the sample was also pre-cracked. Given the low strength of tape cast green bodies, this problem presents a great hurdle in hot pressing tape cast parts. Attempts may be made with other, more fluid powders (less extreme than the use of milling media) to transmit the pressure between the upper and lower loading rams without cracking the green body. More attempts at pressureless sintering cycles could also be made. Finally, a different material, which does not require pressure-assisted densification, may be considered.

66

4.3.4 Initial Mechanical Testing Results

The first six samples of SiAlON microstructure composites densified (all of the 3-3 connectivity) were tested in equibiaxial flexure (ring-on-ring method). Strength values were low, ranging from 120 MPa to 270 MPa. Furthermore, no toughening characteristics were observed in the load/deflection curves of the specimen. Two examples of the load/deflection curves of 3-3 connected α-/β- SiAlON microstructure composites are shown in Figure 4.9.

Figure 4.9. Typical load/deflection curves of the α-/β-SiAlON 3-3 microstructure composites tested. No “pop-in” features are observed, which is indicative of catastrophic failure.

67

Neither show any “pop-in” characteristic of increased toughening in the material as was observed in the alumina microstructure composites. This is indicative of catastrophic failure in these samples. The low flexure strength and catastrophic nature of failure is most likely due to the “pre-cracking” seen in these samples from hot-pressing.

4.4 Conclusion

Microstructure composites of α-/β-SiAlON were designed to exploit the high hardness of the α phase and the high toughness of the β phase. The chemical formulas of the individual phases, Y2/3Si9Al3ON15 and Si5AlON7, were selected to optimize these properties. 7 wt% yttrium nitrate was added to the β phase to encourage the anisotropic growth of the β grains. The fracture toughnesses of the α and β phases were 2.5 MPa⋅m1/2 and 3.8 MPa⋅m1/2, respectively.

Many processing issues were encountered in fabricating microstructure composites of these SiAlON phases. Failure in both the green state and during densification was observed. Failure in the green state occurred during co-casting, lamination, and binder burnout. Furthermore, extensive degradation of the samples was observed during hot-pressing. While these issues were resolved, the issue of the sample “pre-cracking” during hot-pressing was not. Due to the weakness of green, tape-cast parts, all specimens were cracked during some stage of hot-pressing by the loading ram. Further work is required to solve this processing issue.

4.5 References

1. T. Ekström and M. Nygren, “SiAlON Ceramics,” Journal of the American Ceramic Society 75[2], 259-76 (1992). 2. P. Sajgalik, Z. Lences, and M. Hnatko, “Nitrides,” in Ceramics Science and Technology, Volume 2, edited by R. Riedel and I-W. Chen, 59-89 (2010). 3. G. Z. Cao and R. Metselaar, “α’-SiAlON Ceramics: A Review,” Chemistry of Materials 3, 242-252 (1991). 4. M.I. Jones, H. Hyuga, and K. Hirao, “Optical and Mechanical Properties of α/β Composite SiAlONs,” Journal of the American Ceramic Society 86[3], 520-22 (2003).

68

5. A.A. Kudyba-Jansen, H.T. Hintzen, and R. Metselaar, “The Influence of Green Processing on the Sintering and Mechanical Properties of β-SiAlON,” Journal of the European Ceramic Society 21, 2153-2160 (2001).

69

5 Future Work

The idea of microstructure composites remains a young concept with great potential for adaption and application to new materials. For the equiaxed/textured alumina microstructure composites studied in this work, more mechanical characterization could further reveal the effect of connectivity on mechanical behavior. Adapting the ring-on-ring concept to measure fracture toughness of these composites could illuminate the ability of the textured microstructure to toughen alumina. Furthermore, changing the size and scale of the microstructure composites and their constituent components reveals endless possibilities for designing highly complex and tailored architectures. This could involve changing the thicknesses of the component stripes or changing the thicknesses of the each of the layers for further crack deflection. Also, other materials could be used to make microstructure composites to exploit desired properties for various applications. Regarding the α-/β-SiAlON microstructure composites, further work could solve the problem of pre-cracking during hot-pressing. Ascertaining at which point during the hot- pressing schedule that cracking occurs is a necessary step towards achieving this goal. With this knowledge, changes could be made to the hot-pressing schedule, to the die, or to the embedding material to mitigate this problem. Embedding the samples with different materials should also be tested. Pressureless sintering could be further examined for the system as well.

70