Mechanical Properties of Ceramics and Composites

Very few evaluations of effects of particle parameters on compressive strength, ballistic performance, or wear have been made. One set of tests on crystallized glasses at room temperature showed a substantial range of compressive strengths, including some fairly substantial strengths to 100 MPa (up to 150% of the parent glass, Figure 10.2) at 22°C, with varying rates and types of decreases as test temperatures increased (Fig. 11.13). Similarly a limited study showed compressive strengths of up to 500+ MPa in a synthetic composite of SiC platelets in a glass matrix at 22°C, with marked decreases as test temperatures increased, but with composite strengths still well above those of the matrix glass alone. Composites of laminar particles of refractory metal and ceramic layers have shown respectable strengths, e.g. > 1 GPa can be obtained, which is probably due to the finer particle (and grain) size. However such strengths were also a function of the laminar particle dimensions, showing again that other microstructural dimensions can play a role in mechanical properties.

One limited evaluation showed that compressive fatigue crack propagation occurred in a ceramic whisker composite, but not always unfavorably relative to the matrix alone. Limited testing has shown that the ballistic stopping power of a dense, fine particle (and grain) size composite approached performance levels of established monolithic ceramics for armor.

More extensive, but still limited, wear testing of similar composites clearly shows that they can be competitive to established wear resistant ceramics. This is corroborated by some commercial production and use of ceramic composites for wear and related applications, e.g. Al2O3-TiC for cutting tools and various wear applications and Al2O3-SiC whisker composites for cutting tools and critical wear components for dies for deep drawing of beverage and other cans with an integral bottom.

Even less testing of the above properties has been made at elevated temperature. Testing of a set of Li2O-SiO2 crystallized glasses showed net decreases in strength to the limit of testing (700°C, Fig. 11.13) as was expected. However there was a diversity of behavior at more modest temperatures, including widely varying rates of strength decrease with increasing temperature and some modest, temporary strength rises to maxima, some modest, temporary minima, or both. These trends suggest a fairly diverse range of effects, probably due to mismatch stresses.

Little or no data exists on effects of particle or grain shape or orientation on properties of composites addressed in this section. However, while investigation is limited, results clearly show that composites can have good hardness, respectable compressive strengths and ballistic performance, and good wear resistance, all primarily associated with finer microstructures commonly achievable in composites. Thus further investigation and development should be fruitful.

F.Mechanisms Controlling Normal Brittle Tensile Strength Failure of Monolithic and Composite Ceramics

Given the common assumption that crack bridging, related R-curve effects, and resultant increased toughness with larger scale crack propagation determined tensile strengths and reliability of materials exhibiting such behavior of monolithic and many composite ceramics, it is useful to review the combined evidence for monolithic and composite ceramics on this and the alternate mechanism. As noted in earlier chapters, many investigators conducted simple, especially indentation, toughness tests and showed or noted observations of crack bridging in their monolithic, and especially composite, ceramics, assuming that these arrested crack observations showed behavior across a range of crack velocities and sizes. However, extensive evidence shows that such effects are far more limited in their control of strengths and reliability than was previously thought, applying primarily to strengths retained after serious damage from thermal stress or shock or mechanical impact, i.e. when cracks on the scale of those showing bridging-R-curve effects are developed in components or test specimens. Before proceeding to this summary, it is important to note that two related factors played a major role in the over emphasis of the role of these effects, which are an important component of crack propagation in brittle materials, in determining tensile strength and reliability of some monolithic and many composite ceramics. First was a frequent and often extensive neglect or rejection of the literature, e.g. the rediscovery (and naming) of bridging phenomena, i.e. neglecting earlier observations of bridging in both ceramics and rocks, and the rejection of fractography as a viable tool for corroborating and understanding observations. Second was a narrow range of testing, i.e. limiting or precluding opportunities for testing the self-consistency of data, and interpretation and concepts, with both being important added motivations for this book.

Turning to the summary of evidence seriously questioning, or contrary to, the application of bridging and related, typically large crack phenomena, to normal small-crack strength behavior of monolithic and composite ceramics, major factors are summarized in Table 12.1 for brittle fracture, e.g. at nominally 22°C, under the two headings of test issues and strength/toughness behavior. The most significant and extensive evidence of major discrepancies are those of the microstructural dependences of toughness and strength and of the inconsistencies between strengths and toughness often encountered at high toughness where crack bridging, branching and R-curve effects are most substantial. However, the test issues are important, e.g. starting with the observation that toughness values derived from fractography of failed components or test specimens should be, and are generally found to be, most consistent with strength [33,34]. Further, effects of surface finishing and

TABLE 12.1 Summary of Factors Questioning or Contrary to the Application of Large-Scale Crack Bridging and R-Curve Effects to Normal Small-Crack Strength Behavior of Dense Monolithic and Composite Ceramics

(A) Test Issues

(B) Strength/Toughness Behavior

crack velocity on bridging and related effects are a fruitful area for further research. Again, fractography is important, e.g. in this case via effects on mirror and related dimensions.

Another major reason for not accepting large crack toughness results reflecting R-curve and related effects for predicting most strength behavior is the evidence and results for the alternative mechanism extensively addressed in this book. This is based on modest size flaws introduced, especially from machining, that control most strengths, as is shown by extensive fractographic studies. Extensive studies also show that machining controls strengths of monolithic ceramics via effects of both the abrasive size and its direction in machining relative to subsequent stressing to measure strength [33]. Though there has been less such study for ceramic composites (presumably because of the assumption that the real control of strengths was via toughness as measured by normal large crack tests), more limited studies clearly show similar machining dependence of composite to monolithic ceramics. These observations are consistent with, and explain, the long established strength–G-1/2 behavior of monolithic ceramics, as well as the essentially identical strength–D-1/2 behavior shown for ceramic composites in this book. In both cases the thesis is that machining flaws introduced depend not only on the machining conditions (i.e. abrasive, machining depth of cut, speed, etc.) but also on material parameters, specifically, Young’s modulus, hardness, and toughness of the material locally around induced flaws, (e.g. per Eqs. (3.2) and (5.2).

Three factors show that such machining-induced flaws are consistent with the microstructural dependence of tensile strength of monolithic and composite ceramics and the generally limited or no effect of large crack R- curve effects on tensile strength. First, such flaws are of modest size, e.g. commonly from 10 m to a few tens of microns, reflecting the transient indentation of abrasive particles and associated crack generation, as well as probable local elevated temperatures and some residual surface compressive stresses. Small, rather than large, crack toughness values should correlate with the formation of such flaws. Thus propagation of such flaws to determine strength is also likely to be controlled by toughness values for small to modest crack sizes, hence explaining the general lack of R-curve effects and thus frequent opposite trends of large crack toughnesses and strengths. Second, the formation of such machining flaws is consistent with the microstructural dependence of tensile strength due to second-order effects of microstructural dependence of flaw size c on body properties per Eq. (3.2), especially due to effects of hardness (H). Thus since c varies inversely with H, and H varies inversely with D and G, c increases as D and G increase so that tensile strength decreases with increasing D and G, as is broadly observed for monolithic and composite ceramics (the former only having a G dependence). The size of ma- chining-induced flaws also varies directly with E and inversely with K, but E

is not basically dependent on D or G, and small crack toughness values typically have limited microstructural dependence, especially on D and G. Thus even though the dependences of c on E and K are to different powers, their net microstructural dependence is typically limited. Such machining-induced flaw formation also explains the generally poor strengths of platelet composites, since the large platelet–matrix interfaces commonly provide favorable locations for flaw formation and subsequent propagation to failure. Third, there is typically a minimum in c at intermediate volume fractions of dispersed phase that are associated with the typical maximum in toughness (Figs. 9.16 and 9.18, for both smaller and larger cracks, but often more pronounced for the latter). However, this c minimum may be shifted in φ value, level, and shape by effects of E and H dependences on φ and associated microstructural effects.

Again, when cracks causing failure are larger, e.g. such as those controlling residual strengths after serious thermal stress or impact damage, or from very extensive SCG, R-curve and related effects are probably pertinent to such mechanical behavior. This is the case for continuous ceramic fiber composites where simple, e.g. machining, flaws do not grow continuously, accelerating to failure, thus determining strengths. Behavior of whisker, and probably discontinuous short-fiber, composites may be somewhat in between, though the weight of evidence indicates that whisker composites are much closer in behavior to particulate than continuous fiber composites.

Thus, in summary, there are extensive reasons for discounting most bridging and related R-curve effects in large scale crack toughness measurements of brittle fracture, on normal small crack strengths of monolithic and composite ceramics as summarized in Table 12.1. Though studied much less, there is also evidence that such large scale crack effects often also occur in some porous bodies, mainly at intermediate porosity levels, but that this again has at best limited strength effects, mainly with large cracks, e.g. from serious thermal shock or impact damage, not for normal strength behavior [41,42]. Changes to plastic deformation, e.g. by slip in single crystals and grain boundary sliding, in monolithic ceramics change toughness and strength behavior, reducing differences between large and small crack behavior in some cases, but introducing other differences, probably reflecting effects of crack velocity and strain rate effects. Though data is not as extensive for elevated temperature effects on ceramic composites, they show similar changes and differences as for monolithic ceramics. Thus while there are differences between lower and higher temperature behaviors in different ceramic systems, all raise similar issues of differences between large and small scale crack behavior whether due entirely to brittle fracture processes or to processes involving some plastic deformation, making both areas for further study.

IV. NEEDS AND OPPORTUNITIES TO UNDERSTAND AND BETTER USE THE GRAIN AND PARTICLE DEPENDENCE OF MECHANICAL PROPERTIES OF CERAMICS AND CERAMIC COMPOSITES

A.Testing and Evaluation Needs

There are a variety of needs to improve understanding of the microstructural dependence of properties of both monolithic and composite ceramics, and thus their performance. Clearly a basic one is better microstructural characterization using both direct and indirect methods, the latter via property measurements. More direct documentation of grain and particle character is needed. Even approximate or average values are often not given, nor is the origin of many given values (especially what factor was used to convert linear intercepts to “true” sizes). Almost no data addresses size distributions, less exists for qualitative and none for quantitative characterization of shape, and only limited data exists on grain or particle orientation, reflecting additional needs. Improvements in these measurements should be accompanied by better descriptions of specimen fabrication methods and parameters, since these indicate microstructural characteristics, and the use of such indirect characterization increases as its relation to specific microstructural measurements increases. Modern computerized stereology measurements are an important aid in better characterization.

An even more critical need is for a much wider range of property measurements; determining the extent of microcracking is an important goal. While this is needed better to document properties central to defining mechanisms, it is also an important factor in indirect microstructural characterization. Thus, for example, rather than assuming isotropic bodies and properties in many cases where some anisotropy may exist, determining the degree of isotropy of properties can be valuable. However, the most critical need is for much more comprehensive property measurement. Thus failure to measure elastic moduli along with toughness or strength, and to measure only one of the latter two, are serious constraints on understanding both the properties and the mechanisms. Preferably both static and wave methods should be used to determine at least E, since the former, while often less accurate for absolute values, is an important indicator of microcracking, while the latter methods typically are not but are often better for absolute moduli (before any microcracking). However, use of acoustic emission to detect microcracking is often also valuable, especially in conjunction with other tests such as elastic property changes. Further, given the variability of different toughness tests, it is desirable to measure toughness by at least two or three tests, e.g. indentation, indentation fracture, and possibly a third test. Combinations of other tests, e.g. compressive, hardness, and fatigue tests are also valuable.

There are five other expansions of testing that can be valuable. The first is of

other, often nonmechanical, properties, e.g. of thermal or electrical conductivities, especially of composites, since these can often more clearly reveal microstructural factors, especially the onset of percolation, that can also affect mechanical properties, e.g. creep. Second is a broader range of microstructural variables, e.g. of G, D, or volume fraction second phase. The importance of this is clearly illustrated by the very limited amount of data showing the particle size dependence of tensile strengths (Figs.9.2,9.15,9.16). Third is the closely related aspect of testing composites of the same composition and raw materials, but reflecting different mixing or fabrication methods or both, and hence variations in spatial distributions of microstructures. Fourth, measuring tensile strengths with more than one surface finish and as a function of grinding direction relative to test bar axes can be valuable, e.g. showing the degree of strength dependence on machining flaws. This is particularly important for composites, given the importance of machining flaws in their mechanical behavior and the limited characterization and evaluation of machining effects on their behavior. Fifth is a broader range of temperature testing, especially testing of mechanical properties at modest temperatures, instead of assuming (falsely in some important cases) that properties do not change significantly at modest temperatures. While this is particularly true for tensile strength (Figs. 6.12,6,14,6.15,6.18 ), it can also apply to other mechanical properties such as elastic moduli (Fig. 6. 18), hardness, and compressive strength (Fig. 7.6).

B.Fabrication and Processing Opportunities to Improve Ceramics

Briefly consider some broader aspects of fabrication on the development and understanding of the mechanical behavior of ceramics, especially composites. While conventional consolidation of powders is the dominant method of fabrication and will continue to be so, variations on this as well as other fabrication technologies offer opportunities for producing novel bodies, especially composite ones, for study as well as possible production. One such step is extension to finer powders and resultant microstructures, especially on a nanoscale. Thus such bodies can extend properties such as strengths and hardnesses to higher levels provided that contaminants on the very high surface area powders can be adequately removed during consolidation while maintaining desired fine microstructures. This has often not been the case, especially for nominally single phase bodies such as MgF2 (Fig. 3.24) and TiO2 (Fig.4.5), but it has more commonly been achieved in composites, since higher processing temperatures can better drive off adsorbed species, while the composite structure results in mutual restriction of grain and particle growth. However, the large compaction ratios of most, if not all, very fine powders provide substantial production challenges, which may be reduced or circumvented in some cases by alternant fabrication and processing techniques.

One alternate technique is reaction processing [43,44], such as is used for

fabricating mullite-ZrO2 and various oxide–nonoxide composites. While this is more limited in the compositions to which it is applicable, it is useful for a substantial range of compositions and can offer one to three possible important advantages. The first is that in principle it can produce finer grain and particle sizes for given starting particle sizes, since the resultant body microstructuure is generally formed entirely by nucleation and growth of new phases from the reaction, which if controlled should lead to finer, more uniform microstructures. A key identified need to achieve this is to prevent or limit the extent of transient liquid phase formation, since this greatly increases grain or particle growth, often in a heterogeneous fashion. The second potential advantage is that greater microstructural uniformity can be achieved and that this can translate into greater mechanical reliability . The third potential advantage is that the raw material costs for reactants is often substantially lower than for the resultant product phases (Table 10.1), thus making reaction processed composites somewhat more economical. Some systems may also allow more effective consolidation that can also lower costs.

There are two other alternate fabrication routes, namely melt processing and CVD, that, though they are also more restricted in compositions, can offer similar advantages. Thus PSZ polycrystals and especially single crystals as well as ZTA eutectic bodies are examples of promising composites that have been made with encouraging results via melt processing. While thermal stresses and pores from intrinsic liquid–solid density differences and extrinsically from exsolved gases are a serious challenge, various conventional and novel solidification methods offer important opportunities [43–49]. Conventional directional solidification as for PSZ crystals from skull melting and single crystal eutectics are key examples of important successes, including some of the best high temperature strengths (Figs. 11.11,11.12 ). Similarly, though probably less recognized, there are important opportunities for preparation of fine uniform microstructures, again especially for composites via CVD [46,47]. Thus codeposition of different compositions have been explored with some promising results, e.g. the formation of TiN precipitates from 3 to 15 nm in dimensions in a Si3N4 matrix [47].

V.SUMMARY

This book has extensively reviewed and discussed the dependence of primarily mechanical properties of dense monolithic ceramics and ceramic composites on respectively grain and particle parameters, mainly size, but also to the extent feasible on shape and orientation. Such information, though often not adequately addressed in studying the mechanisms of mechanical behavior, is an essential component of knowledge of the microstructural dependence of properties that is necessary to understanding and improving their performance.

The other key microstructural component is porosity, which has been addressed in a companion book [42].

These books show that microstructural understanding, and ultimately design and control of microstructure, is key to good materials. However, this has to be intimately coupled with fabrication and processing technology, which still holds great promise for further development and invention to produce novel microstructures. This book addresses many of these opportunities, a major component of which is composites, but the reader is also referred to some of the earlier discussions of opportunities [43–46] and more recent ones [48,49].

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