Mechanical Properties of Ceramics and Composites

the colony structure. In the above cases the WOF showed little or no decrease and often substantial increase as test temperature increased, as was discussed in Chap. 11, Sec. III.E.

F.Ceramic–Metal Particulate and Wire Composites

Ceramic–metal composites are a subject too large and diverse to address comprehensively here, but there are three sets of such composites that are of value to note here. The first is composites consisting of metal particles dispersed in a ceramic matrix, i.e. analogous to ceramic particulate composites discussed earlier. In such composites, the particles are typically totally isolated from each other and the environment (except for those at the surface), so some of these composites also can compete directly with some all-ceramic materials. The second set are composites in which a minimum amount of metal phase is used for densification and bonding, in particular cermets, especially WC-Co, which besides being an important set of composites and competing with other all-ceramic materials is similar to some of these in terms of microstructure, e.g. of sintered Si3N4. The third set of ceramic–metal composites briefly considered here are those with metal wires, filaments, or fibers in a ceramic matrix. Note that many ceramic–metal composites were investigated before fracture toughness and related evaluation was established, but strengths were evaluated. Thus representative examples of such studies are summarized in Chap. 9, Sec. III.G, which further show that different mechanisms may control strength from those that control toughness, e.g. as shown for monolithic ceramics.

A variety of metal particles, e.g. Al [195], Ag [196], Fe [34,81–83], Ni [197], Mo [198,199], NiAl [200], Pb [201], and W [202], have been introduced in pure or alloy form into polycrystalline, and often silicate glass, matrices. Besides the usual dispersed particle and composite parameters, another factor for some of these (and some all-ceramic composites) is the degree of bonding between the particle and the matrix, which for oxide matrices can often be controlled by oxidation of the particle surfaces, as noted in Sec. III. Thus, for example, Krstic et al. [203] showed that hot pressed composites of silicate glass matrices with (100–200 m) particles of partially oxidized Al or Ni could substantially increase toughness (DCB), e.g. increases > sixfold, from < 1 to > 6 MPa·m1/2 for φ= 0.2. They noted that such increases, which entailed ductile fracture-toughening in the wake region of the crack, required minimizing the elastic and expansion mismatch stresses between the particles and the matrix and having a good bond between the glass and metal (via oxidation of the metal particle surfaces). Substantial fracture examination showed more extensive crack–particle interactions, e.g. crack bowing between particles when the glass matrix had lower expansion than the particles, since the latter then generally act more as pores due to shrinkage away from the matrix. Similar fracture occurred

when the glass expansion was greater than the metal particles and there was not a strong bond between the glass and the particles. Elimination of expansion mismatches, but having elastic mismatches, reduces fracture complexity, i.e. cracks propagated around the particle poles due to stress concentrations there. No significant mismatches of either type resulted in less tortuous fracture in the glass, but random intersections with metal particles allowed more impact of plastic flow with good particle matrix bonding.

Tuan and Brook [197] made composites of Al2O3 with Ni by sintering (in a reducing atmosphere) mixtures of fine NiO particles in the Al2O3 matrix (achieving densities of 98–99% of theoretical for sintering respectively at 1600 and 1700°C at φ=0 and 95 and 92% at φ=0.25). The resulting Ni particles ranged in size (D) from 0.4 to 2–3.3 m as φ increased from 0.01 to 0.25 and sintering increased from 1600 to 1700°C. Toughness (I) increased with both φ and D, with 40–80% of the increases attributed to ductile bridging of the crack in the wake region based on both direct observations of this and their toughness increases being a linear function of (φD)1/2 as predicted by a model developed by Ashby et al. [201]. Tuan and Brooks also observed that residual oxygen in their Ni particles probably substantially raised its yield stress as sintering temperature increased, adding further to the toughening as predicted by Ashby et al.’s model (the yield stress appears with φ and D under the square root). They also showed that the fraction of particles actually intersected by the crack was typically 1/3 the potential interactions, similar to observations of Jessen and colleagues discussed below. Also note that while Krstic et al. [203] did not vary particle size, and their work was done before the model of Ashby et al. was developed, their results appear generally consistent with the model, e.g. as indicated by their high level of toughness increases with larger particles and somewhat higher volume fractions of them.

Breval et al. [204] also made Al2O3-Ni composites, but by hot pressing of powder obtained via Ni(NO3)2 additions to an aluminous sol, resulting in dispersed Ni particles from 20 nm to 50 m, with the sizes not changing significantly as the Ni content was increased from 10 to 50 w/o, but the number of finer particles increased. Although elastic moduli decreased significantly as the amount of Ni increased due to the lower moduli of Ni ( 1/2 that of Al2O3) as well as increasing porosity as the Ni content was increased, fracture toughness more than doubled on going from 0 to 50 w/o Ni, i.e. from 3–4 to 8.3 MPa·m1/2.

Flinn et al. [195] concluded that the toughening mechanism of Al2O3-Al composites made by the directed oxidation of Al, i.e. Lanxide material, is due to ductile bridging of the crack in the wake region and that the model of Ashby et al. is a reasonable guide to the toughening of this material (see Fig. 8.20). However, they concluded that three factors of metal hardening, debonding from the matrix, and the extent of plastic elongation needed further study and documentation. Rice [15,18] also showed that a factor in this material (and others) that

FIGURE 8.20 Fracture of early Lanxide Al2O3-Al composites. SEM fracture photos showing general fracture with a substantial and complex pattern of Al (light material) and its extensive necking down and rough transgranular fracture associated with inclusions in the grain. Such effects are factors in the doubling of toughness. (Photo from Ref. 189.)

needs to be considered is the location of the dispersed phase and its relation to fracture mode in composites with polycrystalline matrices. Thus some inclusions may be incorporated within grains and some at grain boundaries, in which case they only affect the resultant fracture process to the extent that it is transgranular (some of which occurs in Lanxide material, Fig. 8.20) or intergranular.

Jessen and Lewis [81,82] specifically studied the degree of interaction of cracks with dispersed spherical Fe-Ni-Co alloy particles in a glass matrix as occurred in DCB toughness tests. Specimens were fabricated with either uniform volume fractions of 0, 0.1, and 0.25 or with sections of differing volume fraction along the specimens and hence of crack propagation. Tests were conducted so toughness would be measured in various sections of the layered specimens, i.e. in some cases after the crack had propagated through some layers. The result was that the layer through which the crack first propagated determined the subsequently measured toughness value. Thus specimens with the initial section of crack propagation having no dispersed particles always gave the toughness for glass without any metal particles (i.e. 0.7 MPa·m1/2) regardless of whether the KIC value was actually measured in a subsequent layer of the specimen with

metal particles. On the other hand, when the initial section of the specimen had metal particles, the toughness values obtained always reflected that of the initial layer regardless of subsequent layers in which toughness was measured (i.e. where the crack went catastrophic). Thus toughness values were respectively 1.7 and 2.4 MPa·m1/2 for specimens having initial sections of crack propagation respectively with φ= 0.1 and 0.25 regardless of what compositions followed that in which the actual toughness was measured. In other words, the initial material in which the crack propagated determined its subsequent behavior regardless of the subsequent compositions of these composites encountered.

Yun and Choi [202] showed that (I) toughness of composites sintered from AlN with W additions (φ 0–0.2) giving W particles mostly ≤ 1 m (and some reaction phases and reduced AlN G) increased modestly from 1.8 to 2.4 MPa·m1/2 (but strengths decreased from 230 to 200 MPa, which was attributed to some increase in P and second phases as φ increased).

Consider briefly limited results on ceramics with nano scale dispersed metal particles. Nawa et al. made nanocomposites of Mo particles in 3Y-TZP [198] and Al2O3 [199] matrices using Mo powder (0.2 to > 1 m) mixed with comparable sized Al2O3 (α or γ) or TZP powders then hot pressed. Mo (5 v/o) composites using γ-Al2O3 powder gave an α-Al2O3 matrix with considerable growth of the Al2O3 grains (to 8–9 m, and some elongation, e.g. to 15 m), but with some MoO2 and the Mo as elongated layers around parts of the grains and many cracks and branches associated with the grains and Mo resulting in increased (SEPB) toughness of 7.1 MPa·m1/2 (IF did not work on these samples, but no increase in strength was found). Use of α-Al2O3 powder gave much finer alumina grains (e.g. 1/2–3 m) with many nanometer Mo particles within the grains (consistent with the Mo acting as a grain growth inhibitor) as well as some Mo particles > 1 m at grain boundaries, with both the Al2O3 grain and the Mo particle sizes increasing as the hot pressing temperatures increased. Fracture toughness (IF) increased as both the volume fraction Mo increased over the range investigated of 20 v/o as well as with hot pressing temperature, hence with increases of both Al2O3 grain and the Mo particle sizes (but again with strengths showing the opposite trends). While the fracture mode of their Al2O3 was intergranular at finer G and transitioning to transgranular fracture at the larger G, the fracture mode of the composites was predominantly transgranular, even at the finest G, i.e. similar to other composites (Section IV).

Nawa et al.’s 3Y-TZP- Mo composites gave submicron intergranular Mo particles around submicron TZP grains for φ < 0.3, but with further increased Mo content fine (e.g. ≤ 10 nm) intragranular particles increasingly appeared, as did larger, more elongated, often interconnected and polycrystalline intergranular Mo particles. IF toughness tests gave toughness values up to 18 MPa·m1/2 at φ= 0.5, but this was considered uncertain since CNB tests gave a maximum at φ= 0.7 of 11.4 MPa·m1/2 (consistent with a strength maximum of 2 GPa at φ= 0.7).

Sekino et al. [205] made Al2O3-Ni composites (φ= 0–0.2) by either mixing larger NiO particles with fine ( 0.2 m) Al2O3 or by addition of NiO via Ni(NO3)2, with both mixtures reduced to Ni and densified by hot pressing. The two processes producing respectively Ni particles averaging 180 nm (and 1/2 interand intragranular) and 100 nm (with much more intragranular particles) both resulted in modest (I) toughness increases, respectively of 25% and 12% for φ= 0.2. However, strengths were at a significant maximum of 1 and 1.1 GPa respectively for the two processes, with these values being consistent with the reductions of G in the Al2O3 from 3.4 m for no Ni to 2.3 m for 5 v/o Ni from NiO and 0.9 m for Ni from the nitrate.

Chou et al. [200] showed an almost linear increase in NB toughness as additions of NiAl particles (D 5.5 m) increased from 3.7 to 9 MPa·m1/2 as the addition increased from 0 to 50 v/o. These increases were attributed to crack deflection with some contribution of pullout of elongated NiAl particles and some plastic deformation of them.

Next consider ceramic–metal composites consisting of ceramic grains that are partially or completely bonded with a metallic grain boundary phase, usually of limited volume fraction, i.e. cermets, of which the WC-Co is best known and most extensively studied. While much of the extensive development work on this system was done before the development of fracture mechanics and fracture toughness measurements and studies, a few reasonably comprehensive studies addressing WC particle (grain) size have been conducted. Though there is some scatter as expected, these show toughness increasing with both volume fraction Co and the WC particle size (Fig. 8.21). This correlation might suggest ductile deformation of the Co phase as a toughening mechanism per the above ce- ramic–metal particle composite behavior, e.g. since increased Co content and WC size should also correlate with increased dimensions of the Co phase. Though the WC grain/particle size range is not sufficient to test this suggestion seriously, it should be noted that some ductile behavior of the Co bonding phase and resultant toughening has long been indicated. More recent study of Marshall et al. [206] has shown that a zone of irreversible strain exists around cracks, mostly in the wake region, which is a major source of higher toughness and is attributed to plasticity and possibly some microcracking in the WC and a TiC cermet studied. This zone was much larger than plastic zones previously thought or indicated by finite element analysis, e.g. over an order of magnitude larger than the WC grain size. However, it was also noted that the zone size diminished with crack velocity, especially for unstable crack growth, indicating that the dynamic toughness may be substantially lower. This may well be an important factor in why increased toughness as the WC particle size increases has limited or no benefit to strength, which decreases as the WC grain/particle size increases (Fig. 9.18). His observations are probably pertinent to the reaction processed ZrCZrB2 (platelet)- Zr composites of Claar et al. [161] noted in the previous section.

FIGURE 8.21 Fracture toughness of WC-Co versus Co content for various WC grain sizes. (Data from Pickens and Gurland [72,73] and Chermant and Osterstock [76].)

More limited studies of other cermets also provide some evidence for ductile behavior of the bonding phase, e.g. of Cr3C2-Ni and increased toughness as the Ni content increases (but limited decrease in toughness as the content of TiC in lieu of Ni increases) [207]. Han and Mecholsky [208] showed that hot pressed composites with WC-10 w/o Co as the matrix with dispersion of 10 w/o of Nb (17 v/o, D 100 m) or Mo (13 v/o, D 5 or 55 m) particles had increased toughness. Specifically they showed that the absolute value of the toughnesses was proportional to the product of the metal particle yield stress and (Dφ)1/2, while the increment of toughness was proportional to the latter square root term, hence generally consistent with the model of Ashby et al. [201], suggesting plastic deformation as a mechanism.

The above WC-Co and related composites imply higher toughness with a continuous metal phase, as would seem logical and is supported by other data. Thus Trusty and Yeomans [209] showed that hot pressed Al2O3-Fe composites made via wet milling yielding a mainly discontinuous distribution of 20 v/o of Fe doubled the toughness to 6.2 MPa·m1/2 (from 3.1 MPa·m1/2 for the pure

Al2O3), while dry milling that yielded a much more continuous Fe phase tripled toughness to 9.6 MPa·m1/2. Pyzik and Beaman’s [210] reaction processed composites of B4C-Al with an isolated B4C structure generally had higher toughness, e.g. a maximum of 12 MPa·m1/2, versus 8.2 MPa·m1/2 for bodies with a connected B4C structure.

A few brief observations are in order for composites of ceramic matrices with metal wires, fibers, or filaments, since these have been of considerable, especially earlier interest. Originally this interest was motivated by the concept of refractory wires providing toughness and related fracture resistance to ceramic matrices, which in turn would provide environmental protection for the metal wires. This was at least part of the motivation for considerable work on such ceramic matrix composites, e.g. made by consolidating powders around wires, investigation of a number of ceramic–metal eutectic systems, or infiltration of porous ceramic preforms with molten metal. However, subsequent showing that cracks commonly formed in the ceramic matrix which allowed environmental access to the wires and development of other systems greatly reduced work on ceramic–metal wire and related composites. Three examples with some pertinence to the microstructural theme of this book are noted here. Zwissler et al. [36] showed NB toughness of composites of 0–15 v/o chopped stainless steel wires (6, 12, or 25 m dia.) increased linearly by up to 100%, similar to but greater than increases in E or strength, apparently due to the absence of cracking (Sec. III; Chap. 9, Sec. III.G). Simpson and Wasylyshyn [211] showed that the work of fracture of Al2O3 with 0–12 v/o chopped Mo wires (3.16 mm long and 0.08 mm dia.) increased to 250 times that of Al2O3 alone, but strengths were reduced, indicating microcracking. Brennan [212] showed that addition of 25 v/o of Ta wires (0.63 or 1.27 mm dia.) to hot pressed Si3N4, while improving impact resistance and hence toughness, still had extensive cracks from the impact (and poor flexure and especially tensile strength). Second, Flinn et al. [195] showed that ductile bridging of the crack wake region was a major factor in the high fracture resistance with increasing crack extension, e.g. to values of 10–25 MPa·m1/2 in composites with 20 m dia. Al alloy fibers introduced by melt infiltration of a dense Al2O3 preform with aligned channels. Other aspects of ceramic wire or fiber composites pertinent to the issue of fiber sizes and mismatches in properties are briefly discussed in Chap. 9, Sec. III.G.

VI. GENERAL DISCUSSION

A full discussion of the mechanical properties of composites and their microstructural dependence requires more extensive property and microstructural characterization. An important component is other mechanical property, especially tensile strength, and temperature dependence of properties, which are addressed in Chaps. 9–11, so only a partial discussion focused on toughness, based

mostly on large cracks, is presented here. Such cracks often give very different results from smaller cracks, as is noted in Chapters 2, 3, and 9, a key factor being the scale of the crack to that of the microstructural features impacting crack propagation. When the latter are no longer small relative to the crack size, behavior of large cracks may no longer be pertinent to small-scale cracks, i.e. much toughness data with large cracks may not be consistent with strength data reflecting smaller crack behavior, as is shown in Chaps. 3 and 9.

Beginning with macroscopic behavior and progressing to microstructural aspects, consider first the effect of elastic property changes on toughnesses. As noted earlier, since K=(2Eγ) and γE, K E, so increasing E in a composite should increase K. The data for many composites is generally consistent with this, some with quite close correlations, e.g. for BaTiO3 with small amounts of SiC (Fig. 8.19), but there are also important deviations from this, especially for many ceramic–metal composites (Sec. V.F). Except for composites of ceramics with lower E or of metals with higher E, composite moduli are decreased, but toughness is often increased, due to other toughening mechanisms, e.g. plastic deformation, as is discussed below. A specific issue where the composite E and K increase is the extent of K increases versus E values. Becher et al. [164] reported that while composites of two silicate glasses and mullite, or alumina matrices with the same 20 v/o SiC whisker additions, all showed toughness increases, the increases were linear as a function of (EC/EW)1/2, where the subscripts C and W refer to the composite and the whisker, i.e. greater increases as the modulus of the matrix increased. Data of Table 8.3, while providing some possible agreement, also shows some substantial disagreement with this, poorer results in most Si3N4-SiC whisker composites being an important example. This probably reflects other factors such as thermal expansion differences, which Becher et al. also recognized as a source of variation.

Turning to another bulk property correlation, consider ductility of the dispersed phase as in the case of ceramic–metal composites, where both modeling and experiments are reasonably consistent and provide some guidance for improved understanding of mechanisms in other composites. Thus not only do basic models and experiments generally agree, the models predict trends based on material (yield) and microstructure, i.e. on a (φD)1/2 dependence, that have been corroborated by data. Further, clear microstructural observations of ductile deformation of embedded metal particles have been given. However, we need better understanding of particle–matrix bonding variations, including effects of size and expansion differences, and better definition of the effects of yield stresses and yield extent versus properties of the matrix. Experiments with ceramic fillers of some ductility, e.g. halides, may be of interest scientifically and practically, e.g. to lower density and have better dielectric behavior.

Consider now the effects of thermal expansion differences, which while a bulk property difference, manifests itself primarily as a microstructural factor,

moving the discussion to a focus on microstructural issues. A summary of average expansion differences for some key composites is given in Table 8.3, focused on estimates of particle sizes for spontaneous microcrack generation. Though such estimates are uncertain due to probable volume fraction effects and neglect thermal expansion anisotropies of the noncubic constituents as well as elastic anisotropies of all constituents, they indicate that all these composites may exhibit some microcracking. Many composites have not been examined thoroughly for microcracking prior to loading, and especially as a function of loading (e.g. via static modulus decreases and acoustic emission). However, microcracking has been clearly established and studied in SiC-TiB2 composites [213,214], suggesting that it may be more common in a number of other similar composites. Though studied less, microcracking has also been reported in Al2O3- SiC composites by Nakahira et al. [123,124]. In crystallized glasses, differences in thermal expansion (and elastic moduli) may exist between the crystallized particles and the matrix to contribute to microcracking, as may shrinkage strains from the glass–crystal phase transformation strains. Microcracking begins at a minimum particle size and generally increases in effects as D increases (Table 8.3), passing through a maximum and then decreasing in its effects on toughness (e.g. Figs. 2.10, 2.15). However, strength trends often do not reflect toughness trends as measured by large cracks.

Turning to other microstructural aspects of toughening of composites, consider the implications of the microstructural parameters indicated by models summarized in Table 1. Transformation toughening, e.g. in ZTA, is clearly con-

Table 8.3 Summary of Microcracking Parameters for More Investigated Ceramic Composites

Source: Ref. 15, published with permission of Ceramic Engineering and Science Proceedings.

The matrix or major constituent is given first. Δα = the mismatch in polycrystalline thermal expansion coefficients between the composite constituents, T = the approximate temperature below which stress from the expansion mismatch begins to build up, E is the approximate Young’s modulus, γ is the approximate fracture energy for microcracking on the particle scale, and DS is the resulting particle size for spontaneous generation of microcracks (i.e. in the absence of any contribution of external, e.g. applied, stresses).

sistent with the general trends predicted. However, quantitative details are somewhat uncertain, since increased strengths due to reduction of G and surface compressive stresses from machining have not been accounted for in transformation toughening theories. Further, the modeling has generally assumed noninteractive particles, which is often at best uncertain.

Crack deflection has no obvious particle size dependence, and so could be a factor in many composites, as has been suggested by a number of investigations (e.g. Ref. 215). Crack pinning should increase with decreasing D, but probably with some material-dependent particle size, which has been indicated as a factor in some composites. Crack branching may originate from different sources, e.g. from some pinning as well as microcracking. Similarly, crack bridging may be in part related to both of the preceding mechanisms but is a separate mechanism based on its occurrence and effects in the wake region and related R-curve effects, which are typically observed in ceramic composites. Though there are uncertainties in the relation of surface observations of bridging to actual fracture away from the surface (Chap. 2, Sec. II.A), there is no doubt that bridging plays a role in the large crack toughness of many ceramic composites. Again however, there are serious questions of the degree to which such large-scale crack bridging in toughness tests versus effects of such phenomena on strengths that are typically controlled by small cracks with a different propagation history from large cracks for toughness measurements can occur.

Finally consider pullout, which is clearly established as a major factor in the noncatastrophic failure of ceramic fiber composites (where cracks are typically large and probably experience periodic acceleration and deceleration, i.e. similar to large crack toughness tests). Less well established is the extent of pullout mechanisms for whisker and especially platelet composites. While exposure of part of platelet morphology on composite fracture surfaces has led some to suggest pullout as a toughening mechanism in platelet composites, it seems that this partial exposure of platelets may reflect partial microcracking or simple intergranular fracture. Pullout has also been cited as an important toughening mechanism in whisker composites based in part on exposure of parts of whiskers on fracture surfaces. While at least some of the whisker fracture exposure probably again reflects intergranular fracture, there is other evidence to support pullout as a mechanism in whisker composites. A key factor in the latter is the decreased effects of whiskers with increased oxide content on their surface and hence presumably stronger matrix–whisker bonding, since increased matrix–fiber bonding is well known to reduce pullout toughening in fiber composites. While decreasing fiber sizes appear to increase pullout effects (based on surface area), there may be lower limits to this, so the increasing effectiveness of larger whiskers may not be counter to pullout. However, the increase in whisker toughening with increasing Al2O3 matrix grain size [83,162,163,170] is opposite to its effects on strength.