Mechanical Properties of Ceramics and Composites

FIGURE 11.4 Young’s modulus and internal friction versus quench temperature difference for 10 and 100 quenches of Al2O3-20 v/o SiC whisker composites. Note that there is some fatigue effect, i.e. further reduction with repeated quenches at the same temperature difference, but the major effect is from increased quench temperature differences. (From Ref. 8.)

potential needs for improved radar windows and domes for more extreme thermal environments. They first investigated Al2O3-BN composites made by hot pressing powders of the ingredients giving Al2O3 grain sizes ranging from < 1 to > 5 m and BN particle thicknesses of 0.1–0.2 m and diameters of 5 m [39,40]. Subsequent similar processing with mullite matrices resulted in similar microstructures, but probably more toward a 5 m matrix grain size (promising results were also demonstrated with similar Si3N4-BN composites, but they were not pursued because of expected performance and subsequently demonstrated processing advantages of the oxide matrix, especially mullite-based, composites). Properties of these composites are summarized in Table 11.1, showing substantial anisotropy as expected from substantial orientation of the BN platelets during hot pressing, as well as reasonable strengths attributed to the substantially finer BN size than in earlier composites (Fig. 11.5).

Four additional aspects of the above Al2O3- and mullite-BN composite studies should be noted. First, Coblenz [41] conceived of reaction processing

aDensity (ρ ), Young’s modulus (E), fracture toughness (K), flexure strength (σ), and critical temperature difference on quenching into a room temperature water bath for significant loss of strength ( TC); and refer respectively to properties measured with stress parallel and perpendicular to the hot pressing axis, i.e. normal to the plane of preferred orientation. bAl2O3 + 30 v/o BN made by hot pressing powders of mullite and BN with resultant anisotropy. Source: Ref. 39.

cMullite + 30 v/o BN made by hot pressing mullite + BN powders and hence resultant anisotropy. Source: Ref. 39.

dData for mullite + 30 v/o BN made by reaction hot pressing ingredients to produce mullite + BN in situ and hence resultant approximate isotropy. Source: Ref. 41.

eBodies were ~ isotropic (Fig. 11.5C).

that produced more homogeneous and isotropic mullite+ BN composites with lower raw materials costs and similar or easier processing that gave good properties (Table 11.1). Second, better properties and overall performance were found to correlate with more homogeneous, and thus generally finer, microstructures, especially with regard to the matrix grain size (Fig. 11. 5) [39–42].

Third, other compositions have been investigated, e.g. Coblenz and Lewis [41] showed that reaction processing of mullite + 30 v/o BN with 28 v/o residual Si3N4 increased initial strengths by 10–15% and TC by the same or greater percent, e.g. to ≥450°C. Goeuriot-Launay et al. [43] prepared bodies of 70 v/o Al2O3 + 30 v/o γ-ALON with increasing volume percents of BN decreasing starting strengths and increasing TC that are overall consistent with those of Lewis et al. [39] (Fig. 11.6). Mazdiyasni and Ruh [9] showed that Si3N4+ 20 w/o BN hot pressed composites had TC of 700–800°C versus 600°C with no BN and that there were (higher and) more gradual changes in internal friction and less tendency for catastrophic failure in the composite. Valentine et al. [44] showed that SiC-BN composites had increasing TC with increasing BN content similar to, but greater than, that for Al2O3-BN composites (Fig. 11.6).

FIGURE 11.6 Critical quench temperature difference for Al2O3-BN-based composites of Lewis et al. [39] and Goeuriot-Launay et al. [43] and of SiC-BN composites of Valentine et al. [44], and for starting strengths of the former two composites. Note that (1) the upper and lower points show the range of data [39] and the vertical bars the standard deviations [43], and (2) the trends for both properties are overall consistent between the two alumina-based composites, especially for starting strengths (as are also the trends for SiC-BN), and roughly for

TC, more so for the lower values of Lewis et al. [39].

D.Thermal Shock and Related Mechanical Fatigue

The fourth aspect of Al2O3- and mullite-BN composites that should be noted is their thermal shock fatigue, that is, their resistance to degradation and failure under repeated thermal shocking. While this topic has received very little attention, especially for composites designed to give higher strengths than conventional refractories, limited data shows that this is an important subject. Lewis and Rice [45] first investigated this for Al2O3- and mullite-BN composites and showed that just a few repeated thermal shock cycles progressively reduced the retained

while water effects were not the cause of the observed fatigue effect, further testing of materials such as soda lime glasses more susceptible to SCG and to substantially more cycles than the ten they used may be useful to explore possible more limited effects of such SCG.)

Lewis and Rice further showed that various ceramics ranged between the above glass and glass-containing aluminas with no observable to substantial multiple quench fatigue effect. Thus a commercial crystallized cordierite–glass body (Pyroceram 9606 ) showed a modest effect via increased scatter of retained flexural strengths, but no clear trend for strength reduction for multiple

strength. More extreme were results for an Al2O3-25 v/o TZP composite where five quench cycles reduced TC from 250 to ≤ 175°C and net retained strengths by 1/2, i.e. similar to effects with the aluminaand mullite-BN composites. The common theme they saw running through all of these tests was the expected increase in microcracking from none in the borosilicate glass to limited amounts in the PSZ, more in the Lucalox, and much more in the composites with TZP or BN. They thus proposed that the quench fatigue effects arose from increasing extent and effect of microcracking, e.g. via the sequence shown schematically in Fig. 11.8.

In order to determine further the causes of the multiple quench fatigue of ceramics, Lewis and Rice [46] conducted some simple mechanical fatigue tests to ascertain how much of this effect was true thermal shock fatigue and how much was basic mechanical fatigue. Having no specific fatigue testing, they conducted some zero-tension tests using repeated flexure with a conventional mechanical test machine at both 20–25 and -196°C (i.e. the latter in liquid N2) on several ceramics and a granite. Results from two coarser grain materials, Lucalox alumina and the granite, both of which should have some microcracking, showed particular effects with failures at 2–100 cycles at 80–100% of their nominal flexure, i.e. single loading, flexure strength. Both showed failure in fewer cycles at -196°C versus at 20–25°C and that the number of cycles to failure at the nominal flexure strength increased as the thickness of the specimen increased, especially for the granite. The latter was suggested as possibly being due to stress gradients inherent in flexure (and also thermal shock) tests or to specimen compliance changes leading to increased specimen deflections by the test machine on each cycle. (The former may be a function of grain size relative to specimen dimensions, i.e. with more effect in the granite with larger G than the alumina.) Greater fatigue effects at the lower temperature clearly argue against any effect of SCG, since this is greatly, if not totally, suppressed at

FIGURE 11.8 Schematic of possible progressive grain (or particle) boundary microcracking during both tensile loading and unloading to cause zero-tension mechanical fatigue and thermal shock fatigue in multiple quenching. The concept is based on local grain (particle) stress relaxations due to microcrack formation or extension on loading resulting in incompatibilities inhibiting closure of such microcracks on unloading such that further microcrack extension or forming occurs so the specimen locally does not return to the same local stress state on unloading that it was at before loading. (From Ref. 45,46. Published with permission of Ceramic Engineering and Science Proceedings.)

–196°C, while the lower temperature enhances mismatch stresses between grains due to the thermal expansion anisotropy of Al2O3 grains and some of the phases (grains) in granite as well as differential expansion between different phases in the granite.

Based on the above observations, Lewis and Rice [45] proposed that both thermal shock and pure mechanical fatigue resulted from microcracking from mismatch stresses between grains (Fig. 11.8). They recognized that the key to a cyclic fatigue process in a brittle material with macroscopic elastic behavior required a mechanism of progressive damage on a local, e.g. a microstructural, scale with cyclic loading. Some specific scenarios for progressive microcrack extension on each loading cycle were proposed, based on tensile loading forming or extending microcracks such that elastic relaxation of the grains abutting the microcrack result in incompatibilites resisting crack closing on unloading. When such incompatibilities are sufficient to cause further extension or genera-

tion of microcracks during unloading, a basic mechanism exists for progressive degradation on subsequent cycles, since the body locally does not return to the same stress state that existed before the loading and unloading. Such microcrack effects should be sensitive not only to the degree and nature of the mismatch stresses between grains, particles, or both, but also on their grain and particle sizes, as well as on local porosity and boundary phases.

Some other observations have been made of repeated thermal shocking, e.g. of mullite-ZrO2 composites (Fig. 11.3) and of SiC whisker composites. Schneibel et al. [47] investigated the cyclic thermal shock resistance of two commercial Si3N4 and Al2O3-SiC composites (two with whiskers and one each with particles or fibers) using 10 or 100 shocks by quenching bend bar specimens from 1200°C into a room temperature fluidized bed and then measuring flexure strength at 22°C. This procedure, which is less severe than quenching into room temperature water, showed substantial thermal shock fatigue effects in only the Al2O3 whisker and fiber composites.

E.Tensile Strength

Turning now to the temperature dependence of tensile (flexure) strength, consider first glass matrix composites, where limited data is available mainly or only for crystallized glasses. Northover and Groves’ [10] study of LiO2-Al2O3-SiO2 (LAS) bodies showed that fracture toughness and flexural strength both decreased to minima of 2/3 their room temperature values at 500°C and then increased back to, or slightly above, the room temperature values at the limit of their testing of 1000°C (Fig. 11.1). Govila et al. [11] showed similar strength behavior for similar LAS bodies, i.e. a strength minimum at 600°C followed by a sharp maximum at 1000°C, which was similar to that of the parent glass, but with strengths generally twice as high as the minimum, and especially the maximum, extended to higher temperatures (Fig. 11.9). They also demonstrated that the maximum was associated with softening of the residual glass matrix and resultant intergranular SCG in the crystallized body, while subsequent fracture was transgranular.

Borom [48] reported flexure strengths of two Li2O-SiO2-based crystallized glasses decreasing substantially with modest temperature increases. One body decreased from 190 MPa at 22°C to a minimum of 140 MPa at 400°C, i.e. a decrease of 24%, while the other had a greater decrease from 235 MPA to 130 MPa at 550°C, which was believed to be at or near a minimum for it. These decreases of 25 and 45–50% are far greater than the 5% decrease in Young’s moduli. Tests of both parent glasses and a simulated matrix glass for one crystallized body all showed similar strengths at 22°C of 95 MPa that increased with test temperature by 30% at 400°C, i.e. nearly the same as the crystallized bodies at this temperature. This showed that the real

FIGURE 11.9 Strength versus test temperature for the starting Li2O-Al2O3-SiO2 glass and after crystallization. Note the significant increase in strength and reduced dependence on temperature in the crystallized material. (From Ref. 11. Published with permission of the Journal of Materials Science.)

strength decreases of the crystallized bodies was greater due to the increased glass strength, which was attributed to reduced stress corrosion and the onset of plastic flow (which was observed in macroscopic behavior at 450°C).

Consider next Al2O3-ZrO2 composites, starting with sintered-HIPed bodies of 10, 20, or 40 w/o Al2O3 with respectively 90, 80, or 60 w/o 2Y-TZP of Tsukuma et al. [6]. These composites showed similar strengths and trends with test temperature (Fig. 11.10), with a trend for strengths to increase with increasing Al2O3 content, especially at the highest temperature (1000°C). Thus bodies with no Al2O3, i.e. pure 2Y-TZP, had 30% lower strength at 22°C, with the same or greater difference at 1000°C. Greater strengths with Al2O3 versus pure 2Y-TZP was also shown in bodies that were only sintered, but with lower strengths and less difference between bodies with and without Al2O3, e.g. 20% difference. Note that the strength minimum at 400°C corresponds to a toughness minimum at the same temperature, probably with a more rapid decrease in E at 200–300°C, which were noted respectively in Secs. A and B. These in turn are probably related to similar deviations seen in ZrO2 (Fig. 6.18). Results of Govila [49] for HIPed 3Y-TZP with 20 w/o Al2O3, while showing no decrease in strength at 200°C, showed overall similar strength trends with test temperature, as well as a temporarily more rapid strength decrease between 200 and 400°C

FIGURE 11.10 Strengths of Al2O3-ZrO2 composites versus test temperature [1,6,48,49]. Note the similarity between the behavior of the (PRD) fiber [1] and bulk bodies [6,49, 50].

(Fig. 11.10). Results of Shin et al. [50] for HIPed Al2O3-15 v/o ZrO2 showed similar relative strength decrease with increasing test temperature and lower overall strength as expected, while data of Lavaste et al. [1] for PRD166 fibers of Al2O3+ 20 w/o TZP showed similar strength and E decreases (Sec. A) till 1000°C, i.e. 1–20%, then a more rapid decrease (Fig. 11.10). Comparison of the latter two results again showed substantial similarities between fiber and bulk ceramic mechanical property trends, e.g. as discussed in Chap. 3, Sec. III.A. Much less decrease, e.g. 20%, in the high 22°C strengths of 330 MPa to 120022°C for reaction processed composites of mullite-6 w/o Al2O3 + 31 w/o ZrO2 is consistent with their strengths not being derived from transformation toughening but probably in part from finer grain and particle sizes maintained at the elevated temperature [51].

Turning to other polycrystalline composites, Niihara et al. [18] showed that their composite of Al2O3 with 5 v/o SiC (2 m) particles having a strength of 500 MPa at 22°C, and the same or slightly less at 600°C, increased slightly to a modest maximum at 1000°C and then rapidly decreased to < 200 MPa at nearly 1300°C. This maximum and the subsequent rapid decrease in strength correspond with an increasing upward swing of fracture