Mechanical Properties of Ceramics and Composites

Again, within each of these topics the data, though limited, is treated in the same general order for composites in earlier chapters, i.e. glass matrix, polycrystalline oxide, all nonoxide, platelet, whisker, eutectic, and then ceramic–metal composites.

II.THEORETICAL BACKGROUND

There is very little specific theoretical background for the focus of this chapter, particle (and matrix grain) dependence of properties, i.e. again similar to the situation for monolithic ceramics in Chapters 6 and 7, but with less certain guidance for composites. Again the most specific guidance is for elastic moduli, but with three added uncertainties and variations over those for monolithic ceramics, namely the specifics of the combination of the moduli of the constituents to yield those of the composite as at room temperature, possible changes due to chemical interactions of the constituents, and variation in moduli of each constituent with temperature. The latter may vary for the constituents, especially when the refractoriness of the constituents varies widely, e.g. for a glass matrix with refractory dispersed phase or a refractory ceramic matrix with less refractory metal dispersed phase.

There is more guidance for the thermal shock behavior of composites, since its dependence on other properties such as tensile strength, Young’s modulus, and thermal expansion, as well as thermal conductivity in some cases, is generally accepted (see Chap. 6, Sec. V). However, predicting these composite properties from those of the composite constituents introduces an added degree of uncertainty, which is compounded if microstructural mechanisms such as crack branching and especially significant microcracking occur.

Trends for other mechanical properties may be indicated by their dependence on or correlation with elastic moduli, especially Young’s modulus, but this poses the same or more severe uncertainties as for monolithic ceramics. Thus while the trend for fracture toughness would be predicted by that of Young’s modulus for simple elastic fracture, the occurrence of behavior such as microcracking and crack branching or bridging, which are composite, microstructure, test, and crack size and character dependent, present complications. Resultant uncertainties are compounded for behavior that only correlates with elastic properties such as hardness and compressive strength, and the relation between the latter two properties in view of the limited and uncertain theoretical dependence of hardness on composite structure and temperature. Wear and related behavior that typically depend on elastic moduli, hardness, and fracture toughness thus reflect the compounded uncertainties of each of these properties and their temperature dependence in predicting trends, thus typically limiting prediction to only rough trends.

III.DATA REVIEW

A.Elastic and Related Behavior

One of the limited studies of the temperature dependence of elastic moduli of ceramic composites is the experimental and summarized literature data for Young’s modulus of alumina fibers (FP) and Al2O3 +20 w/o Y-ZrO2 (PRD) versus temperature to 1200–1300°C of Lavaste et al. [1]. The FP fiber had G 0.5 m and the PRD 166 fiber respective Al2O3 grain and ZrO2 particle sizes of 0.34 and 0.15m, with both fibers averaging nominally 18 m diameter and having very limited residual porosity. They reported E 360 GPa for the PRD fiber at room temperature, in good agreement with the rule of mixtures expectation, i.e. Eq. (8.1) with E 400 and 220 GPa respectively for Al2O3 and ZrO2. Their E value for PRD fibers was very close to that reported by Romie [2] of 380 GPa at room temperature, but both were higher than the value reported from another study [3] of 280 GPa. However, the latter data and that of Lavaste et al. [1] are much closer at higher temperatures, e.g. being respectively 300 and 325 GPa at 800°C and crossing each other twice between 1000 and 1200°C, giving respective values of 190 and 159 GPa at 1200°C. These values are bracketed (except for the one low room temperature value for the PRD fiber) by values for FP fibers of 405 and 300 GPa at room temperature, with the higher FP value being consistent with bulk alumina data. However, the PRD fiber data showed faster E decrease above 1000–1100°C than the alumina fiber (FP) and bulk alumina [4]. This faster decrease of the Al2O3-ZrO2 (PRD) fiber probably reflects at least in part the finer grain and particle sizes and resultant enhanced grain boundary sliding, and probably faster lower temperature decreases in E (Fig. 6.18).

French et al. [5] compiled Young’s modulus data for both the components

of, as well as for composites of Al2O3-50 v/o ZrO2 and Al2O3-50 v/o Y3Al5O12 (YAG), showing that E for both the composites was very close to the rule of mix-

tures [Eq. (8.1)] to the limit of the data presented (1200°C). Tsukuma et al. [6] showed an overall nearly linear decrease in E for 2Y-ZrO2 + 20 w/o Al2O3 composite to 800°C, except for a temporarily faster decrease between 200 and 300°C (again consistent with Fig. 6.18). The latter corresponded with a marked internal friction peak just past 200°C.

Gadkaree [7] measured Young’s and shear moduli and Poisson’s ratio versus temperature for composites of 25 v/o of SiC whiskers hot pressed in either a cordierite or an anorthite matrix. The cordierite matrix composite gave moduli decreases of 11% to 1100°C and the anorthite matrix composite decreases of5–7% to 1000°C, with indications of accelerating decreases above 1000°C for the former and especially above 1100°C in the latter. The respective changes of Poisson’s were + 4% and – 11%, but they are of some statistical uncertainty. Application of the Halpin–Tsai model for fiber composites gave mixed predictions at room temperature and were apparently not applied to the temperature

1.4 MPa·m1/2 at 22°C and decreasing rapidly above 850°C to 0.25 MPa·m1/2 at 1100°C. These studies, especially that of Govila et al., showed that the increase in toughness (and strength) are due respectively to plastic blunting of cracks and high-temperature SCG.

Beall et al. [13] showed the (CNB) fracture toughness of an extensively crystallized canasite glass decreasing from nearly 5 MPa·m1/2 at 22°C to a pronounced minimum of 1 MPa·m1/2 at 600°C and then increasing to 2 MPa·m1/2 at the limit of testing at 800°C (Fig. 11.1). They attributed the marked decrease in toughness with increasing temperature to decreasing microcracking, which they believed was an important factor in the high toughness at 22°C.

Turning to non-glass-based composites, Brandt et al. [14] reported that the

NB toughness of a commercial cutting tool composite of Al2O3-4 w/o ZrO2 decreased from 4.3 MPa·m1/2 at 22°C by only 10% by 800–1000°C and then by

50% by 1200°C. The latter marked decrease was attributed to the onset of high-temperature slow crack growth based on both the substantial decrease from

50% transgranular fracture at room temperature and especially deviations from linear load deflections at stress intensities > 2 MPa·m1/2 at 1200°C.

French et al. [5] measured both CNB and IF toughness for both the con-

stituents and their composites of Al2O3-50 v/o ZrO2 (CZ) and Al2O3-50 v/o Y3Al5O12 (YAG) to 1200°C, all having 2 m grain and particle sizes. CNB tests showed the Al2O3 toughness decreasing from 3.7 to 2.25 MPa·m1/2 with most of the decrease by 800–1000°C and a similar 60% decrease for CZ from 2 to

1.25 MPa·m1/2, while the composite, though starting lower, decreased less (from

3.1 to 2.5 MPa·m1/2), ending up > the Al2O3. In contrast to this, CNB tests showed YAG toughness independent of temperature at 1.8 MPa·m1/2, but with a modest maximum of 2.2 MPa·m1/2 at 1000°C; the YAG composite toughness was constant at 2.8 MPa·m1/2, except for a modest decrease between 1000 and 1200°C. On the other hand, IF tests typically showed that toughness values for all constituents and composites at 22°C were 40% < CNB values with similar

trends with temperature, except for ZrO2 and Al2O3-50 v/o ZrO2, which showed less decrease with increasing temperature and increases at higher temperatures to

respectively equal and exceeding the toughness of the Al2O3. The authors also showed that their Al2O3 had essentially all intergranular fracture while the cubic ZrO2 and especially YAG had mainly transgranular fracture, and the composites had fracture modes intermediate between the constituents at 22°C. However, all bodies showed essentially exclusive intergranular fracture by 800°C.

Tsukuma et al. [6] reported that the NB toughness of their HIPed 2Y-TZP+

20 w/o Al2O3 decreased from 10 MPa·m1/2 at 22°C to a minimum of 6.5 MPa·m1/2 at 400°C, through a maximum of 8 MPa·m1/2 at 700°C and then decreasing to

6 MPa·m1/2 at 1000°C. Whether the minimum is related to the more rapid decrease in E between 200 and 300°C (Sec. III.A) is not known, but this minimum also appears to correspond to strength minima at 400°C (Fig. 11.10).

Orange et al. [15] showed NB toughness of reaction processed mullite + 20 v/o ZrO2 decreased to a minimum at 400°C and then rose to a maximum at 800°C (Fig. 11.2), similar to the above 2Y-TZP+ 20 w/o Al2O3. Two variations of the composite produced lower toughness values (e.g. to 4.5 MPa·m1/2 at 22°C) but with minima and maxima at the same temperatures as for the highest toughness, as was also true for strengths, i.e. they correlated with toughness trends. However, a much more pronounced NB toughness and a lesser strength minimum, and an implied toughness maximum, at ≥ 200°C higher temperature, was reported by Leriche et al. [16,17] for similar reaction processed bodies. These have some similarities with, and differences from, the more limited strength data on such materials.

Niihara et al. [18] showed that I and IF toughness of their hot pressed Al2O3+ 5 v/o SiC (2 m) composite was constant at 4.3 MPa·m1/2 ( 10% above the Al2O3 without SiC) till 800°C and then increased to 5 MPa·m1/2 at 1200°C (while the Al2O3 decreased to 3 MPa·m1/2). Fractography showed clear SCG on 1200°C fractures.

Data of Baldoni et al. [19] for an Al2O3 -30 v/o TiC composite showed IF toughness decreasing from 3.5 MPa·m1/2 at 22°C by 10–15% by 800–1000°C

FLEXURE TOUGHNESS (MPa-m1/2)

FIGURE 11.2 Fracture toughness versus test temperature of mullite alone and mullite-ZrO2 composites made by reaction of zircon and alumina [15] and for another similar reaction processed composite (along with strength data) [16,17]. Note the significant minima and subsequent maxima but with different temperatures and property values of occurrence.

and then increasing 10% to 4 MPa·m1/2 at 1200°C, the latter associated with SCG. These trends are consistent with those of Brandt et al. [14] for NB toughness of a commercial cutting tool composite of Al2O3-30 w/o Ti(C,N) being constant at 3 MPa·m1/2 to 1000°C, and with intergranular fracture increasing substantially from 50% at 22°C.

Baldoni et al.’s [19] data for a Si3N4 -30 v/o TiC composite showed IF toughness decreasing from 4.7 MPa·m1/2 at 22°C by 15% at 800°C, then increasing to 6.9 MPa·m1/2 at 1200°C, the latter associated with SCG. Tai and Watanabe [20] showed that I toughness of their WC+10 w/o Co-20 w/o Al2O3 composite was twice that of the Al2O3 alone ( 9 versus 4 MPa·m1/2) to 300 °C; it increased to nearly 11 MPa·m1/2 at 400°C and then decreased back to 8 MPa·m1/2 at their limit of testing of 600°C (while that for Al2O3 decreased slightly).

Toughness data for composites of mullite with SiC whiskers showed higher toughness at 1300 versus 22°C along with strength data (Sec. III.E).

C.Thermal Shock

Before considering the effects of temperature on tensile or flexure strength, the subject of effects of transient exposure to higher temperatures, namely thermal shock degradation, is considered. As discussed in Chap. 6, Sec. V, there are two cases to be dealt with, one of avoiding any significant loss of strength as a result of thermal shock exposure, and the other surviving thermal shock with some limited but useful load carrying ability. In the first case it is commonly of interest to start with and maintain substantial strengths, e.g. at least several hundred MPa, while in the latter case starting, and especially postshock, strengths of the order of 10 to 100 MPa are common. Refractories are a common and very important manifestation of the latter.

While the subject of refractories is a large and specialized subject not treatable here, a few comments on the relevance of composite structures are pertinent. High thermal shock resistance is achieved in many refractories by making them quite porous (which also aids their thermal insulation), but there are important cases where extensive porosity is either not sufficient or not suitable, e.g. in steel processing, especially in continuous casting, which places severe thermal shock (and other) demands on refractories. Refractories to meet such demands are commonly made with substantial contents, e.g. 25–40%, of sizeable (e.g. ≥ 100 m) graphite flakes in matrices of such materials as SiC or Al2O3 [21]. Such refractories (often containing 15–30% porosity) have such high thermal shock resistance that use of the thermite reaction of iron oxide and aluminum metal powder to produce molten iron in seconds is needed to test their thermal shock resistance. Such microstructures of coarse graphite flakes in denser refractory carbide matrices or of coarse (e.g. several hundred m dia.) BN flakes in oxide

626 Chapter 11

matrices were also investigated for missile nose tips for military needs [22]. However, more recent work has been heavily focused on much finer composite microstructures in the search for bodies capable of sustaining substantial thermal shock but retaining considerably more strength.

That considerable improvement can be made in thermal shock resistance via composite microstructures was shown by studies of Fairbanks et al. [23] and of Oguma et al. [24] of respectively a cordieriteand a canasite-based glass before and after crystallization. Both showed that crystallization increased the critical quench temperature difference ( TC) for strength loss on quenching into a room temperature water bath twofold (respectively from 130 to 350°C and from 100°C for the parent glass and to 200°C for the crystallized canasite glass), < the threefold increases in strengths, despite some increase in E (and thermal expansion, at least for the canasite system). This doubling of TC was consistent with the strengths divided by the product of the Young’s modulus and the thermal expansion coefficient, especially as shown by more complete characterization of the canasite system. Further, both systems showed ≥ threefold increases in retained strengths for quenching above TC, i.e. from 5–10 to 25–30 MPa, which correlated with toughness increases of twofold, or substantially more, as crack sizes increased in the crystallized bodies. This crack size dependence of toughness was cited as important in the improved shock resistance of the cordierite system [23]. One uncertainty left in the cordierite system is why annealing the parent glass below the nucleation temperature resulted in not only a doubling of the starting

lar to those of strength and fracture energy as a function of the volume fraction of unstabilized ZrO2 (of high uniformity in size and dispersion, e.g. Fig. 9.6). However, these results were from tests in which the quench medium was boiling water, rather than water at room or lower temperature, and boiling water does not give as severe a thermal shock. Thompson and Rawlings [26] also showed some modestly increased thermal shock resistance, i.e. a TC of270°C for Al2O3-20 w/o ZrO2 (+ 3 m/o Y2O3) by quenching into an ethylene glycol–water solution at 20°C, which again represents a less severe shock. They also showed that acoustic emission and other acoustic techniques gave additional information, including on precursor events to thermal shock damage and its sources.

Other tests of similar Al2O3-ZrO2 composites quenched into water at ≤ 22°C do not show such improvements. Thus Tomaszewski [27] showed that TC remained at 200°C with increased addition of unstabilized ZrO2 to Al2O3 despite starting strengths first increasing and then decreasing with increasing ZrO2 content for quench tests into room temperature water. Beyond 20 w/o ZrO2 there

was essentially no change in the low strengths ( 100 MPa) with quenching. Sintering the composites in vacuum, which resulted in substantial stabilization (due

3Y-TZP-20 w/o Al2O3 quenched into room temperature water.

Lutz et al. [29] showed that TC for their composites consisting of agglomerates of one Al2O3-ZrO2 composition designed to give greater microcracking in a matrix of another Al2O3-ZrO2 composition designed to give less microcracking was also typically 200°C when the bodies had strengths > 200 MPa. Composites giving limited or no loss of strength on quenching into room temperature water all had lower starting strengths, typically 100–180 MPa. The lower strength and good retention of it after thermal shock was attributed to extensive microcracking, as shown by dye penetrant tests. Again note that this microcracking only occurred when some limited porosity remained in the bodies, especially in the agglomerates for greater microcracking, and was not present in HIPed bodies (though whether the primary effect is residual porosity or possible changes in ZrO2 stabilization due to some reduction in HIPing cannot be unequivocally ascertained).

Yuan et al. [30] thermally shocked mullite-ZrO2 composite bars by heating them to 1200°C and then removing them from the furnace and exposing them to a stream of room temperature air, giving an estimated cooling rate of 50°C/sec. Postshock measurement showed substantial, 50%, strength loss with no ZrO2 and progressively less loss as ZrO2 content increased, with no loss, or possible increase of, strength at higher v/o ZrO2. Strength losses were generally somewhat greater with three versus one thermal shock (Fig. 11.3). Young’s modulus tests showed some decrease on shocking with no ZrO2, no change with some, and some possible limited increase on shocking at higher ZrO2 contents, again with some greater effects with three versus one shock. Toughness tests after shocking showed some possible limited, e.g. 10%, increases on shocking. Ishitsuka et al. [31], using quenching into a water bath at 0 (not 22)°C, showed TC was 250°C for 3Y-TZP alone and 375°C for mullite alone and 3Y TZP-50 v/o mullite composites, with similar low, e.g. 100 MPa, residual strengths. Composites with higher TZP contents showed promise for possible higher TC and residual strength values, consistent with their higher toughnesses. Though the latter are not fully reflected in strengths (Chap. 9, Sec. III.B), the effect of toughness is consistent with serious thermal shock reflecting larger crack toughnesses values, while strength frequently does not. Orange et al. [15] showed that TC increased from 250+ to 300+°C and retained strengths increased from 75 to 150 MPa as starting strength decreased for water quench tests of their reaction processed mullite + 20 v/o ZrO2 composites. The above

FIGURE 11.3 Strength and Young’s modulus at 22°C of mullite-ZrO2 composites versus volume percent ZrO2. Data [30] plotted for both as-processed bars and after bars were subjected to three air quench thermal shocks. Arrows with the asprocessed E values indicate probable corrections for residual porosities of 2.6–8.7%.

results are consistent with microcracking, which is expected in mullite-ZrO2 composites, and has been directly shown in directionally solidified mullite-ZrO2 composites [32].

Kim and Kim [33] showed that the critical TC for Al2O3-SiC composites made by directed oxidation of Al metal was possibly increased by 10–20% over the value of 200°C for Al2O3 alone, especially as the SiC particle size decreased from 100 to 13 m (which increased starting strengths, Fig. 9.15). However, retained strengths after thermal shock damage were increased by addition of SiC particles, especially with coarser SiC particles, e.g. 100% for 100m particles, i.e. the opposite trend from that for starting strengths. Residual Al

in the composite was seen as not playing a significant role in the composite thermal shock results.

Landon and Thevenot [34] reported that both starting strengths and the critical quench temperature difference for significant strength loss generally decreased in AlN-SiC bodies, e.g. from 800 MPa and 350°C to 450 MPa and300°C as the SiC content decreased from 90 to 10%, with no clear distinctions between whether α or β-SiC was used.

Tiegs and Becher [35] reported that retained strengths in thermal shock tests of Al2O3-20 v/o SiC whiskers progressively increased from 620 MPa with no shock to 700 MPa with a quench of 900°C, but there were slightly reduced retained strengths to 550 MPa with 10 quenches of 900°C. However, these tests were conducted by quenching into boiling instead of room temperature water baths, which as indicated earlier for Al2O3-ZrO2 gives less severe thermal shock results.

Lee and Case [8], while also showing limited effects of additional thermal shocks, showed more changes in thermal shock testing of the same composition whisker composites, with quenching into a room temperature water bath. They evaluated changes in Young’s modulus and internal friction showing respectively decreases and increases as the quench temperature differential increased, with both showing substantial rates of change for quenches of > 310°C (Fig. 11.14). These results indicated limited improvement of thermal shock resistance of such composites over that of Al2O3 itself, e.g. a TC of 200°C (Fig. 6.21). Lack of a precipitous drop in E, as commonly shown for strengths, and a rise in internal friction, is more realistic, e.g. since the catastrophic drop in strength at TC is apparently an artifact of testing limited numbers of specimens and at limited temperature intervals as noted in Chap. 6, Sec. V. The relationship of the above thermal shock behavior to a report of acoustic emissions in similar composites of Al2O3-20 v/o SiC whiskers increasing significantly during cooling from thermal cycling to temperatures > 500°C [36] is uncertain.

Consider now composites of either graphite or BN flakes (or platelets) in various matrices, which have been suggested by development of refractories as noted earlier. A study of the underlying mechanical behavior of Al2O3–graphite refractories has been reported by Cooper et al. [37] and other aspects of laboratory Al2O3–graphite composites were outlined in Chap. 9, Sec. III.C. Composites of various oxides such as BeO, Al2O3, MgO, and ThO2 [22, 38] with additions of 5–30 v/o of large BN flakes (e.g. 75–400 m) were investigated and showed significantly improved thermal shock resistance over the matrix alone but at substantial loss of initial strengths. It was also noted that addition of as little as 3 v/o W (1–2 m) particles to MgO substantially increased thermal shock resistance over the matrix alone, but strength data was again not reported [22].

More extensive study has been made of composites of fine BN platelets in various, especially Al2O3 or mullite, matrices by Rice and colleagues based on