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interrelations are also probably a major factor in the nonuniqueness of bridging results, which raise key questions for strength and life predictions. Thus much more work is needed in evaluating these and other interactions, which requires more comprehensive tests and characterization to evaluate self-consistency of results. This should entail a broader range of test parameters, e.g. environment, specimen configurations and sizes, surface finishes, loading conditions (e.g. biaxial), with different materials and microstructures especially grain structures. Broader evaluation of fractography (e.g. fracture mode and character), acoustic emission, and crack velocity, as well as properties such as E, damping, and strength is needed. A key aspect of broader evaluation should be evaluation of the toughness–tensile strength relationship, with fractography to determine fracture origins as a key tool.

The primary need is for better perspective, particularly that the crack sizes used in crack propagation–toughness tests are pertinent to the flaw sizes controlling strength, i.e., give similar microstructural effects of strength controlling flaws, e.g. typically from machining. While much work has assumed that large crack test results predict small flaw effects, it has been stressed in this and subsequent chapters that this is not necessarily so; in particular it is only true if the results with test cracks are applicable to strength controlling flaws. This is usually true if similar actual cracks/flaws are used, as in fractographic determinations of crack propagation and toughness behavior, or if the scale of both the test cracks and strength flaws relative to the strength controlling microstructure give similar statistical samplings on both, e.g of grains or other resultant phenomena such as crack deflection, branching, or bridging. This is typically, at least approximately, the case for TZP bodies due to their finer G and is often similarly true for many Si3N4 bodies. The second factor of microstructural impact on flaws controlling strength, especially from machining, is particularly important, since it has been widely neglected, i.e. the focus on understanding strength dependence on microstructure has been sought primarily via its effects on toughness. However, as is shown extensively in Chaps. 3 and 8, body microstructure has an important and often dominant effect on the microstructural dependence of strength, via effects on strength controlling flaws introduced (e.g. as noted in the anisotropy of oriented hot pressed Si3N4), a key factor often lost on a frequently singular focus on toughness.

Preferred orientation of grains, especially elongated ones, though widely neglected, occurs to varying, often significant, extent depending on material and especially fabrication. While much of the resultant effects, such as anisotropy in toughness, are significant, their effects may be greater with intergranular fracture, which is often a function of grain boundary phases. Further, anisotropy of some properties such as toughness and especially strength, while in part determined by properties along appropriate single crystal axes, may often be complicated by other polycrystalline factors such as

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preferred orientation of anisotropically shaped pores. Fractography is thus a critical tool in resolving these, especially strength, effects, as it is in so many other cases.

NOTE

After completing this chapter, a few additional and important references were obtained, whose results pertinent to this chapter are briefly summarized. Donners [268] conducted a more comprehensive study of fracture of MnZn ferrites. He corroborated the opposite change in fracture mode, i.e. from more transgranular for SCG to more intergranular for fast fracture reported in such ferrites reported by Beauchamp and Monroe [52, 53] instead of the usual reverse of this, and cites evidence that this is related to effects of stoichiometry. He also showed that (1) fracture mode could thus vary with location on the fracture surface as a function of distance from the specimen surface due to reduced H2O diffusion, (2) fracture toughness decreased from 2 to 1.3 MPa·m1/2 as the relative humidity increased from 0 – 10 to 100%, (3) NH3 had the greatest effect on SCG, H2S, the next, with NO and CO having a similar effect as H2O, and (4) large, exaggerated grains (e.g. >> 100 µm due to Al2O3 from contact with refractories) were preferred fracture origins. A threshold for SCG was also reported, as was fracture initiation for areas of exaggerated grain growth. He also cited date of Tanaka et al. [269] showing fracture on both (100) and (110) crystal planes with respectively 1–1.3 and 0.9–1 MPa·m1/2, consistent with a preference for (110) fracture.

Swab et al. [270] have shown that fracture initiation in dense, transparent AlON with large G ( 150–200 µm) was from one or a few transgranularly fractured grains followed by mainly intergranular fracture indicating another case of the reversal of the typical SCG to fast fracture mode transition. They also showed that isolated large grains in dense, transparent MgAl2O4 with a bimodal grain distribution (G - 5–20 and 200 µm) were frequent fracture origins, apparently with transgranular fracture, with more intergranular fracture of the surrounding finer grains.

Further testing of electric field effects on fracture of piezoelectric ceramics shows wider and conflicting results. Thus, while some, e.g. Park and Sun [271] showed that positive fields normal to a crack aid propagation while negative ones retard propagation, and questioned the use of stress intensity factors in such cases, Fu and Zhang [272] found crack propagation enhanced by either field polarity, e.g. apparently consistent with effects of ac fields [237]. However, these and other differences, and reports of microcracking and fatigue, respectively implying and showing nonreversible effects, indicate that this is a broader and more complex area of research that requires much more comprehensive study and analysis than it has generally received so far.

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Finally, further note the very high dimensional stability of materials, which consists of two aspects: (1) the precision elastic limit (PEL) commonly defined as the stress to produce a detectable positive residual strain, e.g. 1 ppm, upon removal of a temporarily applied stress and (2) dimensional stability (DS) with a stress applied for extended times, often expressed as a percentage of the PEL, i.e. a “creep” resistance. While the PEL definition is based on stability to 1 part in 10-6 some (e.g. optical and gyroscope) applications, may require higher stability, e.g. 1 part in 10-7–10-8. PEL values for metals are a fraction of their yield stress, e.g. < 1/2, while the more limited values for ceramics may be their ultimate, usually true tensile, strength, but can be less [273–275]. Thus, three measurements showed the PEL of sintered BeO varying from 70–81% of their true tensile strengths (of 105–145 MPa) and one measurement in compression gave a PEL of 82% of the ultimate (1.7 GPa) [273]. The DS of the three tensile tests were 50–80% of their tensile PEL, while the one compressive test was 6–12% of the compressive PEL, i.e. twice the stress level for compression versus tension. Other tests of fused SiO2 and a low expansion, highly crystallized glass (Cer-Vit®) showed elastic behavior to a strain sensitivity of 5x10-8 to stresses of 34 MPa (which was true for specimens with etched or unetched surfaces, with some specimens failing at such a stress) [275]. Such tests may be valuable for detecting effects such as SCG (though the similarity of etched and unetched results with glassy materials may question this) and microcracking, e.g. as may occur in BeO (in tension and possibly compression).

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