Mechanical Properties of Ceramics and Composites

140

Chapter 3

FIGURE 3.4 Continued

Grain Dependence of Ceramic Tensile Strengths at 22°C

141

FIGURE 3.5 Strength of dense MgO versus G-1/2 from various studies, including substantial earlier data of Rice for as-hot-pressed (AS HP), hot pressed and annealed (HP & A) and hot extruded [48, 50–52] (all with as-sanded surfaces); earlier data of Vasilos and colleagues [53, 54] as well as data for single crystals as-cleaved, machined, or recrystallized [47, 48] or translucent-transparent bodies as hot-pressed (with different grinding directions [13] relative to the tensile axis, TA) or hot-pressed and annealed. Also shown are hot pressed data of Bradt et al. [25] (ground or lapped perpendicular to the tensile axis with grits as shown) and Evans and Davidge [55] for transparent MgO (with as-sawn surfaces); and sintered data of Harrison [56] (with various surface finishes), as well as more recent data of Nishida et al. [57].

nificant σ effects of grinding direction, especially at finer G (3–20 µm), and (2) much lower slope (and lower σ) finer G branches. Such finer G branches were further (1) shown by fractography separating edge and tensile surface failures [17] (i.e. similar to Al2O3 studies) and (2) indicated in more recent analysis of MgO data, especially that of MgO with second phases (much of them at grain boundaries) [52]. Similarly Bradt et al.’s [25] tests of hot pressed MgO ground

142

Chapter 3

FIGURE 3.6 Fractographs of recrystallized MgO samples showing fracture initiation associated with slip bands blocked at grain boundaries. (A) Fracture origin from two transgranular fractured grains showing etched slip band across the left grain terminating at the origin, which is internal, as was typical for machined surfaces due to work hardening of the surface grains. (B) Less common origin from grain boundary surface, again with associated slip bands (arrows) and internal location. (From Ref. 52, After Rice [52] published with the permission of the Journal of the American Ceramic Society.)

Grain Dependence of Ceramic Tensile Strengths at 22°C

143

parallel or perpendicular to the tensile axes with various grit sizes again indicate lower strengths with coarser grits, especially at finer G. These all indicate a transition to preexisting flaw control of strength at finer G and more competition between failure from such flaws and microplasticity at larger G, with harsher surface finish and more residual porosity or other phases shifting the balance in favor of flaws.

Extensive fractography of specimens from hot deformed, recrystallized MgO showed fracture origins from slip bands (Fig. 3.6). Some of these were implied based on expected association with cleavage steps [48, 52]. Etching corroborated such cleavage step–slip band association as well as frequent internal fracture origins being due to work hardening of surface grains due to mechanical finishing [49].

III.CUBIC CERAMICS AND RELATED MATERIALS FAILING FROM PREEXISTING FLAWS

A.Cubic Alkaline Earth Halides and Oxides

While most, if not all, dense alkali halides show microplastic controlled strength (and often some macroplasticity), alkaline earth halides have shown no evidence of this at room temperature [58]. Single crystals and large G bodies that have been investigated, mainly CaF2 and some BaF2 and SrF2, have shown three characteristics of failure from preexisting flaws. The first is greater strength sensitivity to surface finishing than expected for microplastic controlled strength, e.g. doubling single crystal and fully dense (press forged) polycrystal strengths [59, 60]. The second is substantial identification of flaws at fracture origins that were consistent with fracture stress and toughness ( 0.4 MPa·m1/2) values [61]. Third is some increase in strengths with limited decreases in G (e.g. to 100–200 µm) as well as some strengths higher and lower than single crystals at large G ( 1 cm in test bars from fusion cast billets with bar dimensions 1/2 cm). Fractography showed clear cases of single crystal fracture mirrors, i.e. mist and hackle formation within a single grain (Fig. 3.7) [61], showing that failure occurred entirely within such a large grain rather than propagating subcritically into surrounding grains.

Limited data exists for some other common single oxide (machined) ceramics that is not inconsistent with the flaw model or provides some support for it. Thus NiO strength data of Spriggs et al. [62] and of Harrison [63] individually are not inconsistent with the flaw model (but are difficult to compare since they used respectively diametral and flexure tests). However, both showed clear deviations to lower strengths at finer grain sizes achieved, e.g. respectively 0.5–1.5 and < 5 µm by higher pressure hot pressing or hot pressing versus sintering. Thus Spriggs et al.’s bodies hot pressed at higher pressures (70–140 MPa) only had

144

Chapter 3

FIGURE 3.7 Optical fractograph of a higher strength, larger grain, dense CaF2 from a surface (probably machining) flaw near the test bar edge, photo bottom center, at start of edge rounding. Note grain boundary near right photo edge and the onset of mist and hackle within the grain of origin.

about 2/3 the strength, 240 MPa (P 0, G 1.5 µm) versus samples hot pressed at lower pressures, despite the latter having larger G, e.g. 3-fold. Such lower strengths with finer G as higher hot pressing pressures were used were also shown by Spriggs et al. [64] for other oxide ceramics, Al2O3, Cr2O3, TiO2, and MgO, as is discussed in Sec. V.A.

ThO2 data of Knudsen [65] illustrates a typical problem with much, particularly earlier, data, namely substantial P, especially at finer G. However, correction to P= 0 gives results reasonably consistent with the flaw model, e.g. suggesting a finer G branch of lower but >0 slope transitioning to the larger G branch at G 10–15 µm (Fig. 3.8). Even more limited uniaxial flexure data [28, 67] for dense, transparent sintered Y2O3 shows decreasing strength with increasing G, as does biaxial flexure data of Rhodes et al. [68]. Both studies clearly show failure initiation in larger G bodies occurring from machining flaws substantially smaller than the grains, even with the size of surface grains typically being substantially reduced by machining (Fig. 3.9A and B). On the other hand,

Grain Dependence of Ceramic Tensile Strengths at 22°C

145

FIGURE 3.8 Strength versus G-1/2 for sintered ThO2 at 22°C of Knudsen corrected for variable porosity (% shown as superscripts) using the exponential relation e-bP, with b 4.4 per his determination [65]. Vertical bars = standard deviations, subscripts = number of tests averaged. Note good consistency between various data points of differing P, as well as reasonable agreement with the one data point of Curtis and Johnson [66]. Also included are data on the same ThO2 at 1000°C [65] and for two grain sizes of dense (transparent)Y2O3 at 22°C [28, 67]. Bars are the standard deviations, and subscripts or numbers adjacent to data points the number of tests.

fractography of the finer grain body showed machining flaw fracture origins somewhat larger than G, indicating they are from the finer G branch, but near its intersection with the larger G branch.

Diametral compression strengths of UO2 by Kennedy and Bandyopadhyay [69] for G 2–20 µm showed a finer G branch with limited, but >0, slope, and probably transitioning to the larger G branch at G 10–20 µm, whether strengths are corrected for the limited P or not. Flexural strengths from three other studies with G 8–50 µm [70–73], again with or without limited corrections for limited P, are more consistent with each other than the diametral strength values, being higher than the latter by 2–3-fold. These studies also indicate possible transitions to a larger G branch in the range of G 10–20 µm, which would be consistent

146

Chapter 3

FIGURE 3.9 Sample fractographs of machining flaw origins in dense (transparent) Y2O3. (A) Smaller, more irregular flaw from machining parallel to the subsequent tensile axis. Note its location in a large (> 300 µm) grain. (B) Elongated flaw from machining perpendicular to the subsequent tensile axis. Note location near the left boundary of a truncated grain. (C) More uniform flaw in a smaller remnant of a larger grain.

Grain Dependence of Ceramic Tensile Strengths at 22°C

147

with strength being controlled by toughness values between those for grain boundaries or single crystals and polycrystalline values.

Previous compilations [74, 75] of the limited data for nearly, or fully, cubic ZrO2 with CaO, MgO, or Y2O3 stabilizers showed a substantial decrease in strength with increasing G over 10 µm, consistent with this data being on the larger G branch as expected (Fig. 3.10A). More recent limited data of Adams et al. [76] (11 wt% Y2O3, G 1–50 µm), though generally of lower strengths (in

A

FIGURE 3.10 Strength and fracture of fully stabilized, machined cubic zirconia bodies. (A) Strength–G-1/2 data for fully and some partially stabilized bodies (see Fig. 3.21 for TZP data), as well as similarly finished, fully stabilized crystals of the same or similar compositions [80]. (B) Optical fractograph of a fracture origin from a smaller surface grain (surrounded by somewhat larger, but still normal, mostly transgranularly fractured, intermediate size grains) and the fracture mirror (brighter region) in an intermediate to larger grain ZrO2 + 1lw/o Y2O3 body. (C) SEM fractograph of a machining flaw fracture origin in a CZ crystal. (D) SEM of fracture origin (center of photo) in a commercial PSZ (Zircoa 1027 with 2.8 w/o, 8.1 m/o, MgO) from an individual grain boundary facet with common excess tetragonal phase and some pores [82]. (Photos (B) and (D) published with the permission of the ASTM [21].)

148

Chapter 3

FIGURE 3.10 Continued.

Grain Dependence of Ceramic Tensile Strengths at 22°C

149

part reflecting no correction for the 1–8%, often heterogeneously distributed porosity, especially at larger G), clearly extends into the finer G branch, e.g. being consistent with an intersection of the branches at G 15 µm. This data and effects of porosity on it are corroborated by data for transparent cubic ZrO2 (+10 m/o TiO2 and 7 m/o Y2O3) of Ahlbom et al. [77] falling along or slightly above the upper bound of Adams et al.’s data. Data of Hague [78] and King and Fuchs [79], having the highest porosity levels (to 15%, but mostly ≤ 8%), is (possibly over) corrected to zero porosity via the exponential relation e-bP, using b = 8, except at the lowest stabilizer contents). Again, note polycrystalline strengths extending substantially below earlier [74] and later, higher quality fully stabilized single crystals of the same or similar compositions (mostly with Y2O3) and surface machining [80]. PSZ polycrystalline data (i.e. for cubic grains with tetragonal precipitates) shows similar larger G dependence, but with higher strengths, as expected from transformation toughening, and thus steeper slopes [81–83]. (Such PSZ polycrystalline strengths are well below those for PSZ crystals of similar compositions and surface finishes, but are consistent with strengths of TZP bodies, Fig. 3.21). Fractography showed fracture initiation in intermediate G Y-CZ from grains or defects on the scale of the typical cubic grains (Fig. 3.10B and C) and fracture mirrors consistent with such grain scale flaws [21]. Grain scale origins from individual grain boundary facets (with some pores) are found in commercial PSZ (Fig. 3.10D), again showing similarities between cubic and PSZ [83].

Extensive MgAl2O4 data of Bailey and Russell [84, 85] and others in an earlier survey [86] and more recent work [87–92] show two sets of σ–G-1/2 branches, with fine G branches being determined by various fabrication, e.g. finishing, parameters (Fig. 3.11). Though there is considerable variability, all fine G branches show >0 slopes. Most if not all the significant variations of σ with stoichiometery were via G [84, 85], e.g. Bailey and Russell’s MgO rich MgAl2O4 showed G decreasing 10·fold with corresponding σ increases with increasing excess MgO [85]. The data clearly shows σ (even for highly translucent to transparent specimens) in the large G region extending down to40–70% that of low index orientation of stoichiometric single crystals with similar surface finishes. This is also true for Gentilman’s [92] data for fusioncast transparent (2Al2O3·1MgO) large (2–5mm) G specimens tested with grain boundaries perpendicular to the tensile surface. Fig. 3.12 shows one set of the author’s data from Fig. 3.11 along with specific fracture origins. Data of Hou and Kriven [93] for CaZrO3 also shows a two-branch behavior. A single data point of Moya et al. [94], is reasonably consistent with their data but probably reflects common differences in preparation, testing, and characterization (especially G measurement).

Note that there is generally no significant inconsistency between the G dependence of toughness and strength for most cubic materials where there is