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Toughening mechanisms in the high temperature fracture of high purity alumina
Materials Science and Engineering, A 176 (1994) 455-460
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Toughening mechanisms in the high temperature fracture of high purity alumina T. J. Marrow and S. G. Roberts Department of Materials, Oxford UniversiO', Parks Road, Oxford OXI 3f'H (UK)
Abstract The R curve behaviour of long cracks in high purity (99.98%) polycrystalline alumina was investigated using a single-edge precracked beam, at room temperature, 900 °C and 1200 °C. The room temperature R curve was negligible. A strong R curve was observed at 900 °C and 1200 °C. This was associated with a change in slow crack growth mode from mixed intergranular fracture and transgranular cleavage to fully intergranular fracture. Scanning acoustic microscopy conclusively demonstrated asperity contact in the crack wake at high temperature. It is proposed that the high temperature R curve behaviour is due to crack shielding from friction in the crack wake.
I. Introduction
Polycrystalline aluminas are chemically inert and oxidation and abrasion resistant [1[. Except for debased aluminas, which contain a significant glass content (95%-97%), they generally possess excellent compressive strength and creep resistance up to high temperatures [2]. However, structural applications at all temperatures are limited by their low fracture toughness and consequent poor flaw tolerance under tensile loading I3]. The strong R curve behaviour of some microstructures, i.e. rising fracture toughness with crack extension, leads to improved flaw tolerance [4]. Understanding the fracture mechanisms which control the R curve will aid microstructural design of stronger ceramic components. This is particularly important for high temperature applications where fracture strengths can be extremely low [5, 6]. Room temperature R curve behaviour in ceramics, which has been an area of considerable research interest, is generally attributed to crack bridging or a microcrack process zone, depending on the microstructure. The elastic modulus decrease and dilation from microcracking in the crack wake can shield the crack tip from the applied stress [7-9]. Models also consider the additional work of fracture required to extend both the crack and its wake during steady state growth [10]. Any contribution to shielding by the microcracking ahead of the crack tip is negated in alumina by the decreased intrinsic fracture toughness of the microcracked microstructure [711.The effectivness of microcrack shielding depends on the critical stress required to nucleate a microcrack. In alumina this is controlled by the grain boundary fracture stress 19, 11, 121. This 0921-5093/94/$7.~)(i ,',iS'l)l 0921-5093( 93 )02526-9
varies with the residual stress across grain boundaries which arises from anisotropic thermal contraction of adjacent grains on cooling from the processing temperature [9]. The grain boundary fracture stress increases with decreasing grain size. Predicted fracture stresses in fine-grained aluminas are very high, and room temperature microcrack toughening is not considered to be important in aluminas with grain sizes below 20 /~m-30 ,urn [9. 13, 14]. Room temperature microcrack toughening is an important toughening mechanism in coarse-grained aluminas [9]. Room temperature crack bridging in the crack wake and R curve behaviour have been observed in finegrained aluminas with grain sizes between 4 um and 2(1 /~m[15-17]. It has been proposed that local compressive residual stresses prevent the fracture of suitably oriented grains [ 181. The intact grains, bypassed by the main crack, bridge the crack wake and shield the crack tip by resisting grain pullout and rotation [17, 181. The magnitude of crack shielding, and thus the effective fracture toughness, increases with increasing crack wake length. Bridging zones several millimetres in length have been observed [16, 17!. Grains cannot bridge the crack when the crack opening displacement exceeds approximately one quarter of the bridging grain size [1 91. The extent of the crack bridging zone is thus a function of crack profile [151, and the crack opening displacement and R curve are interdependent. A "material" R curve, independent of specimen geometry, does not exist [201. There is considerable interest in thc development of theoretical models for crack face traction forces. These may be used to predict the effects of crack size, grain size and specimen geometry on the effective fracture toughness [21, © I cJcJ3 - ~lsl.'\,icr Sccluoi~l. :\11 i-ighl'~ r e s e r v e d
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22]. Accurate model development requires a detailed understanding of the distribution and strength of bridges in the crack wake. High temperature fracture and deformation of alumina is dominated by grain boundary sliding, which occurs to relax the applied stress. Grain boundary sliding at low homologous temperatures, which cannot be accommodated by grain deformation, nucleates intergranular wedge cracks at grain triple junctions [23]. Under compressive loading this leads to stable microcracking and elastic creep [24, 25]. Under tensile loading intergranular slow crack growth can occur by the linking of wedge cracks [24]. The crack velocity depends on the grain boundary strength and the rate and mechanism of grain boundary sliding, which is related to grain boundary structure [26]. High temperature grain boundary failure is favoured by the presence of an amorphous grain boundary layer [24, 27, 28]. R curve behaviour has been observed previously at high temperature in aluminas with a substantial glass content (97%-99.5% purity) [29-31]. This was demonstrated to be caused by processes in the crack wake [30] and has been attributed to crack face adhesion by viscous glass [31] or friction from interlocking grains [30]. In this work we investigate the high temperature R curve behaviour of a high purity (99.98% purity), polycrystalline fine grained alumina in which glass adhesion should not be an important mechanism. The small glass content weakens grain boundaries at high temperatures to allow grain boundary sliding. This may induce microcracking and the substantial dilation of microcracks by grain boundary sliding, compared with elastic relaxation of grain boundary stresses of room temperature microcracking [7], may induce considerable crack shielding. Alternatively, crack deflection during intergranular slow crack growth, also encouraged by grain boundary sliding, may form crack bridges. High resolution reflection scanning acoustic microscopy (SAM) is employed here to identify whether microcracking or crack bridging is the dominant toughening mechanism at high temperature. SAM images are particularly sensitive to the presence of surface-breaking cracks [32, 33].
Hightemperaturefracture of alumina
Ca, balance A1). An amorphous grain boundary layer has been observed in a similar high purity alumina [34]. The microstructure was revealed by thermal etching (1450 °C in air for 1 h) of mechanically polished surfaces (Fig. 1). The grain size was approximately 3/~m-5/~m, although some elongated grains up to 10 ktm in size were observed. The presence of glass was confirmed by etching in hydrofluoric acid at room temperature, selectively attacking regions of high glass content. The grain boundaries were lightly etched in some areas. 2.2. Fracture toughness testing
Fracture toughness measurements used single-edge precracked beams (nominally 2.5 mm by 2.5 mm by 25 mm) in four-point bending (outer span, 22 mm; inner span, 12 mm). Through-thickness, straight, planar precracks were produced via coalescence of radial cracks from a line of indentations, made at 0.5 mm intervals with a Vickers' diamond under a load of 20 kgf. This specimen geometry provides a wellcharacterized applied stress intensity factor which is insensitive to small changes in crack length. Substantial slow crack growth is thus encouraged, allowing the R curve behaviour of long cracks to be determined. All fracture toughness tests were conducted in air using a high purity alumina jig in a high temperature furnace (maximum temperature 1500 °C), constructed by Severn Science, which was mounted on a digitally controlled, servo-electric Instron mechanical testing machine. Tests were performed under computer control with simultaneous load and cross-head disNace-
2. Experimental details and results 2.1. Material
The alumina, obtained from Vesuvius Zyalons and manufactured by hot isostatic pressing at 1600 °C for 1 h, had a nominal purity of 99.98%. The principal impurities, silicon, sodium and potassium, are all glassforming elements and a thin amorphous grain boundary layer was expected (manufacturer's analysis: 50 ppm Si, 40 ppm K, 12 ppm Na, 6 ppm Fe, 1 ppm
Fig. 1. Polycrystalline alumina microstructure (thermally etched).
T. J. Marrow, S. G. Roberts
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High temperamre fracttlre qf alumina
ment data logging using either load (0.01N m i n - ~ - 2 0 N min ~) or displacement (0.001 mm m i n - ~ - I mm rain t) feedback control. Tests were conducted at room temperature (approximately 20 °C), 900 °C and 1200 °C. Specimens annealed at 900 °C and 1200 °C after precracking were also tested at room temperature. Fracture surfaces were gold sputtered and examined using scanning electron microscopy (SEM) to determine the crack length at fracture and the extent of preceding slow crack growth. Tests which did not exhibit a satisfactory crack were discarded.
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ture at 900 °C showed a small amount of transgranular cleavage, but was mostly intergranular. Fast fracture at 1200 °C was fully intergranular. The crack front at the slow crack growth-fast fracture transition was indi-
2.3. Fracture toughness results The room temperature fracture toughness was approximately constant (average fracture toughness 5 MPa m ~!2)for crack extensions of up to 0.5 mm, with no measurable R curve behaviour (Fig. 2). The average precrack length was 0.35 mm. SEM examination showed a region of mixed intergranular and transgranular fracture (Fig. 3 and Fig. 4(a)) between the transgranular precrack and the transgranular fast fracture region (Fig. 4(b)). Decreasing loading rate increased the extent of the intergranular-transgranular region by slow crack growth during the test. Crack growth velocities were not determined. Annealing at 900 °C and 1200 °C had no effect on the fracture toughness. Increasing slow crack growth with decreasing loading rate was also observed at 900 °C and 1200 °C. The slow crack growth paths at 900 °C and 1200 °C were completely intergranular (Fig. 5). The fast frac-
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Fig. 3. A precracked fracture toughness specimen broken at room temperature, illustrating precrack (region a), slow crack growth re, on (region b) and fast fracture (region c).
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02 {}4 06 os l Crack extension (Aa, mm) Fig. 2. The effect of crack extension Aa, via slow crack growth during the fracture test, on the effective fracture toughness K at room temperature (25 °C, []), 900 °C ( A ) and 1200 °C (n).
Fig. 4. Room temperature fracture surfaces: (a) slow crack growth by intergranular fracture and transgranular cleavage: (b) fast fracture by transgranular cleavage.
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High temperature fracture of alumina
Fig. 5. High temperature fracture surface of slow crack growth by fully intergranular fracture ( 1200 °C).
cated by a change in the macroscopic fracture plane. R curve behaviour was observed at both 900 °C and 1200 °C (Fig. 2). Effective fracture toughness values of up to 9.5 MPa m 1/2 were observed, whereas the intrinsic fracture toughness value at 1200 °C was less than 3 MPa m ~/2. 2.4. Scanning acoustic microscopy Precracked samples were stressed at 900 °C and 1200 °C, at an applied stress intensity factor of approximately 3 MPa m l/2, until relaxation was observed by a change in the cross-head displacement under constant load. After grinding and careful polishing the crack profiles were examined using a Leitz ELSAM high resolution reflection scanning acoustic microscope before thermal etching for SEM examination [35]. An intergranular stable crack growth region was observed at both temperatures which showed discrete zones along the crack path of good acoustic transmission of surface Rayleigh waves. These were indicated by the absence of fringes parallel to the crack image in the scanning acoustic microscope (Fig. 6). Fringes require good reflection and poor transmission of Rayleigh waves [33, 35]. The corresponding features were identified by SEM (Fig. 7). Transmitting regions
Fig. 6. Scanning acoustic microscope image of the crack profile produced by slow crack growth at 1200 °C. Regions without fringes indicate high Rayleigh wave transmission across the crack. Positions 1 and 2 correspond to Figs. 7(a) and 7(b).
were not observed along the room temperature transgranular precrack, which was strongly reflecting. There was no evidence of crack blunting and microcracking was not observed in either specimen.
3. Discussion The average room temperature fracture toughness value (5 MPa) is slightly higher than average values reported in the literature for fine aluminas [30, 36] (4 MPa ml/2-4.5 MPa m~/2). Negligible R curve behaviour has been observed previously in fine-grained aluminas [37, 38]. The grain size is below the theoretical size limit for microcrack toughening [9, 13, 14], which is confirmed by the absence of observable microcracking. The high fraction of transgranular
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High temperature fractlwe (~['alumina
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Fig. 7. Scanning electron microscope images of asperity contact at the areas of high Rayleigh wave transmission observed in Fig. 6. The crack path was produced by slow crack growth at 1200 °C. (a) and (b) correspond to positions 1 and 2 respectively in Fig. 6.
cleavage during the environmentally assisted slow crack growth may discourage the formation of crack bridges. A bridging zone might also saturate at a small size for a long crack in a fine-grained alumina beam specimen. For a small fraction of bridging grains, a systematic examination using repeated sectioning of the crack profile would be required to confirm the presence of room temperature crack bridging in this particular combination of material and test specimen geometry. Crack deflection by elongated grains may also contribute to the higher fracture toughness value 139]. The high temperature tests, performed over a range of strain rates, fall on a single curve of effective fracture toughness as a function of crack growth. The extent of crack growth increased with decreasing loading rate. This is interpreted as R curve behaviour via slow crack growth, which is strain rate controlled, and a fast fracture mechanism which is strain rate insensitive over these test conditions, lntergranular slow crack growth at both 900 °C and 1200 °C is considered to be driven by grain boundary sliding [24], assisted by the presence of an amorphous grain boundary glass. Crack blunting, which was not observed in these tests, may occur at very low loading rates in some specimen geometries [40]. A study is currently in progress to confirm the presence of grain boundary glass using high resolution
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transmission electron microscopy. Grain boundary glass has been observed in a similar alumina [34]. Fractographic examination using SEM showed no evidence to support adhesion of intergranular fracture surfaces (Fig. 4(a)), as expected from the low content of glass forming elements. Crack face adhesion by viscous glass is therefore not considered to be responsible for the R curve behaviour. The zones of high acoustic transmission, observed using SAM (Fig. 6), in the high temperature slow crack growth region could be correlated directly with SEM observations which suggest asperity interlocking and frictional contact (Fig. 7). Physical contact across the crack would increase transmission and decrease reflection of surface Rayleigh waves [35]. The high temperature slow growth crack path was always intergranular and crack bridging by intact grains was not observed. It is proposed that the convoluted crack path of intergranular slow crack growth causes crack bridging via asperity contact and interlocking, rather than by bridge creation from the bypassing of unfawmrably oriented grains by the main crack. Microcracking was not observed by SAM or SEM at any temperature. Although closed microcracks would be difficult to image using SAM, owing to their high transmission and poor reflection [35], they would necessarily have a negligible shielding effect [7]. The similarity of the R curves at 900 °C and 1200 °C (Fig. 2) may arise from similarities in fracture mode, crack length and specimen geometry. However, the transgranular cleavage during fast fracture at 90t) °C suggests that the intrinsic fracture toughness at 900 °C may be higher than at 121)0 °C and this similarity of R curves may be just coincidental. Further tests would be required to determine accurately the absolute effect of crack bridging by contact in the crack wake on the effective fracture toughness.
4. Conclusion
Fine-grained high purity polycrystalline alumina does not possess a strong R curve behaviour at room temperature for long cracks propagating by environmentally assisted slow crack growth. Room temperature slow crack growth occurs by both intergranular fracture and transgranular cleavage. Microcrack shielding does not occur and the effect of any crack bridging is small. Strong R curve behaviour is induced at 900 °C and 1200 °C when crack extension occurs via fully intergranular slow crack growth. Intergranular crack propagation may be driven by non-accommodated grain boundary sliding, assisted by the deformation of a thin amorphous grain boundary layer. The intrinsic fracture
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toughness decreases at high temperature, accompanied by increased intergranular fast fracture. High resolution reflection S A M demonstrated that crack bridging occurred at high temperatures via the interlocking and frictional contact of asperities in the crack profile. Scanning electron microscope examination alone did not allow their positive identification as crack bridges. T h e R curve arises from traction behind the crack tip which increases with increasing crack length. T h e r e was no evidence to support the alternative high temperature toughening mechanisms of crack face adhesion or microcrack shielding. This combined analysis using S A M and SEM may be of considerable use in the study of R curve mechanisms in other materials, such as platelet-reinforced ceramics and composites, where a detailed understanding of crack bridge distribution is required. Optimizing the crack bridging mechanism may produce microstructures possessing high levels of flaw tolerance. Strong R curve behaviour is required if the decrease in fracture strength with increasing flaw size is to be offset by the increased effective fracture toughness.
Acknowledgments T h e authors would like to thank Professor R. J. Brook for provision of laboratory facilities and Dr. G. A. D. Briggs for the use of the scanning acoustic microscope and for helpful discussions. Thanks are also due to V. Luprano and Dr. C. Lawrence. T h e research was funded by SERC Grant G R G 2853.
References 1 J. B. Wachtman (ed.), Structural Ceramics, Academic Press, London, 1989. 2 J. Lankford, J. Mater. Sci., 16(1981) 1567-1578. 3 J.P. Singh, Adv. Ceram. Mater., 3(1988) 18-27. 4 S. J. Bennison and B. R. Lawn, J. Mater. Sci., 24 (1989) 3169-3175. 5 S.M. Oh, J. Mater. Sci., 12(1977) 405-406. 6 H. P. Kirchner and R. M. Gruvner, Mater. Sci. Eng., 13 (1974) 63-69. 7 A. G. Evans and K. J. Faber, J. Am. Ceram. Soc., 67(1984) 255-260. 8 P. G. Charalambides and R. M. McMeeking, J. Am. Ceram. Soc., 71 (1988) 465-472. 7 A. G. Evans and Y. Fu, in W. D. Kingery (ed.), Advances in Ceramics 10, American Ceramics Society, 1983, pp. 697-719.
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10 B. Budiansky, J. W. Hutchinson and J. C. Lambropoulds, Int. J. Solids Struc., 19 (1983) 337-355. 11 J. Lankford, J. Mater. Sci., 13 (1978) 351-367. 12 A. Krell and G. Kirchoff, J. Mater. Sci. Lett., 4 (1984) 1524-1526. 13 N. Claussen, B. Mussier and M. V. Swain, J. Am. Ceram. Soc., 65(1982) C14-C16. 14 R. W. Rice, S. W. Freiman and P. E Becher, J. Am. Ceram. Soc., 64(1981) 345-350. 15 J. R6del, J. E Kelly and B. R. Lawn, J. Am. Ceram. Soc., 73 (1990) 3313-3318. 16 P. L. Swanson, C. J. Fairbanks, B. R. Lawn, Y. W. Mal and B. J. Hockey, J. Am. Ceram. Soc., 70(1987) 279-289. 17 G. Vekenis, M. E Ashby and P. W. R. Beaumont, Acta Metall. Mater., 38(1990) 1151-1162. 18 M.V. Swain, J. Mater. Sci. Lett., 5(1986) 1313-1315. 19 R. W. Steinbrech, A. Reichl and W. Schaarw~ichter, J. Am. Ceram. Soc., 73(1990) 2009-2015. 20 R. W. Steinbrech, E Deuerler, A. Reichl and W. Schaarw~ichter, in D. Taylor (ed.), Proc. Conf. on Science of Ceramics 14, Institute of Ceramics, 1988, pp. 659-664. 21 X.Z. Hu, E. H. Lutz and M. V. Swain, J. Am. Ceram. Soc., 74 (1991) 1828-1832. 22 J. Llorca and R. W. Steinbrech, J. Mater. Sci., 26 (1991) 6383-6390. 23 A.H. Heuer, J. Am. Ceram. Soc., 63 (1980) 53-58. 24 A.G. Evans, Acta Metall., 28 (1980) 1155-1163. 25 A. Venkateswaren, K. Donaldson and D. P. H. Hasselman, J. Am. Ceram. Soc., 71 (1988) 565-567. 26 R. Raj and M. F. Ashby, Metall. Trans. A, 2 (1971) 1113-1127. 27 P.F. Becher, J. Mater. Sci., 6 ( 1971 ) 275-280. 28 W. R. Cannon and T. G. Langdon, J. Mater. Sci., 18 (1983)
1-50.
29 H. Weininger, K. Kromp and R. E Pabst, J. Mater. Sci., 22 (1987) 1352-1358. 30 R. E. Grimes, G. P. Keikar, L. Guazzone and K. W. White, J. Am. Ceram. Soc., 74(1991)1399-1404. 31 A. Bornhauser, K. Kromp and P. E Pabst, J. Mater. Sci., 20 (1985) 2586-2596. 32 A. Quinten and W. Arnold, Mater. Sci. Eng. A, 122 (1989) 15-19. 33 G. A. D. Briggs, Acoustic Microscopy, Monographs on the Physics and Chemistry of Materials, Vol. 47, Oxford University Press, London, 1992. 34 M. P. Harmer, in W. D. Kingery (ed.), Advances in Ceramics 10, American Ceramics Society, 1983, pp. 679-696. 35 T.J. Marrow, V. Luprano and S. G. Roberts, J. Am. Ceram. Soc., 76 (11)(1993)in press. 36 B. Mussier, M. V. Swain and N. Claussen, J. Am. Ceram. Soc., 65 ( 1982 ) 566-572. 37 K. Ikeda and H. Igaki, J. Am. Ceram. Soc., 70 (1987) C29-C30. 38 H. Weininger and K. Kromp, J. Mater. Sci., 21 (1986) 411-418. 39 K. T. Faber and A. G. Evans, Acta Metall., 31 (1983) 565-576. 40 W. Blumenthal and A. G. Evans, J. Am. Ceram. Soc., 67 (1984) 751-759.
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Toughening mechanisms in the high temperature fracture of high purity alumina T. J. Marrow and S. G. Roberts Department of Materials, Oxford UniversiO', Parks Road, Oxford OXI 3f'H (UK)
Abstract The R curve behaviour of long cracks in high purity (99.98%) polycrystalline alumina was investigated using a single-edge precracked beam, at room temperature, 900 °C and 1200 °C. The room temperature R curve was negligible. A strong R curve was observed at 900 °C and 1200 °C. This was associated with a change in slow crack growth mode from mixed intergranular fracture and transgranular cleavage to fully intergranular fracture. Scanning acoustic microscopy conclusively demonstrated asperity contact in the crack wake at high temperature. It is proposed that the high temperature R curve behaviour is due to crack shielding from friction in the crack wake.
I. Introduction
Polycrystalline aluminas are chemically inert and oxidation and abrasion resistant [1[. Except for debased aluminas, which contain a significant glass content (95%-97%), they generally possess excellent compressive strength and creep resistance up to high temperatures [2]. However, structural applications at all temperatures are limited by their low fracture toughness and consequent poor flaw tolerance under tensile loading I3]. The strong R curve behaviour of some microstructures, i.e. rising fracture toughness with crack extension, leads to improved flaw tolerance [4]. Understanding the fracture mechanisms which control the R curve will aid microstructural design of stronger ceramic components. This is particularly important for high temperature applications where fracture strengths can be extremely low [5, 6]. Room temperature R curve behaviour in ceramics, which has been an area of considerable research interest, is generally attributed to crack bridging or a microcrack process zone, depending on the microstructure. The elastic modulus decrease and dilation from microcracking in the crack wake can shield the crack tip from the applied stress [7-9]. Models also consider the additional work of fracture required to extend both the crack and its wake during steady state growth [10]. Any contribution to shielding by the microcracking ahead of the crack tip is negated in alumina by the decreased intrinsic fracture toughness of the microcracked microstructure [711.The effectivness of microcrack shielding depends on the critical stress required to nucleate a microcrack. In alumina this is controlled by the grain boundary fracture stress 19, 11, 121. This 0921-5093/94/$7.~)(i ,',iS'l)l 0921-5093( 93 )02526-9
varies with the residual stress across grain boundaries which arises from anisotropic thermal contraction of adjacent grains on cooling from the processing temperature [9]. The grain boundary fracture stress increases with decreasing grain size. Predicted fracture stresses in fine-grained aluminas are very high, and room temperature microcrack toughening is not considered to be important in aluminas with grain sizes below 20 /~m-30 ,urn [9. 13, 14]. Room temperature microcrack toughening is an important toughening mechanism in coarse-grained aluminas [9]. Room temperature crack bridging in the crack wake and R curve behaviour have been observed in finegrained aluminas with grain sizes between 4 um and 2(1 /~m[15-17]. It has been proposed that local compressive residual stresses prevent the fracture of suitably oriented grains [ 181. The intact grains, bypassed by the main crack, bridge the crack wake and shield the crack tip by resisting grain pullout and rotation [17, 181. The magnitude of crack shielding, and thus the effective fracture toughness, increases with increasing crack wake length. Bridging zones several millimetres in length have been observed [16, 17!. Grains cannot bridge the crack when the crack opening displacement exceeds approximately one quarter of the bridging grain size [1 91. The extent of the crack bridging zone is thus a function of crack profile [151, and the crack opening displacement and R curve are interdependent. A "material" R curve, independent of specimen geometry, does not exist [201. There is considerable interest in thc development of theoretical models for crack face traction forces. These may be used to predict the effects of crack size, grain size and specimen geometry on the effective fracture toughness [21, © I cJcJ3 - ~lsl.'\,icr Sccluoi~l. :\11 i-ighl'~ r e s e r v e d
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22]. Accurate model development requires a detailed understanding of the distribution and strength of bridges in the crack wake. High temperature fracture and deformation of alumina is dominated by grain boundary sliding, which occurs to relax the applied stress. Grain boundary sliding at low homologous temperatures, which cannot be accommodated by grain deformation, nucleates intergranular wedge cracks at grain triple junctions [23]. Under compressive loading this leads to stable microcracking and elastic creep [24, 25]. Under tensile loading intergranular slow crack growth can occur by the linking of wedge cracks [24]. The crack velocity depends on the grain boundary strength and the rate and mechanism of grain boundary sliding, which is related to grain boundary structure [26]. High temperature grain boundary failure is favoured by the presence of an amorphous grain boundary layer [24, 27, 28]. R curve behaviour has been observed previously at high temperature in aluminas with a substantial glass content (97%-99.5% purity) [29-31]. This was demonstrated to be caused by processes in the crack wake [30] and has been attributed to crack face adhesion by viscous glass [31] or friction from interlocking grains [30]. In this work we investigate the high temperature R curve behaviour of a high purity (99.98% purity), polycrystalline fine grained alumina in which glass adhesion should not be an important mechanism. The small glass content weakens grain boundaries at high temperatures to allow grain boundary sliding. This may induce microcracking and the substantial dilation of microcracks by grain boundary sliding, compared with elastic relaxation of grain boundary stresses of room temperature microcracking [7], may induce considerable crack shielding. Alternatively, crack deflection during intergranular slow crack growth, also encouraged by grain boundary sliding, may form crack bridges. High resolution reflection scanning acoustic microscopy (SAM) is employed here to identify whether microcracking or crack bridging is the dominant toughening mechanism at high temperature. SAM images are particularly sensitive to the presence of surface-breaking cracks [32, 33].
Hightemperaturefracture of alumina
Ca, balance A1). An amorphous grain boundary layer has been observed in a similar high purity alumina [34]. The microstructure was revealed by thermal etching (1450 °C in air for 1 h) of mechanically polished surfaces (Fig. 1). The grain size was approximately 3/~m-5/~m, although some elongated grains up to 10 ktm in size were observed. The presence of glass was confirmed by etching in hydrofluoric acid at room temperature, selectively attacking regions of high glass content. The grain boundaries were lightly etched in some areas. 2.2. Fracture toughness testing
Fracture toughness measurements used single-edge precracked beams (nominally 2.5 mm by 2.5 mm by 25 mm) in four-point bending (outer span, 22 mm; inner span, 12 mm). Through-thickness, straight, planar precracks were produced via coalescence of radial cracks from a line of indentations, made at 0.5 mm intervals with a Vickers' diamond under a load of 20 kgf. This specimen geometry provides a wellcharacterized applied stress intensity factor which is insensitive to small changes in crack length. Substantial slow crack growth is thus encouraged, allowing the R curve behaviour of long cracks to be determined. All fracture toughness tests were conducted in air using a high purity alumina jig in a high temperature furnace (maximum temperature 1500 °C), constructed by Severn Science, which was mounted on a digitally controlled, servo-electric Instron mechanical testing machine. Tests were performed under computer control with simultaneous load and cross-head disNace-
2. Experimental details and results 2.1. Material
The alumina, obtained from Vesuvius Zyalons and manufactured by hot isostatic pressing at 1600 °C for 1 h, had a nominal purity of 99.98%. The principal impurities, silicon, sodium and potassium, are all glassforming elements and a thin amorphous grain boundary layer was expected (manufacturer's analysis: 50 ppm Si, 40 ppm K, 12 ppm Na, 6 ppm Fe, 1 ppm
Fig. 1. Polycrystalline alumina microstructure (thermally etched).
T. J. Marrow, S. G. Roberts
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High temperamre fracttlre qf alumina
ment data logging using either load (0.01N m i n - ~ - 2 0 N min ~) or displacement (0.001 mm m i n - ~ - I mm rain t) feedback control. Tests were conducted at room temperature (approximately 20 °C), 900 °C and 1200 °C. Specimens annealed at 900 °C and 1200 °C after precracking were also tested at room temperature. Fracture surfaces were gold sputtered and examined using scanning electron microscopy (SEM) to determine the crack length at fracture and the extent of preceding slow crack growth. Tests which did not exhibit a satisfactory crack were discarded.
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ture at 900 °C showed a small amount of transgranular cleavage, but was mostly intergranular. Fast fracture at 1200 °C was fully intergranular. The crack front at the slow crack growth-fast fracture transition was indi-
2.3. Fracture toughness results The room temperature fracture toughness was approximately constant (average fracture toughness 5 MPa m ~!2)for crack extensions of up to 0.5 mm, with no measurable R curve behaviour (Fig. 2). The average precrack length was 0.35 mm. SEM examination showed a region of mixed intergranular and transgranular fracture (Fig. 3 and Fig. 4(a)) between the transgranular precrack and the transgranular fast fracture region (Fig. 4(b)). Decreasing loading rate increased the extent of the intergranular-transgranular region by slow crack growth during the test. Crack growth velocities were not determined. Annealing at 900 °C and 1200 °C had no effect on the fracture toughness. Increasing slow crack growth with decreasing loading rate was also observed at 900 °C and 1200 °C. The slow crack growth paths at 900 °C and 1200 °C were completely intergranular (Fig. 5). The fast frac-
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Fig. 3. A precracked fracture toughness specimen broken at room temperature, illustrating precrack (region a), slow crack growth re, on (region b) and fast fracture (region c).
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Fig. 4. Room temperature fracture surfaces: (a) slow crack growth by intergranular fracture and transgranular cleavage: (b) fast fracture by transgranular cleavage.
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Fig. 5. High temperature fracture surface of slow crack growth by fully intergranular fracture ( 1200 °C).
cated by a change in the macroscopic fracture plane. R curve behaviour was observed at both 900 °C and 1200 °C (Fig. 2). Effective fracture toughness values of up to 9.5 MPa m 1/2 were observed, whereas the intrinsic fracture toughness value at 1200 °C was less than 3 MPa m ~/2. 2.4. Scanning acoustic microscopy Precracked samples were stressed at 900 °C and 1200 °C, at an applied stress intensity factor of approximately 3 MPa m l/2, until relaxation was observed by a change in the cross-head displacement under constant load. After grinding and careful polishing the crack profiles were examined using a Leitz ELSAM high resolution reflection scanning acoustic microscope before thermal etching for SEM examination [35]. An intergranular stable crack growth region was observed at both temperatures which showed discrete zones along the crack path of good acoustic transmission of surface Rayleigh waves. These were indicated by the absence of fringes parallel to the crack image in the scanning acoustic microscope (Fig. 6). Fringes require good reflection and poor transmission of Rayleigh waves [33, 35]. The corresponding features were identified by SEM (Fig. 7). Transmitting regions
Fig. 6. Scanning acoustic microscope image of the crack profile produced by slow crack growth at 1200 °C. Regions without fringes indicate high Rayleigh wave transmission across the crack. Positions 1 and 2 correspond to Figs. 7(a) and 7(b).
were not observed along the room temperature transgranular precrack, which was strongly reflecting. There was no evidence of crack blunting and microcracking was not observed in either specimen.
3. Discussion The average room temperature fracture toughness value (5 MPa) is slightly higher than average values reported in the literature for fine aluminas [30, 36] (4 MPa ml/2-4.5 MPa m~/2). Negligible R curve behaviour has been observed previously in fine-grained aluminas [37, 38]. The grain size is below the theoretical size limit for microcrack toughening [9, 13, 14], which is confirmed by the absence of observable microcracking. The high fraction of transgranular
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Fig. 7. Scanning electron microscope images of asperity contact at the areas of high Rayleigh wave transmission observed in Fig. 6. The crack path was produced by slow crack growth at 1200 °C. (a) and (b) correspond to positions 1 and 2 respectively in Fig. 6.
cleavage during the environmentally assisted slow crack growth may discourage the formation of crack bridges. A bridging zone might also saturate at a small size for a long crack in a fine-grained alumina beam specimen. For a small fraction of bridging grains, a systematic examination using repeated sectioning of the crack profile would be required to confirm the presence of room temperature crack bridging in this particular combination of material and test specimen geometry. Crack deflection by elongated grains may also contribute to the higher fracture toughness value 139]. The high temperature tests, performed over a range of strain rates, fall on a single curve of effective fracture toughness as a function of crack growth. The extent of crack growth increased with decreasing loading rate. This is interpreted as R curve behaviour via slow crack growth, which is strain rate controlled, and a fast fracture mechanism which is strain rate insensitive over these test conditions, lntergranular slow crack growth at both 900 °C and 1200 °C is considered to be driven by grain boundary sliding [24], assisted by the presence of an amorphous grain boundary glass. Crack blunting, which was not observed in these tests, may occur at very low loading rates in some specimen geometries [40]. A study is currently in progress to confirm the presence of grain boundary glass using high resolution
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transmission electron microscopy. Grain boundary glass has been observed in a similar alumina [34]. Fractographic examination using SEM showed no evidence to support adhesion of intergranular fracture surfaces (Fig. 4(a)), as expected from the low content of glass forming elements. Crack face adhesion by viscous glass is therefore not considered to be responsible for the R curve behaviour. The zones of high acoustic transmission, observed using SAM (Fig. 6), in the high temperature slow crack growth region could be correlated directly with SEM observations which suggest asperity interlocking and frictional contact (Fig. 7). Physical contact across the crack would increase transmission and decrease reflection of surface Rayleigh waves [35]. The high temperature slow growth crack path was always intergranular and crack bridging by intact grains was not observed. It is proposed that the convoluted crack path of intergranular slow crack growth causes crack bridging via asperity contact and interlocking, rather than by bridge creation from the bypassing of unfawmrably oriented grains by the main crack. Microcracking was not observed by SAM or SEM at any temperature. Although closed microcracks would be difficult to image using SAM, owing to their high transmission and poor reflection [35], they would necessarily have a negligible shielding effect [7]. The similarity of the R curves at 900 °C and 1200 °C (Fig. 2) may arise from similarities in fracture mode, crack length and specimen geometry. However, the transgranular cleavage during fast fracture at 90t) °C suggests that the intrinsic fracture toughness at 900 °C may be higher than at 121)0 °C and this similarity of R curves may be just coincidental. Further tests would be required to determine accurately the absolute effect of crack bridging by contact in the crack wake on the effective fracture toughness.
4. Conclusion
Fine-grained high purity polycrystalline alumina does not possess a strong R curve behaviour at room temperature for long cracks propagating by environmentally assisted slow crack growth. Room temperature slow crack growth occurs by both intergranular fracture and transgranular cleavage. Microcrack shielding does not occur and the effect of any crack bridging is small. Strong R curve behaviour is induced at 900 °C and 1200 °C when crack extension occurs via fully intergranular slow crack growth. Intergranular crack propagation may be driven by non-accommodated grain boundary sliding, assisted by the deformation of a thin amorphous grain boundary layer. The intrinsic fracture
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toughness decreases at high temperature, accompanied by increased intergranular fast fracture. High resolution reflection S A M demonstrated that crack bridging occurred at high temperatures via the interlocking and frictional contact of asperities in the crack profile. Scanning electron microscope examination alone did not allow their positive identification as crack bridges. T h e R curve arises from traction behind the crack tip which increases with increasing crack length. T h e r e was no evidence to support the alternative high temperature toughening mechanisms of crack face adhesion or microcrack shielding. This combined analysis using S A M and SEM may be of considerable use in the study of R curve mechanisms in other materials, such as platelet-reinforced ceramics and composites, where a detailed understanding of crack bridge distribution is required. Optimizing the crack bridging mechanism may produce microstructures possessing high levels of flaw tolerance. Strong R curve behaviour is required if the decrease in fracture strength with increasing flaw size is to be offset by the increased effective fracture toughness.
Acknowledgments T h e authors would like to thank Professor R. J. Brook for provision of laboratory facilities and Dr. G. A. D. Briggs for the use of the scanning acoustic microscope and for helpful discussions. Thanks are also due to V. Luprano and Dr. C. Lawrence. T h e research was funded by SERC Grant G R G 2853.
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