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Crack bridging by SiC fibers during slow crack growth and the resultant fracture toughness of SiCSiCf composites
Scripta Mctallurgica erhhterhl~a,
Pergamon
Vol. 33, No. 12, pp. 2067-2072,199s Elsevier Science Ltd 1995 Acta Metallurgica Inc. F’rinted in the USA All ri&ts reserved 0956-716X/95 $9.50 + .OO
0956-716x(95)00461-0
CRACK BRIDGING BY Sic FIBERS DURING SLOW CRACK GROWTH AND THE RESULTANT FRACTURE TOUGHNESS OF SiC/SiCf COMPOSITES Christopher R. Jones, Charles H , Henager, Jr.* and Russell H. Jones* Associated Western Universities, Richland, WA 99352 * Pacific Northwest Laboratory, Richland, WA 99352 (Received June 9,1995) (Revised August 3, 1995) Iutroduction
Ceramic matrix composites (CMCs) offer the possibility of high-strength, corrosion-resistant, hightemperature materials with a Im&ire m&stanceadequate for use as structural materials in a variety of systems. Reinl?orcementof a brittle ceramic matrix material by brittle fibers or whiskers separated from the matrix by a weak interface contributes to fracture resistance through reinforcement pull-out, crack bridging, crack deflection, and matrix micro-cracking [ 11.Materials reinforced with continuous fibers can exhibit fracture toughness values in the range 15-30 MPa miR.This allows their consideration as advanced structural materials where use of monolithic ceramics would not be appropriate. If, however, CMCs are to be used in systems where long-term stability is required, they must be resistant to time-dependent crack growth (SCG, subcritical crack growth or slow crack growth). The mechanical response during I&&ire in continuous-fiber-reinforced CMCs is dominated by crack wake bridging by long fibers [l]. Fibers bridging the crack are stressed and at high temperatures can exhibit creep, time-dependent interface fracture, time-dependent interface creep, and possibly v&o-elastic effects at the interface. Previous work by Henager and Jones [2,3] showed that during SCG, the transition from stage II cracking to stage III occurred at a higher stress intensity (K) than expected from four-point-bend fracture tests. This pointed to the possible increase in the toughness of a CMC at high temperatures under slow crack growth (SCG) conditions. Fett et al. found an increase in the toughness of a monolithic ceramic after SCG due to “crack wall interactions behind the crack tip relative to a ceramic without these interactions and an increase in the effective crack fmnt length due to an increase in the surface roughness” [4]. The effect of time-dependent fiber bridging was shown in previous work, which suggested that high temperature plasticity of fibers in the crack-wake bridging zone controls the velocity of subcritical cracks in CMCs [2,3]. In CMCs, the zone exhibiting this bridging is increased during SCG because the fibers relax by high-temperature creep. This paper explores the effect of the bridging on resultant toughness after SCG by comparing the toughness of samples, shown by KQ (peak load fracture toughness) and maximum load, with cracks of varying lengths produced by machined notches (no bridging) and SCG (bridged). Exnerimental Procedures
Composites of Nicalon fiber cloth (O/90“) and CVI P-SIC with both C and BN interfaces were used for these tests. Before the S-Sic CVI fabrication step, 1.O um C and 0.4 pm BN (nominal thicknesses) coatings were 2067
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deposited onto the 15 pm diameter Nicalon fibers. The composites, fabricated by RCI of Whittier, CA, are eight-ply huninates, 4 mm in thickness and supplied as 4 mm x 13 mm x 50 mm bars. The larger bars were cut into single-edge-notched bend bar (SENB) specimens of 4 mm x 5.5 mm x 50 mm. The SENB specimens were loaded in four-point bending with a fully articulated SIC bend fixture having a lower span of L = 40 mm and an upper span of I_&, or 20 mm The tests were performed inside a sealed mullite tube in a split clamshell furnace. The load was applied by a SIC pushrod attached to the top sealing flange of the sealed tube using a flexible metal bellows. Specimen midpoint displacement was measured by an alumina tube, holding a type-K thermocouple, spring loaded against the bottom of the bend specimen. This deflected a strain gage extensometer positioned beneath the SIC support tube. The tests were run at 1100°C in gettered Ar. Four-point fracture tests were run using machine-notched and SCG specimens. In this case, the samples were loaded at a crosshead speed of 8 x 10M4 mm/s until failure. This was done to determine P,, (maximum load) and Kc. In one series the sample was machine notched with a diamond blade saw to fractional crack lengths (alpha = aiW), where a is the crack length and W is the specimen width, ranging from 0.17 to 0.7 and then tested. The notch had a radius of approximately 0.33 mm. In the second series, a crack, or damage zone, was slowly grown (SCG) from a starting machined-notch equal to alpha of 0.17 to varying lengths, ranging from alpha equal to 0.3 to 0.6. The SCG process was achieved by loading the sample to a K of 7.5 MPa m”’ and then holding the load within a range of&l 0 N until the desired crack length was reached. The crack length was determined from the instantaneous compliance curve (time-dependent deflection-applied load curve) as discussed in Ref. 3. The instantaneous compliance versus crack length relation was determined experimentally by sectioning several SCG-cracked specimens. Once the correct crack length was reached, the sample was unloaded and then fractured in the same manner used for the machine-notched samples. Experimental Results A comparison of the load-displacement curves of machine-notched and SCG samples at various values of alpha in Fig. 1 shows that the maximum load decreases with increasing notch depth in the machine-notched samples. This is the response expected without bridging. However, in the SCG samples (Fig. 2) the maximum load remains almost constant. The secant compliances of the various SCG specimens are similar presumably because the fibers are the primary load-bearing component and they were unbroken following slow crack growth (see Discussion). When the load-displacement curves at an alpha of 0.4 are compared (Fig. 3) the SCG sample reaches a maximum load about twice that of the machine notched sample. Finally, the Kc values are compared in Fig. 4. The Kc values for these tests were determined from the peak load (P,,J the sample had reached. Kc is then calculated horn
KQ = P_Y,(a)/(BdW) where P,, is the peak load, B is the specimen width, W is the specimen thickness, and Yr(a) is the dimensionless stress intensity factor for the SENB specimen in pure bending [2]. The values of alpha used for computing Kc are the initial values of the notch depth or damage zone length. For the SCG specimens alpha was calculated from the experimentally-derived compliance-crack length curve and represents the original notch depth plus the depth of the multiply-cracked damage zone [3]. The machine-notched samples form a baseline Kc that is independent of alpha, within acceptable scatter. The SCG samples, on the other hand, are not only apparently tougher, starting at a higher Kc, but also exhibit increasing Kc with increasing alpha. The peak stress, shown in Fig. 5 and calculated as op = (3 L P-)/(4 B Wz), shows the opposite behavior with increasing alpha to that of Kc, shown in Fig. 4. The results of this approach to analyzing the data reveal that the &a&me strength of the samples with the machined notch decreases because the cross-sectional area decreases, while for the SCG samples the peak stress is essentially constant, independent of the depth of the slowly grown crack.
Vol. 33, No. 12
TOUGHNESS
OF SiCYSiC, COMPOSITES
2069
Displacenlcnt (nun) Figure 1. Comparison of loaddisplacement curves for varying notch depths,values of alpha, of the machine-notchedspecimens.
Discussion
The magnitude of & increased with increasing alpha for the SCG samples but remained constant for the machine-notched samples. The peak fracture stress remained constant with increasing alpha for the SCG samples but decreased for the machine-notched samples. We speculate that the primary cause of this behavior is crack-wake bridging by the Sic fibers in the SCG samples which shields or reduces the crack-tip stress intensity and serves as the primary load-bearing component during fracture. The slower the crack is grown, the more time the fibers have to creep and bridge the crack rather than fracture. This effect is demonstrated in Fig. 5, which suggests that the load-carrying capacity of the sample is unatfected if the fibers can creep and not fracture during the SCG process. This explanation requires the fibers to be the primary load-bearing component in these composites. Cross-sectional examinations of SCG specimens have revealed that the damage zone region consists of multiple matrix cracks emanating form the notch root but that the bridging fibers are intact throughout the damage zone [3]. This work suggests that the definition of a crack in these composites is based on fiber damage and not matrix damage. When the samples are tm&umd at a constant crosshead speed that exceeds the creep rate of the fibers, the bridging zone is shorter than it is during SCG. This occurs because time is insufticient for stress relaxation resulting in more fi-actured fibers and a shorter bridging length. The increased G of the SCG samples is due
Load
Displncenlent (nun) Figure 2. Comparison of load-displacementcurves for varying eraok lengths,where the cracks were grown by SCG.
2070
Vol. 33, No. 12
TOUGHNESS OF SIC/SE, COMPOSITES
SCC 1000-_:.
800--
./ ._.’
alpha = 0.43 A.-’“. .. *.._ ..
,:’
... ... .
.:’
..
0.1
0.3
0.2
Displecemnt
0.4
0.5
(wm)
Figure 3. Comparison of load-displacement curves for machine-notched and SCG specimens for essentially equal crack lengths.
to bridging that formed during the period of slow crack growth. However, the machine notched specimens start off with no bridging and form only a short bridging zone during the fracture test; they therefore manifest a lower IQ An estimate of the bridging length from a bridged crack model [3] is about 100 pm for the case where there is no fiber creep. A bridging length this short is consistent with the vahd test results f?om the alpha = 0.7 machine-notched specimen. Obviously, as alpha is increased towards unity there will be an alpha for which the bridging zone will be longer than the remaining ligament and the notched test will be invalid. Apparently, this has not happened yet for alpha = 0.7, which has a remaining ligament of 1.65 mm (1650 pm). The increase in I$, or constant peak stress, for samples with cracks grown by SCG results thorn timedependent bridging by the Sic fibers. Henager and Jones [2] modeled time-dependent bridging effects by assuming the fibers contribute time-dependent pinching forces along the crack wall. Experimental creep data for Nicalon fibers were used to derive the creep relaxation of the pinching forces. In this analysis, the bridging length increased from 100 pm at no relaxation to 700 pm for 1000 s fiber relaxation. This analysis clearly demonstrated an increase in bridging length attributable to fiber creep. The similarity of the reloading secant compliances of the SCG specimens (Fig. 2) raises an interesting point. Although there was a significant differences in the total deflection during SCG and, hence, the timedependent instantaneous compliances of these three specimens, the reloading compliances are nearly identical. In fact, the compliances are the same as the alpha = 0.18 machined-notch specimen. During SCG testing, the
00 0.1
Figure4. CompubivaluesofKQasaibcticnofalpha(a~) cracks.
0.1
0.3
0.4
0.5
Alpha
(a/W)
0.6
0.7
0.8
comparing the toughness of samples machine-notched and those with SCG
Vol. 33, No. 12
TOUGHNESS OF Sic/Sic,
2071
COMPOSITES
Peak SlRSS (MM
0.1
0.2
0.3
0.4
0.5
0.6
0.7
0.8
Normalized Crack Length (a/W)
Figure 5. tlmp&on
ofXiiiC
compcde peak stress a~ a function of a/w for machine notched (broken fibers) and slow crack growth
(crack bridging fibers intact).
fibers for alpha = 0.47 have crept more, the damage zone (crack) is longer, and the crack mouth is more open than for the others. However, the compliance on reloading reflects the original machine-notch depth of alpha = 0.17 plus the fully-bridged damage zone. Therefore, the fully-bridged damage zone contributes very little to the time-independent reloading compliance whereas it dominates the time-dependent instantaneous compliance during SCG testing. The peak stress data in Fig. 5 show a slight increase in peak stress with increasing alpha. There is sufficient scatter in the data to make it difficult to state whether or not this is an actual increase. These composite materials are very inhomogeneous and for a given specimen there may be difference in local densities or fiber volume fractions that can account for this scatter (see also Fig. 2). This effect may be real since the time-independent bridging zone length, which controls the fracture strength and is suggested to be - 100 pm during fast fracture, is much shorter than the fully-bridged damage zone developed during SCG. Presumably more fibers are fractured during the final fast fracture for the SCG specimens with the largest damage zone which could result in an increased fracture stress. A second possible cause for the observed behavior may be crack bifurcation (multiple cracking). The branching of cracks would cause a greater compliance for a given crack length, because no single crack-tip would experience the entire applied strain. The result of this greater compliance would be increased toughness and higher KQ. The compliance could continue to increase with crack length as cracks continued to branch, producing an exponentiahy increasing number of cracks. It is doubtful, however, that crack bifurcation could produce an increase in KQ of the magnitude observed. It most likely contributes to the increase but is probably not the entire cause. Holmes and Chermant [S] summarized creep damage in continuous fiber ceramic matrix composites according to the ratio between the fiber and matrix creep rates (CMR). For composites with a CMR < 1, creep damage consists of fiber/matrix debonding and time-dependent fiber tiacture while for composites with a CMR > 1 creep damage consists of matrix fracture and fibers bridging the cracks. Matrix microcracking and fiber bridging can also occur for composites with a CMR ( 1 for creep at stresses above the matrix microcracking stress. Holmes and Chennant consider Nicalon fiber reinforced-WI SIC composites as having a CMR < 1, and with this consideration the present results fit the latter case, where CMR < 1 but the stress at the tip of the crack exceeds the matrix microcracking stress. It is also likely that at 1 lOO”C, Nicalon reinforced-WI Sic has a CMR > 1 because Nicalon has a low creep strength at 1100°C. Holmes and Chermant argue that a composite optimized for creep resistance should have a CMR < 1 because fiber &&tre is more desirable than matrix t?acture, assuming the composite is loaded at a stress less
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than the matrix microcracking stress. However, stress concentrations, random overload stresses, thermal cycles, and cyclic stresses can all lead to matrix microcracking for materials with a CMR < 1 and a uniform stress less than the matrix microcracking stress. Also, matrix fracture at elevated temperatures can expose the fibers and fiber/matrix interfaces to the chemical environment. The current results demonstrate that, given matrix microeracking, the composite mechanical properties are not degraded if fiber stress relaxation results in an adequate bridging length. Time-dependent environmental effects can degrade the bridging effectiveness by reaction with the fiber/matrix interface as demonstrated by Jones and Henager [6,7]; however, approaches are being developed to protect the interface and to develop chemically stable interfaces in the event of matrix microcracking at elevated temperatures.
Summaw Ceramic matrix composites show the uncommon and promising property of increasing their toughness and maintaining their load-canying capacity with increasing matrix crack lengths under conditions of subcritical crack growth (SCG). As the crack grows, formation of a more extensive bridging zone allows it to withstand far greater stress intensities than would be expected horn fracture tests. Crack bifurcation possibly contributes to this effect but is not likely to be the primary cause. The results of these tests are encouraging and expand the possible uses of these types of materials in applications that require long-term stability under stress. Acknowledements Special thanks are given to J. L. Humason for his testing expertise. This work was supported by Basic Energy Sciences under U.S. Department ofEnergy (DOE) Contract DE-AC06-76RL0 1830 with Battelle Memorial Institute, which operates Pacific Northwest Laboratory for DOE.
1. 2. 3. 4. 5.
6. 7.
A. G. Evans, J. Am Ceram. Sot., 73 (1990) 187. C. H. Henager, Jr., and R. H. Jones, Mater. Sci. and Eng., Al66 (1993) pg. 211. C. H. Henager, Jr., and R H. Jones, J. Am Cemm. Sot. 77 (1994) pg. 2381. T. Fett, J. Am_ Ceram. Sot., 75 (1992) pg 3133. J. W. Holmes and J. L. Chemunt in “Proceedings of HTGMCI, High Temperature Ceramic Matrix Composites”, The 6th European Confknce on Composite Materials, September 20-24, 1993, Hordeaux, France, R. Naslain, J. Lamon end D. Doumeingts, Eds., Woodhead Publiig, J&l., Abmgton, Cambridge, C816AH, England, 1993, pp. 633-647. C. H. Heauger, Jr. and R H. Jones, in “Ceramic Transactions, Vol. 38: Advances in Ceramic-Matrix Composites”, N. P. Hamal, et al., Ekls., The American Ceramics Society, Westewille, OH, 1994, pp. 3 17-328. R H. Jones, C. H. Henager, Jr., end C.F. Wmdisch, Jr., Mater. Sci. and Eng., in press.
Pergamon
Vol. 33, No. 12, pp. 2067-2072,199s Elsevier Science Ltd 1995 Acta Metallurgica Inc. F’rinted in the USA All ri&ts reserved 0956-716X/95 $9.50 + .OO
0956-716x(95)00461-0
CRACK BRIDGING BY Sic FIBERS DURING SLOW CRACK GROWTH AND THE RESULTANT FRACTURE TOUGHNESS OF SiC/SiCf COMPOSITES Christopher R. Jones, Charles H , Henager, Jr.* and Russell H. Jones* Associated Western Universities, Richland, WA 99352 * Pacific Northwest Laboratory, Richland, WA 99352 (Received June 9,1995) (Revised August 3, 1995) Iutroduction
Ceramic matrix composites (CMCs) offer the possibility of high-strength, corrosion-resistant, hightemperature materials with a Im&ire m&stanceadequate for use as structural materials in a variety of systems. Reinl?orcementof a brittle ceramic matrix material by brittle fibers or whiskers separated from the matrix by a weak interface contributes to fracture resistance through reinforcement pull-out, crack bridging, crack deflection, and matrix micro-cracking [ 11.Materials reinforced with continuous fibers can exhibit fracture toughness values in the range 15-30 MPa miR.This allows their consideration as advanced structural materials where use of monolithic ceramics would not be appropriate. If, however, CMCs are to be used in systems where long-term stability is required, they must be resistant to time-dependent crack growth (SCG, subcritical crack growth or slow crack growth). The mechanical response during I&&ire in continuous-fiber-reinforced CMCs is dominated by crack wake bridging by long fibers [l]. Fibers bridging the crack are stressed and at high temperatures can exhibit creep, time-dependent interface fracture, time-dependent interface creep, and possibly v&o-elastic effects at the interface. Previous work by Henager and Jones [2,3] showed that during SCG, the transition from stage II cracking to stage III occurred at a higher stress intensity (K) than expected from four-point-bend fracture tests. This pointed to the possible increase in the toughness of a CMC at high temperatures under slow crack growth (SCG) conditions. Fett et al. found an increase in the toughness of a monolithic ceramic after SCG due to “crack wall interactions behind the crack tip relative to a ceramic without these interactions and an increase in the effective crack fmnt length due to an increase in the surface roughness” [4]. The effect of time-dependent fiber bridging was shown in previous work, which suggested that high temperature plasticity of fibers in the crack-wake bridging zone controls the velocity of subcritical cracks in CMCs [2,3]. In CMCs, the zone exhibiting this bridging is increased during SCG because the fibers relax by high-temperature creep. This paper explores the effect of the bridging on resultant toughness after SCG by comparing the toughness of samples, shown by KQ (peak load fracture toughness) and maximum load, with cracks of varying lengths produced by machined notches (no bridging) and SCG (bridged). Exnerimental Procedures
Composites of Nicalon fiber cloth (O/90“) and CVI P-SIC with both C and BN interfaces were used for these tests. Before the S-Sic CVI fabrication step, 1.O um C and 0.4 pm BN (nominal thicknesses) coatings were 2067
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TOUGHNESS OF SIC/SE, COMPOSITES
Vol. 33, No. 12
deposited onto the 15 pm diameter Nicalon fibers. The composites, fabricated by RCI of Whittier, CA, are eight-ply huninates, 4 mm in thickness and supplied as 4 mm x 13 mm x 50 mm bars. The larger bars were cut into single-edge-notched bend bar (SENB) specimens of 4 mm x 5.5 mm x 50 mm. The SENB specimens were loaded in four-point bending with a fully articulated SIC bend fixture having a lower span of L = 40 mm and an upper span of I_&, or 20 mm The tests were performed inside a sealed mullite tube in a split clamshell furnace. The load was applied by a SIC pushrod attached to the top sealing flange of the sealed tube using a flexible metal bellows. Specimen midpoint displacement was measured by an alumina tube, holding a type-K thermocouple, spring loaded against the bottom of the bend specimen. This deflected a strain gage extensometer positioned beneath the SIC support tube. The tests were run at 1100°C in gettered Ar. Four-point fracture tests were run using machine-notched and SCG specimens. In this case, the samples were loaded at a crosshead speed of 8 x 10M4 mm/s until failure. This was done to determine P,, (maximum load) and Kc. In one series the sample was machine notched with a diamond blade saw to fractional crack lengths (alpha = aiW), where a is the crack length and W is the specimen width, ranging from 0.17 to 0.7 and then tested. The notch had a radius of approximately 0.33 mm. In the second series, a crack, or damage zone, was slowly grown (SCG) from a starting machined-notch equal to alpha of 0.17 to varying lengths, ranging from alpha equal to 0.3 to 0.6. The SCG process was achieved by loading the sample to a K of 7.5 MPa m”’ and then holding the load within a range of&l 0 N until the desired crack length was reached. The crack length was determined from the instantaneous compliance curve (time-dependent deflection-applied load curve) as discussed in Ref. 3. The instantaneous compliance versus crack length relation was determined experimentally by sectioning several SCG-cracked specimens. Once the correct crack length was reached, the sample was unloaded and then fractured in the same manner used for the machine-notched samples. Experimental Results A comparison of the load-displacement curves of machine-notched and SCG samples at various values of alpha in Fig. 1 shows that the maximum load decreases with increasing notch depth in the machine-notched samples. This is the response expected without bridging. However, in the SCG samples (Fig. 2) the maximum load remains almost constant. The secant compliances of the various SCG specimens are similar presumably because the fibers are the primary load-bearing component and they were unbroken following slow crack growth (see Discussion). When the load-displacement curves at an alpha of 0.4 are compared (Fig. 3) the SCG sample reaches a maximum load about twice that of the machine notched sample. Finally, the Kc values are compared in Fig. 4. The Kc values for these tests were determined from the peak load (P,,J the sample had reached. Kc is then calculated horn
KQ = P_Y,(a)/(BdW) where P,, is the peak load, B is the specimen width, W is the specimen thickness, and Yr(a) is the dimensionless stress intensity factor for the SENB specimen in pure bending [2]. The values of alpha used for computing Kc are the initial values of the notch depth or damage zone length. For the SCG specimens alpha was calculated from the experimentally-derived compliance-crack length curve and represents the original notch depth plus the depth of the multiply-cracked damage zone [3]. The machine-notched samples form a baseline Kc that is independent of alpha, within acceptable scatter. The SCG samples, on the other hand, are not only apparently tougher, starting at a higher Kc, but also exhibit increasing Kc with increasing alpha. The peak stress, shown in Fig. 5 and calculated as op = (3 L P-)/(4 B Wz), shows the opposite behavior with increasing alpha to that of Kc, shown in Fig. 4. The results of this approach to analyzing the data reveal that the &a&me strength of the samples with the machined notch decreases because the cross-sectional area decreases, while for the SCG samples the peak stress is essentially constant, independent of the depth of the slowly grown crack.
Vol. 33, No. 12
TOUGHNESS
OF SiCYSiC, COMPOSITES
2069
Displacenlcnt (nun) Figure 1. Comparison of loaddisplacement curves for varying notch depths,values of alpha, of the machine-notchedspecimens.
Discussion
The magnitude of & increased with increasing alpha for the SCG samples but remained constant for the machine-notched samples. The peak fracture stress remained constant with increasing alpha for the SCG samples but decreased for the machine-notched samples. We speculate that the primary cause of this behavior is crack-wake bridging by the Sic fibers in the SCG samples which shields or reduces the crack-tip stress intensity and serves as the primary load-bearing component during fracture. The slower the crack is grown, the more time the fibers have to creep and bridge the crack rather than fracture. This effect is demonstrated in Fig. 5, which suggests that the load-carrying capacity of the sample is unatfected if the fibers can creep and not fracture during the SCG process. This explanation requires the fibers to be the primary load-bearing component in these composites. Cross-sectional examinations of SCG specimens have revealed that the damage zone region consists of multiple matrix cracks emanating form the notch root but that the bridging fibers are intact throughout the damage zone [3]. This work suggests that the definition of a crack in these composites is based on fiber damage and not matrix damage. When the samples are tm&umd at a constant crosshead speed that exceeds the creep rate of the fibers, the bridging zone is shorter than it is during SCG. This occurs because time is insufticient for stress relaxation resulting in more fi-actured fibers and a shorter bridging length. The increased G of the SCG samples is due
Load
Displncenlent (nun) Figure 2. Comparison of load-displacementcurves for varying eraok lengths,where the cracks were grown by SCG.
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Vol. 33, No. 12
TOUGHNESS OF SIC/SE, COMPOSITES
SCC 1000-_:.
800--
./ ._.’
alpha = 0.43 A.-’“. .. *.._ ..
,:’
... ... .
.:’
..
0.1
0.3
0.2
Displecemnt
0.4
0.5
(wm)
Figure 3. Comparison of load-displacement curves for machine-notched and SCG specimens for essentially equal crack lengths.
to bridging that formed during the period of slow crack growth. However, the machine notched specimens start off with no bridging and form only a short bridging zone during the fracture test; they therefore manifest a lower IQ An estimate of the bridging length from a bridged crack model [3] is about 100 pm for the case where there is no fiber creep. A bridging length this short is consistent with the vahd test results f?om the alpha = 0.7 machine-notched specimen. Obviously, as alpha is increased towards unity there will be an alpha for which the bridging zone will be longer than the remaining ligament and the notched test will be invalid. Apparently, this has not happened yet for alpha = 0.7, which has a remaining ligament of 1.65 mm (1650 pm). The increase in I$, or constant peak stress, for samples with cracks grown by SCG results thorn timedependent bridging by the Sic fibers. Henager and Jones [2] modeled time-dependent bridging effects by assuming the fibers contribute time-dependent pinching forces along the crack wall. Experimental creep data for Nicalon fibers were used to derive the creep relaxation of the pinching forces. In this analysis, the bridging length increased from 100 pm at no relaxation to 700 pm for 1000 s fiber relaxation. This analysis clearly demonstrated an increase in bridging length attributable to fiber creep. The similarity of the reloading secant compliances of the SCG specimens (Fig. 2) raises an interesting point. Although there was a significant differences in the total deflection during SCG and, hence, the timedependent instantaneous compliances of these three specimens, the reloading compliances are nearly identical. In fact, the compliances are the same as the alpha = 0.18 machined-notch specimen. During SCG testing, the
00 0.1
Figure4. CompubivaluesofKQasaibcticnofalpha(a~) cracks.
0.1
0.3
0.4
0.5
Alpha
(a/W)
0.6
0.7
0.8
comparing the toughness of samples machine-notched and those with SCG
Vol. 33, No. 12
TOUGHNESS OF Sic/Sic,
2071
COMPOSITES
Peak SlRSS (MM
0.1
0.2
0.3
0.4
0.5
0.6
0.7
0.8
Normalized Crack Length (a/W)
Figure 5. tlmp&on
ofXiiiC
compcde peak stress a~ a function of a/w for machine notched (broken fibers) and slow crack growth
(crack bridging fibers intact).
fibers for alpha = 0.47 have crept more, the damage zone (crack) is longer, and the crack mouth is more open than for the others. However, the compliance on reloading reflects the original machine-notch depth of alpha = 0.17 plus the fully-bridged damage zone. Therefore, the fully-bridged damage zone contributes very little to the time-independent reloading compliance whereas it dominates the time-dependent instantaneous compliance during SCG testing. The peak stress data in Fig. 5 show a slight increase in peak stress with increasing alpha. There is sufficient scatter in the data to make it difficult to state whether or not this is an actual increase. These composite materials are very inhomogeneous and for a given specimen there may be difference in local densities or fiber volume fractions that can account for this scatter (see also Fig. 2). This effect may be real since the time-independent bridging zone length, which controls the fracture strength and is suggested to be - 100 pm during fast fracture, is much shorter than the fully-bridged damage zone developed during SCG. Presumably more fibers are fractured during the final fast fracture for the SCG specimens with the largest damage zone which could result in an increased fracture stress. A second possible cause for the observed behavior may be crack bifurcation (multiple cracking). The branching of cracks would cause a greater compliance for a given crack length, because no single crack-tip would experience the entire applied strain. The result of this greater compliance would be increased toughness and higher KQ. The compliance could continue to increase with crack length as cracks continued to branch, producing an exponentiahy increasing number of cracks. It is doubtful, however, that crack bifurcation could produce an increase in KQ of the magnitude observed. It most likely contributes to the increase but is probably not the entire cause. Holmes and Chermant [S] summarized creep damage in continuous fiber ceramic matrix composites according to the ratio between the fiber and matrix creep rates (CMR). For composites with a CMR < 1, creep damage consists of fiber/matrix debonding and time-dependent fiber tiacture while for composites with a CMR > 1 creep damage consists of matrix fracture and fibers bridging the cracks. Matrix microcracking and fiber bridging can also occur for composites with a CMR ( 1 for creep at stresses above the matrix microcracking stress. Holmes and Chennant consider Nicalon fiber reinforced-WI SIC composites as having a CMR < 1, and with this consideration the present results fit the latter case, where CMR < 1 but the stress at the tip of the crack exceeds the matrix microcracking stress. It is also likely that at 1 lOO”C, Nicalon reinforced-WI Sic has a CMR > 1 because Nicalon has a low creep strength at 1100°C. Holmes and Chermant argue that a composite optimized for creep resistance should have a CMR < 1 because fiber &&tre is more desirable than matrix t?acture, assuming the composite is loaded at a stress less
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Vol. 33, No. 12
than the matrix microcracking stress. However, stress concentrations, random overload stresses, thermal cycles, and cyclic stresses can all lead to matrix microcracking for materials with a CMR < 1 and a uniform stress less than the matrix microcracking stress. Also, matrix fracture at elevated temperatures can expose the fibers and fiber/matrix interfaces to the chemical environment. The current results demonstrate that, given matrix microeracking, the composite mechanical properties are not degraded if fiber stress relaxation results in an adequate bridging length. Time-dependent environmental effects can degrade the bridging effectiveness by reaction with the fiber/matrix interface as demonstrated by Jones and Henager [6,7]; however, approaches are being developed to protect the interface and to develop chemically stable interfaces in the event of matrix microcracking at elevated temperatures.
Summaw Ceramic matrix composites show the uncommon and promising property of increasing their toughness and maintaining their load-canying capacity with increasing matrix crack lengths under conditions of subcritical crack growth (SCG). As the crack grows, formation of a more extensive bridging zone allows it to withstand far greater stress intensities than would be expected horn fracture tests. Crack bifurcation possibly contributes to this effect but is not likely to be the primary cause. The results of these tests are encouraging and expand the possible uses of these types of materials in applications that require long-term stability under stress. Acknowledements Special thanks are given to J. L. Humason for his testing expertise. This work was supported by Basic Energy Sciences under U.S. Department ofEnergy (DOE) Contract DE-AC06-76RL0 1830 with Battelle Memorial Institute, which operates Pacific Northwest Laboratory for DOE.
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6. 7.
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