ournal Inm Ceran. Soc, 82[1]117-28(1999) Creep and Fatigue Behavior in Hi-Nicalon TM-Fiber-Reinforced Silicon Carbide Composites at High Temperatures Shijie Zhu, t, t Mineo Mizuno, "t Yutaka Kagawa, Jianwu Cao, .t Yasuo Nagano, ,t and Hiroshi Kayas Japan Fine Ceramics Center, Nagoya 456, Japan; Institute of Industrial Sciences, University of Tok kyo 106 Japan; and Petroleum Energy Center, Tokyo 106, Japan Monotonic tension, creep, and fatigue tests of a composite interfaces, which allows the intact fibers to bridge crack t the bon). The weak interface can cause a crack to deflect along the where a silicon carbide (SiC) matrix that contains glass- However although the use of weak interfaces can increase the forming, boron-based particulates has been reinforced with fracture toughness and thermal shock resistance, 2 it is not com- Hi-NicalonTM fiber(Hi-NicalonMSiC) were conducted atible with creep and fatigue resistance at high temperature, ir at 1300C, and creep tests also were conducted in argon which demands strong interfaces that resist the nucleation and t 1300C. The ultimate tensile strength (UTS)of the growth of cavities. 3,4 Hi-NicalonTM/SiC composite was similar to that of a Sic The carbon-coating layer in SiC/SiC composites leads to low composite where a pure Sic matrix is reinforced with oxidation resistance at high temperatures in air. a glass Nicalon'M fiber(standard Sic/Sic) and a Sic composite forming, boron-based particulate that reacts with oxygen to where a matrix of glass-forming, boron-based particulates roduce a sealant glass that inhibits oxidation can be added to is reinforced with Nicalon M fiber (enhanced sic/sic); the matrix. 6 This technology is applied to SiC/Sic composites however, the strains at UTS of the Hi-Nicalon/SiC com- The modified SiC/SiC is called an enhanced SiC/SiC compos- osite and the enhanced SiC/SiC composite were much ite. 6, II Enhanced SiC/SiC composites exhibi larger than that of the standard SiC/SiC composite. The temperature(up to 1300 C) properties in alr. abit good high- Youngs modulus of the Hi-Nicalon Sic composite was Because matrix microcracking occurs during the initial ap -140 GPa. which is higher than that of the enhanced Sic/ ication of a creep load, fiber bridging of matrix cracks oper SiC composite(90 GPa) and lower than that of the standard tes during the creep of standard SiC/Sic composites at high C/SiC composite(200 GPa) at a temperature of 1300 C. stresses, although the creep resistance of SiC fibers is lower The minimum strain rates of cyclic creep(fatigue)were than that of the lower than those of static creep. The time to rupture under for the environmental resistance of the composites, if reep loads was slightly shorter than that under fatigue xposed to air. Because the creep of fibers control loads at a given maximum stress. The creep strain rates of crack growth, 24,25 increasing the creep resistance of th the Hi-NicalonTMsiC composite in air were lower than te. moreover those in argon. Consequently, the time to rupture at a given decreasing the creep resistance of the matrix by adding oxides stress in air was longer than in argon. The creep and fa tigue resistance of the Hi-Nicalon Sic composite both also is expected to increase the creep resistance and simul- taneously improve the environmental resistance. We found that were similar to that of the enhanced Sic/SiC composite but the addition of glassy phases in the Sic matrix increased the were much better than that of the standard siC/SiC col creep and oxidation resistance in an enhanced SiC/SiC com- posite in air. However, in argon, the standard SiC/Sic cor osite at 1300.C, compared with that of the standard SiC/SiC osite had the creep rate whereas the enhanced con the highest creep rate. The time to Cyclic fatigue behavior of CMCs at high temperatures is not SiC/SiC composite was the shortes well understood Conditions such as environment factors. cree and the Hi-NicalonTMSiC composite had the longest life of constituents, thermally induced stresses at interfaces. and interfacial sliding resistance may cause the reduction of fatigue life at high temperatures dation of the interfacial sliding resistance have been considered INTHE HE recent decade, the creep and fatigue of continuous- to be the reasons for decreased fatigue resistance at high tem- T-reinforced ceramic-matrix composites(CMCs) have eratures in a standard SiC/SiC composite. 8, 10, 20 en investigated, because these properties are very impor- The presence of a silicon oxycarbide(SiC, O, ) amorphous for the application of CMCs. To obtain high fracture phase in Nicalon TM fibers(Nippon Carbon Co., Tokyo, Ja toughness and thermal shock resistance. CMCs have been de is responsible for the low creep resistance, because of a viscous signed with a weak interface between the fibers and the matrix flow at temperatures as low as 1000%-1200C 6 The SiC,O (e.g, the interface in a SiC/SiC composite is coated with car- phase decomposes, forming SiC and gaseous species such as CO and SiO, whose diffusion through the fiber and reaction with the free carbon are believed to create pores and other defects in the fiber structure. 27 Such a decomposition cause R. Naslain--contributing editor degradations in the strength and the Youngs modulus and affects the fiber creep behavior, even in inert atmospheres. 28-30 The elimination of Sic, O from the fibers by electron irradia- tion under vacuum, instead of curing in air, can improve the of 12 Rtcmiotr ceramic g as ugine approed erroe imgs a. creep resistance I-35 The Youngs modulus also is creased 31 The modified Nicalon TM fibers are called merican Ceramic Ceramics Center NicalonTM fibers. To increase the mechanical properties of SiC/ Industrial Sciences, University of Tokyo SiC composites, Hi-Nicalon M fibers have been used to pEtroleum Energy Center reinforce the enhanced SiC matrix. This paper will present 117
Creep and Fatigue Behavior in Hi-Nicalon™-Fiber-Reinforced Silicon Carbide Composites at High Temperatures Shijie Zhu,†,‡ Mineo Mizuno,*,† Yutaka Kagawa,*,‡ Jianwu Cao,*,† Yasuo Nagano,*,† and Hiroshi Kaya§ Japan Fine Ceramics Center, Nagoya 456, Japan; Institute of Industrial Sciences, University of Tokyo, Tokyo 106, Japan; and Petroleum Energy Center, Tokyo 106, Japan Monotonic tension, creep, and fatigue tests of a composite where a silicon carbide (SiC) matrix that contains glassforming, boron-based particulates has been reinforced with Hi-Nicalon™ fiber (Hi-Nicalon™/SiC) were conducted in air at 1300°C, and creep tests also were conducted in argon at 1300°C. The ultimate tensile strength (UTS) of the Hi-Nicalon™/SiC composite was similar to that of a SiC composite where a pure SiC matrix is reinforced with Nicalon™ fiber (standard SiC/SiC) and a SiC composite where a matrix of glass-forming, boron-based particulates is reinforced with Nicalon™ fiber (enhanced SiC/SiC); however, the strains at UTS of the Hi-Nicalon™/SiC composite and the enhanced SiC/SiC composite were much larger than that of the standard SiC/SiC composite. The Young’s modulus of the Hi-Nicalon™/SiC composite was ∼140 GPa, which is higher than that of the enhanced SiC/ SiC composite (90 GPa) and lower than that of the standard SiC/SiC composite (200 GPa) at a temperature of 1300°C. The minimum strain rates of cyclic creep (fatigue) were lower than those of static creep. The time to rupture under creep loads was slightly shorter than that under fatigue loads at a given maximum stress. The creep strain rates of the Hi-Nicalon™/SiC composite in air were lower than those in argon. Consequently, the time to rupture at a given stress in air was longer than in argon. The creep and fatigue resistance of the Hi-Nicalon™/SiC composite both were similar to that of the enhanced SiC/SiC composite but were much better than that of the standard SiC/SiC composite in air. However, in argon, the standard SiC/SiC composite had the lowest creep rate, whereas the enhanced SiC/SiC composite had the highest creep rate. The time to rupture of the standard SiC/SiC composite was the shortest and the Hi-Nicalon™/SiC composite had the longest life. I. Introduction I N THE recent decade, the creep and fatigue of continuousfiber-reinforced ceramic-matrix composites (CMCs) have been investigated,1–25 because these properties are very important for the application of CMCs. To obtain high fracture toughness and thermal shock resistance, CMCs have been designed with a weak interface between the fibers and the matrix (e.g., the interface in a SiC/SiC composite is coated with carbon). The weak interface can cause a crack to deflect along the interfaces, which allows the intact fibers to bridge crack faces.1 However, although the use of weak interfaces can increase the fracture toughness and thermal shock resistance,2 it is not compatible with creep and fatigue resistance at high temperature, which demands strong interfaces that resist the nucleation and growth of cavities.3,4 The carbon-coating layer in SiC/SiC composites leads to low oxidation resistance at high temperatures in air. A glassforming, boron-based particulate that reacts with oxygen to produce a sealant glass that inhibits oxidation can be added to the matrix.6 This technology is applied to SiC/SiC composites. The modified SiC/SiC is called an enhanced SiC/SiC composite.6,11 Enhanced SiC/SiC composites exhibit good hightemperature (up to 1300°C) properties in air.11 Because matrix microcracking occurs during the initial application of a creep load, fiber bridging of matrix cracks operates during the creep of standard SiC/SiC composites at high stresses, although the creep resistance of SiC fibers is lower than that of the SiC matrix.7 This phenomenon is undesirable for the environmental resistance of the composites, if they are exposed to air. Because the creep of fibers controls matrix crack growth,24,25 increasing the creep resistance of the fibers can improve the creep behavior of the composite. Moreover, decreasing the creep resistance of the matrix by adding oxides also is expected to increase the creep resistance25 and simultaneously improve the environmental resistance. We found that the addition of glassy phases in the SiC matrix increased the creep and oxidation resistance in an enhanced SiC/SiC composite at 1300°C, compared with that of the standard SiC/SiC composite.11 Cyclic fatigue behavior of CMCs at high temperatures is not well understood. Conditions such as environment factors, creep of constituents, thermally induced stresses at interfaces, and interfacial sliding resistance may cause the reduction of fatigue life at high temperatures.8,20–22 Creep of the fibers and degradation of the interfacial sliding resistance have been considered to be the reasons for decreased fatigue resistance at high temperatures in a standard SiC/SiC composite.8,10,20 The presence of a silicon oxycarbide (SiCxOy) amorphous phase in Nicalon™ fibers (Nippon Carbon Co., Tokyo, Japan) is responsible for the low creep resistance, because of a viscous flow at temperatures as low as 1000°–1200°C.26 The SiCxOy phase decomposes, forming SiC and gaseous species such as CO and SiO, whose diffusion through the fiber and reaction with the free carbon are believed to create pores and other defects in the fiber structure.27 Such a decomposition causes degradations in the strength and the Young’s modulus and affects the fiber creep behavior, even in inert atmospheres.28–30 The elimination of SiCxOy from the fibers by electron irradiation under vacuum, instead of curing in air, can improve the creep resistance.31–35 The Young’s modulus also is increased.31 The modified Nicalon™ fibers are called HiNicalon™ fibers. To increase the mechanical properties of SiC/ SiC composites, Hi-Nicalon™ fibers have been used to reinforce the enhanced SiC matrix. This paper will present R. Naslain—contributing editor Manuscript No. 191122. Received December 19, 1997; approved April 10, 1998. This work is part of the automotive ceramic gas turbine development programs at the Petroleum Energy Center. *Member, American Ceramic Society. † Japan Fine Ceramics Center. ‡ Institute of Industrial Sciences, University of Tokyo. § Petroleum Energy Center. J. Am. Ceram. Soc., 82 [1] 117–28 (1999) Journal 117
18 Journal of the American Ceramic Sociery-Zhu et al. Vol. 82. No the creep and fatigue behavior of a Hi-NicalonTM fiber-rein- A controlled-atmosphere furnace(Model MTS 659, MTS forced SiC composite System Corp. was used. For the tests in argon, the chamber Creep and fatigue tests of the Hi-Nicalon TM/SiC composite was first allowed to pump down to 5% Similar studies of the effect of bend- reinforced CMCs do not exist. ASTM Practice C 1275-94(a) mation, which can redistribute the stress state and sometimes lead to notch insensitivity, a bending strain of 5% should no affect the strength distribution The fatigue tests were performed with a sinusoidal de quency of 20 Hz in air. The stress ratio (r), which was defined as the ratio of minimum stress to maximum stress, was 0.1 for fatigue tests. Creep tests were conducted under constant load in both air and argon. Creep strain was measured directly from the gauge length of the specimen by using a contact xtensometer(Model 632.53-F71, MTS System Corp ) which has a measuring range of +2.5 mm over its gauge length of 25 mm. Repeated unloading-reloading, with a rate of 50 MPa/s, was ed to measure the modulus change during the creep tests. The specimens were allowed to soak for >30 min at a temperature of 1300C before starting monotonic-tension, 4um creep, or cyclic-fatigue tests Fig. 1.(a) Satin- woven structure of the Hi-Nicalon TM fiber bundles he steel dummy specimen plied by MTS Sy with astm sta
the creep and fatigue behavior of a Hi-Nicalon™ fiber-reinforced SiC composite. Creep and fatigue tests of the Hi-Nicalon™/SiC composite were both performed in air at the same maximum stresses, to compare time-dependent deterioration with cyclic-dependent damage. Creep tests in pure argon also were conducted at the same temperature to understand the effect of environment on the creep. The mechanical behavior of the Hi-Nicalon™/SiC composite was compared with the enhanced SiC/SiC11 and standard SiC/SiC composites10 to investigate the effects of improved fibers and the matrix on the creep and fatigue resistance at high temperature. II. Materials and Experimental Procedures The composites used in this investigation were processed via the chemical vapor infiltration (CVI) of SiC into satin-woven 0°/90° Hi-Nicalon™ fiber preforms (made by Du Pont Lanxide Composites, Wilmington, DE). Before the infiltration, the preforms were coated with carbon via chemical vapor deposition (CVD), to decrease the interface bonding between the fibers and the matrix, thereby increasing the toughness. The composites, which were processed as 200 mm × 200 mm panels with a thickness of 3.2 mm, contained 40 vol% SiC fibers and 9.7% porosity. The diameter of a fiber was ∼12 mm, and each bundle consisted of 500 fibers. The tensile specimens were machined from the panels using diamond cutting tools. The shape and dimensions of the specimens for the monotonic-tension, creep, and cyclic-fatigue tests have been described by Zhu and co-workers.7,8,10,11 The surfaces of the specimens were machined by using an 800 grit grinding wheel before testing. The specimens were not protected by a final CVI run after machining. All the mechanical tests were performed with a servohydraulic testing system (Model MTS 810, MTS System Corp., Eden Prairie, MN) at a temperature of 1300°C. The monotonic tensile tests were conducted in air under a constant stress rate of 50 MPa/s. The alignment between the upper and lower grips of the load unit was verified using a steel dummy specimen for verification.¶ Analytical and empirical analysis studies concluded that, for negligible effects on the estimates of the strength distribution parameters (for example, the Weibull modulus and the characteristic strength) of monolithic advanced ceramics, the allowable bending percentage, as defined in ASTM Practice E 1012, should not be >5%. Similar studies of the effect of bending on the tensile strength distributions of continuous-fiberreinforced CMCs do not exist. ASTM Practice C 1275-94 adopted the recommendations for the tensile testing of monolithic advanced ceramics. Because CMCs have inelastic deformation, which can redistribute the stress state and sometimes lead to notch insensitivity, a bending strain of 5% should not affect the strength distribution. The fatigue tests were performed with a sinusoidal loading frequency of 20 Hz in air. The stress ratio (r), which was defined as the ratio of minimum stress to maximum stress, was 0.1 for fatigue tests. Creep tests were conducted under constant load in both air and argon. Creep strain was measured directly from the gauge length of the specimen by using a contact extensometer (Model 632.53-F71, MTS System Corp.), which has a measuring range of ±2.5 mm over its gauge length of 25 mm. Repeated unloading–reloading, with a rate of 50 MPa/s, was applied to measure the modulus change during the creep tests. The specimens were allowed to soak for >30 min at a temperature of 1300°C before starting monotonic-tension, creep, or cyclic-fatigue tests. A controlled-atmosphere furnace (Model MTS 659, MTS System Corp.) was used. For the tests in argon, the chamber was first allowed to pump down to <13.3 Pa (100 mtorr), and then the chamber was backfilled with high-purity argon gas. The steps were repeated three times to ensure a thorough purge. Enough argon gas was flowed through the chamber to equal five times the chamber volume. The volume percentage of oxygen in the high-purity argon gas was <1 ppm. After fracture, the specimens were examined by using both optical microscopy and scanning electron microscopy (SEM). III. Results (1) Microstructures and Monotonic Tension The satin-woven structure and matrix microstructure of the Hi-Nicalon™/SiC composite in the original state are shown in Fig. 1. The width of the fiber bundles in the satin-woven structure is ∼1.5 mm (Fig. 1(a)). There are glassy phases in the enhanced SiC matrix (dark phases in the matrix shown in Fig. 1(b)). Stress-versus-strain curves of the Hi-Nicalon™/SiC, enhanced SiC/SiC, and standard SiC/SiC composites at a temperature of 1300°C are shown in Fig. 2. The enhanced SiC/SiC composite consists of normal Nicalon™ fibers and the enhanced SiC matrix, which is the same as that in the HiNicalon™/SiC composite. The curve for the Hi-Nicalon™/SiC ¶ The steel dummy specimen was supplied by MTS System Corp. It was designed to allow the bending strain to be <5%, in accordance with ASTM Standard E 1012-89. Fig. 1. (a) Satin-woven structure of the Hi-Nicalon™ fiber bundles in the Hi-Nicalon™/SiC composite; (b) glassy phases (dark blocks) in the matrix of the Hi-Nicalon™/SiC composite. 118 Journal of the American Ceramic Society—Zhu et al. Vol. 82, No. 1
January 1999 Creep and Fatigue Behavior in Hi-Nicalon/SiC Composites at High Temperatures 300 250 200 9150 1300ˇc 100 Standard SiC/SiC, Ar Enhanced SiC/SiC, Air a Hi-Nicalon SiC, Air 0002 0.004 0006 0.008 Tensile strai Fig. 2. Monotonic tensile stress versus strain in the Hi-Nicalon TM/SiC, enhanced SiC/SiC, and standard SiC/SiC composites at 1300C under a constant stress rate of 50 MPa/s composite indicates linear elastic behavior up to the propor- is lower than that in air (n =9.4). The effect of the environ- tional limit of 70 MPa; this stress is -30% of the ultimate ment on the creep resistance in the Hi-Nicalon TM/SiC compos- tensile strength(UTS). The UTS of the Hi-NicalonTM/SiC com- ite is the same as that in the enhanced SiC/SiC composite. 1 posite is similar to that of the standard SiC/SiC and enhanced Nicalon M/SiC and enhanced SiC/SiC composites are much Enhanced SiC/SiC Compositer nd SiC/SiC composites; however, the strains at UTS of the Hi- (3) Comparison with Standard In argon, the creep rate of the Hi-Nicalon TM/SiC composite 105 MPa, and the tertiary creep strains ap. the enhanced SiC/SiC composites is longer than that of the pear at 90 MPa for static loads. However, the curves under standard SiC/SiC composite cyclic loads show primary, secondary, and tertiary stages at In air, the creep rate of the hi-Nicalon M/SiC composite at stresses of 120 and 90 MPa 1300%C is much lower than that of the standard SiC/SiC com- (Fig. 3), the steady-state or minimum strain rates of cyclic creep(fatigue)are always slightly lower than those of stati composite but is also similar to that of the enhanced SiC/SiC composite(Fig. I1) In air, the cyclic-fatigue life versus the maximum stress of The creep strain rate(e)can be described by the power law the Hi-Nicalon TM /SiC composite is almost the same as that the enhanced SiC/SiC composite; however, this lifetime is ∈=Ao"e much longer than that of the standard sic/Sic 1300°C(Fg.12) where A is a constant, o the stress, n the stress exponent for In summary, the creep and fatigue resistance of the Hi- creep, Q the activation energy for creep, R the gas constant, and NicalonTMSiC composite is similar to that of the enhanced T the absolute temperature. The stress exponent for cyclic SiC/SiC composite but is much better than that of the standard creep(n =9.8)is similar to that for static creep(n = 9.4) SiC/SiC composite in However, in argon, the The time to rupture under creep loads is slightly shorter than deformation resistance is not consistent with the creep rupture that under fatigue loads at a given maximum stress other than resistance. The creep rates, from lowest to highest, are as fol- results of the minimum creep rates(Fig. 4) ite, enhanced SiC/SiC composite. However, the time to rupture The creep strain rates of the Hi-Nicalon TM/SiC composite in of the standard SiC/SiC composite is the shortest, and the Hi- ir compared with those in argon in Fig. 6. The creep rates in Nicalon TM/SiC composite has the longest life al argon are evidently higher than those in air. Consequently, the For the standard SiC/SiC composite, the creep rates at a time to rupture at a given stress in air is longer than that in given stress in argon are lower than those in air(Fig. 13) argon(Fig. 7). The stress exponent for creep in argon(n=5 ly, the time to rupture at a given stress in argon is
composite indicates linear elastic behavior up to the proportional limit of 70 MPa; this stress is ∼30% of the ultimate tensile strength (UTS). The UTS of the Hi-Nicalon™/SiC composite is similar to that of the standard SiC/SiC and enhanced SiC/SiC composites; however, the strains at UTS of the HiNicalon™/SiC and enhanced SiC/SiC composites are much higher than that of the standard SiC/SiC composite. The modulus calculated from the linear portion of the curve is ∼140 GPa. This value is higher than that for the enhanced SiC/SiC composite (90 GPa) and lower than that for the standard SiC/SiC composite (200 GPa) at 1300°C. (2) Creep and Fatigue Creep strains versus time at different maximum stresses in air at 1300°C are shown in Fig. 3. Only transient creep strain exists at stresses >105 MPa, and the tertiary creep strains appear at 90 MPa for static loads. However, the curves under cyclic loads show primary, secondary, and tertiary stages at stresses of 120 and 90 MPa. Regardless of whether cyclic creep strain is larger (at high stress) or smaller (at low stresses) than the static creep strain (Fig. 3), the steady-state or minimum strain rates of cyclic creep (fatigue) are always slightly lower than those of static creep (Fig. 4). The creep strain rate (e . ) can be described by the power law e . = Asn expS− Q RTD (1) where A is a constant, s the stress, n the stress exponent for creep, Q the activation energy for creep, R the gas constant, and T the absolute temperature. The stress exponent for cyclic creep (n 4 9.8) is similar to that for static creep (n 4 9.4). The time to rupture under creep loads is slightly shorter than that under fatigue loads at a given maximum stress other than 105 MPa (Fig. 5). This result is qualitatively consistent with the results of the minimum creep rates (Fig. 4). The creep strain rates of the Hi-Nicalon™/SiC composite in air compared with those in argon in Fig. 6. The creep rates in argon are evidently higher than those in air. Consequently, the time to rupture at a given stress in air is longer than that in argon (Fig. 7). The stress exponent for creep in argon (n 4 5) is lower than that in air (n 4 9.4). The effect of the environment on the creep resistance in the Hi-Nicalon™/SiC composite is the same as that in the enhanced SiC/SiC composite.11 (3) Comparison with Standard and Enhanced SiC/SiC Composites In argon, the creep rate of the Hi-Nicalon™/SiC composite at 1300°C (Fig. 8) is lower than that of the enhanced SiC/SiC composite,11 and the time to rupture of the Hi-Nicalon™/SiC composite is longer than that of the enhanced SiC/SiC composite (Fig. 9). However, although the creep rate of both the Hi-Nicalon™/SiC and the enhanced SiC/SiC composites at 1300°C is higher than that of the standard SiC/SiC composite (Fig. 8), the time to rupture for both the Hi-Nicalon™/SiC and the enhanced SiC/SiC composites is longer than that of the standard SiC/SiC composite (Fig. 9). In air, the creep rate of the Hi-Nicalon™/SiC composite at 1300°C is much lower than that of the standard SiC/SiC composite but is similar to that of the enhanced SiC/SiC composite11 (Fig. 10). The time to rupture of the Hi-Nicalon™/SiC composite is much longer than that of the standard SiC/SiC composite but is also similar to that of the enhanced SiC/SiC composite11 (Fig. 11). In air, the cyclic-fatigue life versus the maximum stress of the Hi-Nicalon™/SiC composite is almost the same as that of the enhanced SiC/SiC composite; however, this lifetime is much longer than that of the standard SiC/SiC composite at 1300°C (Fig. 12). In summary, the creep and fatigue resistance of the HiNicalon™/SiC composite is similar to that of the enhanced SiC/SiC composite but is much better than that of the standard SiC/SiC composite in air. However, in argon, the creepdeformation resistance is not consistent with the creep rupture resistance. The creep rates, from lowest to highest, are as follows: standard SiC/SiC composite, Hi-Nicalon™/SiC composite, enhanced SiC/SiC composite. However, the time to rupture of the standard SiC/SiC composite is the shortest, and the HiNicalon™/SiC composite has the longest life. For the standard SiC/SiC composite, the creep rates at a given stress in argon are lower than those in air (Fig. 13). Consequently, the time to rupture at a given stress in argon is Fig. 2. Monotonic tensile stress versus strain in the Hi-Nicalon™/SiC, enhanced SiC/SiC, and standard SiC/SiC composites at 1300°C under a constant stress rate of 50 MPa/s. January 1999 Creep and Fatigue Behavior in Hi-Nicalon™/SiC Composites at High Temperatures 119
120 Journal of the American Ceramic Society-Zhu et al. Vol. 82. Ne 12 o Fatigue 。 Fatigue 0.8 0.4 ●目 0 500 1000 4000 TIme, s Cree ●cre o Fatigue 04 口 Fatigue 三0.5 02 輯H 15 Time,×104s Time,×1059 Fig. 3. Tensile creep strain versus time in the Hi-Nicalon TM/SiC composite under constant load (creep) and cyclic loading(fatigue)at 1300C in ir at the same maximum stresses(a)150,(b)120, (c)90, and(d)75 MPa) than that in air(Fig. 14). The fatigue life at a gi Iven debonding of the interfaces between the fibers and the matrix maximum stress in argon also is longer than that in air(F occurs. The 0 fibers bridge crack faces and, therefore, de- 15). The effects of the environment on the creep resistance in crease the driving force at the crack tip as a general bridging the standard SiC/Sic co te are opposite to those in the mechanism( Fig. 17). At a temperature of 1300C, the glas Hi-NicalonTM/SiC(Figs. 6 and 7)and the enhanced SiC/SiC ases become liquid, which flow into cracks. At room tem- composites. This result will be discussed in Section(2)of the ature, they become solid again and are situated in the cracks (Fig. 17(b). When the testing was performed in air, oxidation occurred at the interfaces and at the fibers and the matrix(Figs (4 Microscopic Damage and fracture 16 and 17). Fibers can be severely damaged by oxidation at Creep and fatigue cracks are always observed at large pores places near the edge of the specimen or close to the large pores among fiber bundles(Fig. 16). When cracks meet 0 fibers, (Figs. 18 and 16). However, such a severe oxidation is not 300 1300°c.Ai 三107 1000 Ma stres Time to Ri Minimum creep strain rate as a function of the maximum in the hi- the Hi-Nicalon TM/S posite under constant load (creep) cyclic load- lic loading( fatigue)at 1300 C in air
longer than that in air (Fig. 14). The fatigue life at a given maximum stress in argon also is longer than that in air (Fig. 15). The effects of the environment on the creep resistance in the standard SiC/SiC composite are opposite to those in the Hi-Nicalon™/SiC (Figs. 6 and 7) and the enhanced SiC/SiC composites.11 This result will be discussed in Section (2) of the Discussion. (4) Microscopic Damage and Fracture Creep and fatigue cracks are always observed at large pores among fiber bundles (Fig. 16). When cracks meet 0° fibers, debonding of the interfaces between the fibers and the matrix occurs. The 0° fibers bridge crack faces and, therefore, decrease the driving force at the crack tip as a general bridging mechanism (Fig. 17). At a temperature of 1300°C, the glassy phases become liquid, which flow into cracks. At room temperature, they become solid again and are situated in the cracks (Fig. 17(b)). When the testing was performed in air, oxidation occurred at the interfaces and at the fibers and the matrix (Figs. 16 and 17). Fibers can be severely damaged by oxidation at places near the edge of the specimen or close to the large pores (Figs. 18 and 16). However, such a severe oxidation is not Fig. 4. Minimum creep strain rate as a function of the maximum stress in the Hi-Nicalon™/SiC composite under constant load (creep) and cyclic loading (fatigue) at 1300°C in air. Fig. 5. Time to rupture versus the maximum stress in the HiNicalon™/SiC composite under constant load (creep) and cyclic loading (fatigue) at 1300°C in air. Fig. 3. Tensile creep strain versus time in the Hi-Nicalon™/SiC composite under constant load (creep) and cyclic loading (fatigue) at 1300°C in air at the same maximum stresses ((a) 150, (b) 120, (c) 90, and (d) 75 MPa). 120 Journal of the American Ceramic Society—Zhu et al. Vol. 82, No. 1
January 1999 Creep and Fatigue Behavior in Hi-Nicalon/SiC Composites at High Temperatures 121 1300°c, Creep Creep,1300°c,Ar n=5 由 Enhanced siC/ si ◇ Standard sic/SiC 10 a Time to Rupture, s Fig. 9. Time to rupture versus stress in Hi-Nicalon TM/SiC, enhanced SiC/SiC, and standard SiC/SiC composites at 1300oC in argon. 300 104 100 106 1300 C, Creep 1300C,Air 109 Enhanced SiC/SiC Hi- 1010 TIme to Rupture,s 1000 Stress, MPa Fig. 7. Time to rupture versus stress in the Hi-Nicalon TM/SiC com- posite at 1300C in air and argo Fig. 10. Minimum creep strain rate as a function of stress in Hi- NicalonTM/SiC, enhanced SiC/SiC, and standard SiC/SiC composites in air at I300°C 300 日费 Hi-Nicalon TM/SiC Enhanced siC/SIC . o.Standard SiC/SiC SiC/SIC Stress. MPa 103 um creep strain rate as a function of stress in the TIme to Rupture,s Ites at n argon. Fig. 11. Time to rupture versus stress in the Hi-Nicalon TM/SiC, er hanced SiC/SiC, and standard SiC/SiC composites in air at 1300C. widely distributed in the specimens that have been tested in air, because the filling of the glassy phases in the cracks prohibi the diffusion of oxygen along the crack paths IV. Discussion The fiber pullout under fatigue is longer than that under creep(Figs. 19 and 20). Much debris can be observed (1) Modulus Change fracture surfaces(Fig. 20) Modulus degradation in cyclic fatigue has been reported for Figure 21 shows the fibers on the fracture surface covered by unidirectional and laminated ceramic composites at room tem- Fig 22(a). However, a large amount of glassy phases formed growth has been shown to i? emperatures.$,18,19, 21, 41Damage a layer of glassy phases. In argon, there is no fiber oxidation perature, 36-0 and elevated on the interfaces(Fig. 22(b)) CMCs under fatigue on g. w mpany a modulus decrease in
widely distributed in the specimens that have been tested in air, because the filling of the glassy phases in the cracks prohibits the diffusion of oxygen along the crack paths. The fiber pullout under fatigue is longer than that under creep (Figs. 19 and 20). Much debris can be observed on the fracture surfaces (Fig. 20). Figure 21 shows the fibers on the fracture surface covered by a layer of glassy phases. In argon, there is no fiber oxidation (Fig. 22(a)). However, a large amount of glassy phases formed on the interfaces (Fig. 22(b)). IV. Discussion (1) Modulus Change Modulus degradation in cyclic fatigue has been reported for unidirectional and laminated ceramic composites at room temperature23,36–40 and elevated temperatures.5,18,19,21,41 Damage growth has been shown to accompany a modulus decrease in CMCs under fatigue loading.38,40 Fig. 11. Time to rupture versus stress in the Hi-Nicalon™/SiC, enhanced SiC/SiC, and standard SiC/SiC composites in air at 1300°C. Fig. 6. Minimum creep strain rate as a function of stress in the Hi-Nicalon™/SiC composite at 1300°C in air and argon. Fig. 7. Time to rupture versus stress in the Hi-Nicalon™/SiC composite at 1300°C in air and argon. Fig. 8. Minimum creep strain rate as a function of stress in the Hi-Nicalon™/SiC, enhanced SiC/SiC, and standard SiC/SiC composites at 1300°C in argon. Fig. 9. Time to rupture versus stress in Hi-Nicalon™/SiC, enhanced SiC/SiC, and standard SiC/SiC composites at 1300°C in argon. Fig. 10. Minimum creep strain rate as a function of stress in HiNicalon™/SiC, enhanced SiC/SiC, and standard SiC/SiC composites in air at 1300°C. January 1999 Creep and Fatigue Behavior in Hi-Nicalon™/SiC Composites at High Temperatures 121
Journal of the American Ceramic Society-Zhu et al. Vol. 82. N g ● Enhancec Hi- 郐- 60 20 10110210310410510107108 101102103104 Cycles to F Cycles to Failure Fig. 12. Maximum stress versus cycles to failure for fatigue in the Fig. 15. Maximum stress versus cycles to failure for fatigue in the Hi-NicalonTMSiC, enhanced SiC/SiC, and standard SiC/SiC compos- standard SiC/SiC composite in air and argon at 1300C ites in air at I300°C crep,1300°c … Argon 10-10 Stress, MPa mum creep strain rate as a function of stress in the standard SiC/SiC composite at 1300.C in air and argon 30um …Argo Fig. 16. Crack initiated at a large pore in the Hi-NicalonTM/SiC specimen crept in air at 1300C and 150 MPa for 480 s. The crack propagates to a fiber ized by the value obtained from the linear portion of the loop during the first loading)versus the number of cycles is shown in Fig. 24(a). At stresses 2120 MPa, the modulus decreases 102103 rapidly within ten cycles, then gradually decreases, and finally Time to Rupture, s decreases rapidly up to fracture. The shape of the curves is similar to the creep strain curves(Fig. 3). At stresses <105 Fig. 14. Time to rupture versus stress in the standard Sic/Sic com- MPa, the modulus initially remained constant up to 10 cycles posite at 1300C in air and argon and then monotonously decreased. At 75 MPa, the modulus remained constant up to 107 cycles, at which point the test was stopped. When the modulus decreased to 20%40% of the To understand the damage evolution and degradation mecha- original value, the specimens fractured nism during fatigue and creep, elastic moduli were measured. The change of the modulus during creep with time in ai Figure 23 shows the evolution of the stress-strain hysteresis ( Fig. 24(b)is similar to that during fatigue(Figs. 24(a). How- loops. The slope decreases and the width of the loops increases ever, the limiting modulus for fracture is 50%-60%, which is as the number of cycles increases. The former indicates th higher than that under fatigue. This result is the same as the decrease of the modulus, and the latter represents the decrease results for the standard SiC/SiC composite, 0 in which it was of the interfacial sliding resistance. The hysteresis loops move explained by the longer debonding of the interfaces under fa to the right along the strain axis(which is known as ratchetting) tigue. When cree?, 30 MPa, at this stress, the specimen did not because of time-dependent deformation. The modulus(normal with time, even at
To understand the damage evolution and degradation mechanism during fatigue and creep, elastic moduli were measured. Figure 23 shows the evolution of the stress–strain hysteresis loops. The slope decreases and the width of the loops increases as the number of cycles increases. The former indicates the decrease of the modulus, and the latter represents the decrease of the interfacial sliding resistance. The hysteresis loops move to the right along the strain axis (which is known as ratchetting) because of time-dependent deformation. The modulus (normalized by the value obtained from the linear portion of the loop during the first loading) versus the number of cycles is shown in Fig. 24(a). At stresses $120 MPa, the modulus decreases rapidly within ten cycles, then gradually decreases, and finally decreases rapidly up to fracture. The shape of the curves is similar to the creep strain curves (Fig. 3). At stresses #105 MPa, the modulus initially remained constant up to 104 cycles and then monotonously decreased. At 75 MPa, the modulus remained constant up to 107 cycles, at which point the test was stopped. When the modulus decreased to 20%–40% of the original value, the specimens fractured. The change of the modulus during creep with time in air (Fig. 24(b)) is similar to that during fatigue (Figs. 24(a)). However, the limiting modulus for fracture is 50%–60%, which is higher than that under fatigue. This result is the same as the results for the standard SiC/SiC composite,10 in which it was explained by the longer debonding of the interfaces under fatigue. When creep tests are in argon, the modulus can decrease with time, even at 30 MPa; at this stress, the specimen did not Fig. 12. Maximum stress versus cycles to failure for fatigue in the Hi-Nicalon™/SiC, enhanced SiC/SiC, and standard SiC/SiC composites in air at 1300°C. Fig. 13. Minimum creep strain rate as a function of stress in the standard SiC/SiC composite at 1300°C in air and argon. Fig. 14. Time to rupture versus stress in the standard SiC/SiC composite at 1300°C in air and argon. Fig. 15. Maximum stress versus cycles to failure for fatigue in the standard SiC/SiC composite in air and argon at 1300°C. Fig. 16. Crack initiated at a large pore in the Hi-Nicalon™/SiC specimen crept in air at 1300°C and 150 MPa for 480 s. The crack propagates to a fiber being oxidized. 122 Journal of the American Ceramic Society—Zhu et al. Vol. 82, No. 1
January 1999 Creep and Fatigue Behavior in Hi-Nicalon MSiC Composites at High Temperatures Fig. 17. Micrographs showing (a) cracks bridged by 0o fibers and(b)interfaces between fibers and the matrix being debonded in the Hi- NicalonTM/SiC specimen fatigued at 1300 C under the maximum stress of 150 MPa in air for 2.5 x 104 cycles fracture for up to lll h. At 60 MPa, fracture occurs at a critical and propagation of the matrix cracks in the specimens under modulus equal to 80% of the original value If the fibers have a lower creep resistance than the matrix, a trix cracks at stresses <105 MPa, according to the constant- adual decrease in the modulus during creep occurs because modulus stage( Fig. 24 ). Creep occurs during this stage, incu- creep of the bridge fibers transfers stress to the matrix and 2120 MPa, extensive matrix cracks are bridged by hbeaG bating damage for propagation of the matrix cracks. At stress causes matrix cracking and crack growth. 42-47 However, it is not known whether Hi-Nicalon TM fibers have a higher or lower Therefore, creep of the fibers promotes propagation of the creep resistance than the SiC matrix with the additives. Be- cracks, which leads to the decrease of the modulus. In argon, cause reduction of the elastic modulus reflects a multiplication the creep resistance of the fibers is lower. At a stress of 30 ohM (b) Fig. 18. (a) Severe oxidation 18(a).e in the Hi-NicalonTMSiC specimen crept in air at 1300oC and 150 MPa for 480 S; (b) high-magnification image of center region in fis
fracture for up to 111 h. At 60 MPa, fracture occurs at a critical modulus equal to 80% of the original value. If the fibers have a lower creep resistance than the matrix, a gradual decrease in the modulus during creep occurs because creep of the bridge fibers transfers stress to the matrix and causes matrix cracking and crack growth.42–47 However, it is not known whether Hi-Nicalon™ fibers have a higher or lower creep resistance than the SiC matrix with the additives. Because reduction of the elastic modulus reflects a multiplication and propagation of the matrix cracks in the specimens under fatigue tests,38 the first loading did not produce extensive matrix cracks at stresses #105 MPa, according to the constantmodulus stage (Fig. 24). Creep occurs during this stage, incubating damage for propagation of the matrix cracks. At stresses $120 MPa, extensive matrix cracks are bridged by fibers. Therefore, creep of the fibers promotes propagation of the cracks, which leads to the decrease of the modulus. In argon, the creep resistance of the fibers is lower. At a stress of 30 Fig. 18. (a) Severe oxidation damage in the Hi-Nicalon™/SiC specimen crept in air at 1300°C and 150 MPa for 480 s; (b) high-magnification image of center region in Fig. 18(a). Fig. 17. Micrographs showing (a) cracks bridged by 0° fibers and (b) interfaces between fibers and the matrix being debonded in the HiNicalon™/SiC specimen fatigued at 1300°C under the maximum stress of 150 MPa in air for 2.5 × 104 cycles. January 1999 Creep and Fatigue Behavior in Hi-Nicalon™/SiC Composites at High Temperatures 123
124 Journal of the American Ceramic Society-Zhu et al. Vol. 82. No ↑↑↑ 2n Fig 19. Lateral surfaces of Hi-Nicalon TM/SiC specimens fractured in air at 1300C under a maximum stress of 120 MPa((a) fatigue and (b) creep). koum 10o (b) Fig. 20. Fracture surfaces of Hi-Nicalon TM/SiC specimens fractured in air at 1300oC under the maximum stress of 120 MPa(a) fatigue and(b)
Fig. 19. Lateral surfaces of Hi-Nicalon™/SiC specimens fractured in air at 1300°C under a maximum stress of 120 MPa ((a) fatigue and (b) creep). Fig. 20. Fracture surfaces of Hi-Nicalon™/SiC specimens fractured in air at 1300°C under the maximum stress of 120 MPa ((a) fatigue and (b) creep). 124 Journal of the American Ceramic Society—Zhu et al. Vol. 82, No. 1
January 1999 Creep and Fatigue Behavior in Hi-Nicalon MSiC Composites at High Temperatures 20 um Fracture surfaces of Hi-Nicalon TM/SiC specimens creep- d in air at 1300oC and 120 MPa, showing that a glassy phaso the surface MPa, matrix cracking can occur via stress redistribution due te creep of the fibers (2) Effects of Fiber, Fiber Architecture, and Matrix behavior of composites is dependent on the cree of the matrix. the fibers the fiber architecture. and the inter- hases and interfaces between the fibers and the matrix. In the Hi-NicalonTM/SiC composite, the creep behavior of the en- hanced SiC matrix is not known. The amount of additives in the matrix also is not known. Therefore. it is ssible to compare the creep resistance of the matrix to that of the fiber to determine the load-transfer direction during creep. However all the testing stresses for creep of the Hi-Nicalon TM/SiC com- osite in air are higher than the proportional limit, which can be ssumed to be an approximate matrix cracking stress. The ma trix microcracking occurs during the initial application of a creep load; therefore, fiber bridging of the matrix cracks al ways operates, whether the creep rate of the fibers is higher or lower than the matrix. In the crack-bridging mechanism, the matrix-crack growth rate is governed by a process in which the (b) increase in the crack length is accompanied by an increase in the number of fibers that bridge the crack. This phenomenon Fig 22. Cracks and glassy phases(indicated by continues up to a steady-state condition that is produced by the Nicalon M/SiC specimen crept in argon at 1300oC and 45 MPa for competition between creating more bridged fibers as the crack h;(b)high-magnification image of the composite shown in Fig. 22(a) length increases and the fracture or creep of these fibers as more stress is transferred to them by the increased crack opening displacement. Therefore, the effects of free surfaces of the fiber and/or the matrix, which progressively tance of the bridge fibers on creep behavior of the composite closes the porosity and the access for oxygen toward the inter- are expected to be important phase and, consequently, stops the oxidation pro A special result is that the creep resistance of the H The annular porosity around the fibers decreases the Nicalon TM/SiC composite in air is higher than that in argon strength, which decreases as the gauge length increases. 13, 14 enhanced SiC/SiC composite II However, the opposite result is matrix produces a strong interface that is harmful in regard to observed for the creep of the standard SiC/SiC composite both strength and ductility. Therefore, annular porosity around (Figs. 13 and 14) he fibers and the silica formation on the free surfaces of the For the standard SiC/SiC composite, the strength and stress- fiber and/or the matrix both decrease the strength of the com- rupture life in air are always lower than those in argon or under posite. 20, 48-50 This observation is the primary reason for the vacuum, because of oxidation. 20,48-50 The oxidation of the lower creep and fatigue resistance of the standard SiC/SiC standard SiC/SiC composite includes two concurrent phenom- composite in air than in argon. However, the ena: oxidation of the pyrocarbon interphase, which creates an stability of Nicalon TM fiber at high temperature also should annular porosity around the fibers, and silica formation on the be considered. Fibers heat-treated in argon at high tempera-
MPa, matrix cracking can occur via stress redistribution due to creep of the fibers. (2) Effects of Fiber, Fiber Architecture, and Matrix The creep behavior of composites is dependent on the creep of the matrix, the fibers, the fiber architecture, and the interphases and interfaces between the fibers and the matrix. In the Hi-Nicalon™/SiC composite, the creep behavior of the enhanced SiC matrix is not known. The amount of additives in the matrix also is not known. Therefore, it is impossible to compare the creep resistance of the matrix to that of the fibers to determine the load-transfer direction during creep. However, all the testing stresses for creep of the Hi-Nicalon™/SiC composite in air are higher than the proportional limit, which can be assumed to be an approximate matrix cracking stress. The matrix microcracking occurs during the initial application of a creep load; therefore, fiber bridging of the matrix cracks always operates, whether the creep rate of the fibers is higher or lower than the matrix. In the crack-bridging mechanism, the matrix-crack growth rate is governed by a process in which the increase in the crack length is accompanied by an increase in the number of fibers that bridge the crack. This phenomenon continues up to a steady-state condition that is produced by the competition between creating more bridged fibers as the crack length increases and the fracture or creep of these fibers as more stress is transferred to them by the increased crackopening displacement. Therefore, the effects of creep resistance of the bridge fibers on creep behavior of the composite are expected to be important. A special result is that the creep resistance of the HiNicalon™/SiC composite in air is higher than that in argon (Figs. 6 and 7). This result is the same as the creep of the enhanced SiC/SiC composite.11 However, the opposite result is observed for the creep of the standard SiC/SiC composite (Figs. 13 and 14). For the standard SiC/SiC composite, the strength and stressrupture life in air are always lower than those in argon or under vacuum, because of oxidation.20,48–50 The oxidation of the standard SiC/SiC composite includes two concurrent phenomena: oxidation of the pyrocarbon interphase, which creates an annular porosity around the fibers, and silica formation on the free surfaces of the fiber and/or the matrix, which progressively closes the porosity and the access for oxygen toward the interphase and, consequently, stops the oxidation processes.51–54 The annular porosity around the fibers decreases the fiber strength, which decreases as the gauge length increases.13,14 The silica formation on the free surfaces of the fiber and/or the matrix produces a strong interface that is harmful in regard to both strength and ductility. Therefore, annular porosity around the fibers and the silica formation on the free surfaces of the fiber and/or the matrix both decrease the strength of the composite.20,48–50 This observation is the primary reason for the lower creep and fatigue resistance of the standard SiC/SiC composite in air than in argon. However, the thermodynamic stability of Nicalon™ fiber at high temperature also should be considered. Fibers heat-treated in argon at high temperaFig. 22. Cracks and glassy phases (indicated by arrows) in the HiNicalon™/SiC specimen crept in argon at 1300°C and 45 MPa for 65 h; (b) high-magnification image of the composite shown in Fig. 22(a). Fig. 21. Fracture surfaces of Hi-Nicalon™/SiC specimens creepfractured in air at 1300°C and 120 MPa, showing that a glassy phase covered the surface. January 1999 Creep and Fatigue Behavior in Hi-Nicalon™/SiC Composites at High Temperatures 125
126 Journal of the American Ceramic Society-Zhu et al Vol. 82. Ne 0210310436x104 08 06 06 0.4 0.4 02 0 10910110210310410510107 0 0.002 0.004 0006 0008 cle Strain 90M 05 MPa Fig. 23. Evolution of the hysteresis loops during fatigue of the iC composite at 1300C under a ma 12b lution position of the SiC, O, phase, which resulted in the e. the tures have shown severe degradation that is caused by the 12 严1.2 decom icon monoxide(Si ases. 29,30,55 whereas fibers treated in an oxiding environment oxygen gas, air) showed slightly less degradation. 55 Only 25% of the initial strength of the fibers was retained after the treat- ment at 1300@C in argon. s The instability of Nicalon TM fiber at 06 high temperature in argon may be the reason for the similar y reep resistances at low stresses in air and argon(Figs. 13 and 0.4 When the matrix is enl 90 MPa 76 MPa ticulate(borosilicate) that is capable of sealing the outer sur- face and the cracked surface of the specimen at elevated tem- peratures, the oxidation of the interphase becomes minor. Thus 0 instability of Hi-Nicalon M or(Nicalon TM) fiber at high tem- perature in argon is the key factor that causes lower creep and atigue resistance in argon than in air. moreover, the sealing of Time. s cracks by glass may be more effective in air than in argon because there are more oxygen molecules available in air to react with the glass-forming particulate than in argon of the Hi-Nicalon TM SiC composite is almost the same as that Creep of Hi-NicalonTM/SiC, 1300 C, Ar of the enhanced SiC/SiC composite in which the matrix is the same as in the Hi-Nicalon TM/SiC composite but the fiber is icalonTM fiber in air(Figs. 10-12), althoug NicalonTM fiber had a greater creep resistance (one order of magnitude) in comparison to normal NicalonTM fiber 3 This observation indicates that Hi-Nicalon TM fiber does not improv 08F 08 the creep and fatigue resistance of the composite in air, if the fference in fiber architectures(satin-woven structure in the Hi-NicalonTNSiC composite and plain-woven structure in the enhanced SiC/SiC composite) is neglected. In fact, the effects of fiber architecture on the creep and fatigue of the omposites may not be minor. The bending extent of fibers in 120 MPa 60 MPa the satin-woven structure is smaller than that in the plain- 90 MPa woven structure. This observation may lead to less damage in the fiber bundles and. therefore. is beneficial to 0 0 and fatigue properties. The Hi-Nicalon TM/SiC composite and the enhanced SiC/SiC composite have the same creep and Time. 6 atigue resistance because the creep and fatigue resistance of oth composites is controlled by the matrix rather than the ibers or the fiber architecture. This result also explains why ne creep resistance of the Hi-NicalonTM/SiC composite is Fig. 24. Elastic modulus normalized by the value of the modulus higher than that of the enhanced SiC/SiC composite in argon nder the first loading(EE versus (a) the number of cycles for (Figs. 8 and 9) fatigue of the Hi-Nicalon TM/SiC composite in air at 1300C under ( Creep and Fatigue Interaction The interaction of creep-fatigue of CMCs has not been given mum stresses, and (c)the time required for creep of the Hi-NicalonTM/ much attention over the years. Holmes 8 reported that the fa- SiC composite in argon at 1300 C under different maximum stresses
tures have shown severe degradation that is caused by the decomposition of the SiCxOy phase, which resulted in the evolution of carbon monoxide (CO) and silicon monoxide (SiO) gases,29,30,55 whereas fibers treated in an oxiding environment (oxygen gas, air) showed slightly less degradation.55 Only 25% of the initial strength of the fibers was retained after the treatment at 1300°C in argon.55 The instability of Nicalon™ fiber at high temperature in argon may be the reason for the similar creep resistances at low stresses in air and argon (Figs. 13 and 14). When the matrix is enhanced by adding glass-forming particulate (borosilicate) that is capable of sealing the outer surface and the cracked surface of the specimen at elevated temperatures, the oxidation of the interphase becomes minor. Thus, instability of Hi-Nicalon™ or (Nicalon™) fiber at high temperature in argon is the key factor that causes lower creep and fatigue resistance in argon than in air. Moreover, the sealing of cracks by glass may be more effective in air than in argon, because there are more oxygen molecules available in air to react with the glass-forming particulate than in argon. Another special result is that the creep and fatigue resistance of the Hi-Nicalon™/SiC composite is almost the same as that of the enhanced SiC/SiC composite in which the matrix is the same as in the Hi-Nicalon™/SiC composite but the fiber is normal Nicalon™ fiber in air (Figs. 10–12), although HiNicalon™ fiber had a greater creep resistance (one order of magnitude) in comparison to normal Nicalon™ fiber.31 This observation indicates that Hi-Nicalon™ fiber does not improve the creep and fatigue resistance of the composite in air, if the difference in fiber architectures (satin-woven structure in the Hi-Nicalon™/SiC composite and plain-woven structure in the enhanced SiC/SiC composite) is neglected. In fact, the effects of fiber architecture on the creep and fatigue of the composites may not be minor. The bending extent of fibers in the satin-woven structure is smaller than that in the plainwoven structure. This observation may lead to less damage in the fiber bundles and, therefore, is beneficial to the creep and fatigue properties. The Hi-Nicalon™/SiC composite and the enhanced SiC/SiC composite have the same creep and fatigue resistance because the creep and fatigue resistance of both composites is controlled by the matrix rather than the fibers or the fiber architecture. This result also explains why the creep resistance of the Hi-Nicalon™/SiC composite is higher than that of the enhanced SiC/SiC composite in argon (Figs. 8 and 9). (3) Creep and Fatigue Interaction The interaction of creep–fatigue of CMCs has not been given much attention over the years. Holmes18 reported that the faFig. 23. Evolution of the hysteresis loops during fatigue of the HiNicalon™/SiC composite at 1300°C under a maximum stress of 120 MPa in air. Fig. 24. Elastic modulus normalized by the value of the modulus under the first loading (E/E° ) versus (a) the number of cycles for fatigue of the Hi-Nicalon™/SiC composite in air at 1300°C under different maximum stresses, (b) the time required for creep of the Hi-Nicalon™/SiC composite in air at 1300°C under different maximum stresses, and (c) the time required for creep of the Hi-Nicalon™/ SiC composite in argon at 1300°C under different maximum stresses. 126 Journal of the American Ceramic Society—Zhu et al. Vol. 82, No. 1