COMPOSITES SCIENCE AND TECHNOLOGY ELSEVIER Composites Science and Technology 59(1999)833-851 Monotonic tension, fatigue and creep behavior of SiC-fiber reinforced SiC-matrix composites: a review S. Zhua. M. Mizuno b. Y Kagawa..Y. Mutoh Department of mechanical ing, Nagaoka University of Technology, Nagaoka, 940-21, Japan Center. Atsuta-ku. Ne 456, Japan Institute of Industrial Sciences, The University of Tokyo, Tokyo, 106, Japan Received 26 June 1997: received in revised form 18 June 1998; accepted 8 January 1999 Abstract The monotonic tension, fatigue and creep behaviour of Sic-fiber-reinforced Sic-matrix composites( SiC/Sic) has been reviewed Although the short-term properties of Sic/SiC at high temperatures are very desirable, fatigue and creep resistance at high tem- peratures in argon was much lower than at room temperature. Enhanced Sic/SiC exhibits excellent fatigue and creep properties in air, but the mechanisms are not well understood. The present Hi-Nicalon/SiC has similar properties to enhanced Sic/siC, but at higher cost Improvement of Hi-Nicalon/SiC therefore seems necessary for the development of a high-performance SiC/Sic mate- rial.C 1999 Elsevier Science Ltd. All rights reserved. Keywords: A. Ceramic-matrix composite; B. Fatigue: B Creep: SiC/SIC 1. Introduction behavior of CMCs to constituent properties. It is noted that inelastic strains eliminate stress concentrations and The main disadvantages of monolithic ceramics for scaling effects, enabling design procedures to be used structural components are their brittleness and low which are similar to those used for metals eliability. Even in monolithic ceramics with high To obtain high fracture toughness and thermal shock toughness, macroscopic inelasticity has not been resistance, CMCs were designed with a weak interface chieved. Design with such materials must be based on between fibers and matrix, e.g. the interface in SiC/SiC elastic stresses, combined with weakest-link scaling and composites where the fibers are coated with carbon. The extreme-value statistics weak interface can cause a crack to deflect along the Continuous-fiber-reinforced ceramic-matrix compo- interface, permitting intact fibers to bridge crack faces sites( CMCs)are specifically tailored so that crack-wake [1-3]. However, although the use of weak interfaces can processes result in materials with high fracture resis- increase fracture toughness and thermal shock resis- tance[1-10]. One feature of CMCs which is different tance [11], it is not compatible with creep and fatigue from fiber-reinforced plastics and metal-matrix compo- resistance at high temperature, which demands strong sites is that the failure strain of the matrix is lower than interfaces resisting the nucleation and growth of cavities that of the fiber. The distributed matrix cracking and [11] resultant fiber bridging redistribute stresses around In the last decade, cyclic fatigue and creep of continuous- train concentration sites and increase toughness an fiber-reinforced ceramic-matrix composites have been eliability. The stress for the initiation of matrix crack investigated [12-32] since these properties are very and the extent of various damage modes depend on the important for the application of CMCs. The fatigue interfacial bond strength, residual stresses, and sizes of the limit can be as high as 80% of the ultimate tensile stress pre-existing defects. Evans and his colleagues [10]reported at room temperature [19]. One of the mechanisms the methodology for relating the tensile constitutive responsible for the enhanced microstructural damage during fatigue appears to be related to the wear along Corresponding author. Tel:+81-3-3402-6231, ext 2436: fax: +81 the sliding fiber/matrix interface, which may lead to fiber damage(e.g. of a carbon fiber) and lower its failure 0266-3538/99/.see front matter o 1999 Elsevier Science Ltd. All rights reserved. PlI:S0266-3538(99)00014-7
Monotonic tension, fatigue and creep behavior of SiC-®berreinforced SiC-matrix composites: a review S. Zhua , M. Mizunob, Y. Kagawa c,*, Y. Mutoha a Department of Mechanical Engineering, Nagaoka University of Technology, Nagaoka, 940-21, Japan bJapan Fine Ceramics Center, Atsuta-ku, Nagoya, 456, Japan c Institute of Industrial Sciences, The University of Tokyo, Tokyo, 106, Japan Received 26 June 1997; received in revised form 18 June 1998; accepted 8 January 1999 Abstract The monotonic tension, fatigue and creep behaviour of SiC-®ber-reinforced SiC-matrix composites (SiC/SiC) has been reviewed. Although the short-term properties of SiC/SiC at high temperatures are very desirable, fatigue and creep resistance at high temperatures in argon was much lower than at room temperature. Enhanced SiC/SiC exhibits excellent fatigue and creep properties in air, but the mechanisms are not well understood. The present Hi-Nicalon/SiC has similar properties to enhanced SiC/SiC, but at higher cost. Improvement of Hi-Nicalon/SiC therefore seems necessary for the development of a high-performance SiC/SiC material. # 1999 Elsevier Science Ltd. All rights reserved. Keywords: A. Ceramic-matrix composite; B. Fatigue; B. Creep; SiC/SiC 1. Introduction The main disadvantages of monolithic ceramics for structural components are their brittleness and low reliability. Even in monolithic ceramics with high toughness, macroscopic inelasticity has not been achieved. Design with such materials must be based on elastic stresses, combined with weakest-link scaling and extreme-value statistics. Continuous-®ber-reinforced ceramic-matrix composites (CMCs) are speci®cally tailored so that crack-wake processes result in materials with high fracture resistance [1±10]. One feature of CMCs which is dierent from ®ber-reinforced plastics and metal-matrix composites is that the failure strain of the matrix is lower than that of the ®ber. The distributed matrix cracking and resultant ®ber bridging redistribute stresses around strain concentration sites and increase toughness and reliability. The stress for the initiation of matrix cracking and the extent of various damage modes depend on the interfacial bond strength, residual stresses, and sizes of the pre-existing defects. Evans and his colleagues [10] reported the methodology for relating the tensile constitutive behavior of CMCs to constituent properties. It is noted that inelastic strains eliminate stress concentrations and scaling eects, enabling design procedures to be used which are similar to those used for metals. To obtain high fracture toughness and thermal shock resistance, CMCs were designed with a weak interface between ®bers and matrix, e.g. the interface in SiC/SiC composites where the ®bers are coated with carbon. The weak interface can cause a crack to de¯ect along the interface, permitting intact ®bers to bridge crack faces [1±3]. However, although the use of weak interfaces can increase fracture toughness and thermal shock resistance [11], it is not compatible with creep and fatigue resistance at high temperature, which demands strong interfaces resisting the nucleation and growth of cavities [11]. In the last decade, cyclic fatigue and creep of continuous- ®ber-reinforced ceramic-matrix composites have been investigated [12±32] since these properties are very important for the application of CMCs. The fatigue limit can be as high as 80% of the ultimate tensile stress at room temperature [19]. One of the mechanisms responsible for the enhanced microstructural damage during fatigue appears to be related to the wear along the sliding ®ber/matrix interface, which may lead to ®ber damage (e.g. of a carbon ®ber) and lower its failure Composites Science and Technology 59 (1999) 833±851 0266-3538/99/$ - see front matter # 1999 Elsevier Science Ltd. All rights reserved. PII: S0266-3538(99)00014-7 * Corresponding author. Tel.:+81-3-3402-6231, ext.2436; fax:+81- 3-3402-6350
83 S. Zhu et al./ Composites Science and Technology 59(1999)833-857 stress[ 19,20]. The gradual damage growth is accompanied have become attractive ceramic-matrix composites and by a modulus decrease in CMCs under fatigue loading, have already been applied in some fields which has been studied in detail [12, 13, 16, 19, 22]. At high TEM observation and electron diffraction exhibited a temperature, the fatigue limit coincides with the propor- high degree of polycrystallinity, a relatively small crystal tional limit for unidirectional SiC/Si3 N4, tested under the size (3-4 nm)and a lack of preferred orientation for the conditions for which creep occurs [14] Sic component of the fibers in SiC/SiC [33]. The For creep behavior of CMCs, a classification was microstructure of Sic fibers does not change sig roposed [31, 32] in terms of creep mismatch ratio nificantly during fabrication of the composite. There are (CMR), defined as a ratio of the creep rate of the fiber two kinds of Sic grains in the matrix, polyhedral ones to that of the matrix. When CMr l, the main near to the fibers with a size ranging from 10 to 100 damage mechanism occurring during creep is periodic and columnar grains, further away from the fibers and fiber fracture and the creep behavior is governed by up to 500 nm in length. The interfacial layer between the embedded fibers. When CMR> 1, matrix microcrack- matrix and fibers is a carbon layer ing is the dominant mode of damage and creep behavior The carbon layer in Sic/Sic leads to low oxidation is governed by bridging fibers. However, the creep mis- resistance at intermediate and high temperatures [34-39 match ratio provides only a starting point for consider- A glass-forming, boron-based particulate material can ing creep behavior and damage mechanisms, since it be added to the matrix that reacts with oxygen to pro- only depends on the uniaxial creep behavior of the duce a sealant glass that inhibits oxidation of the carbon constituents. Because of local stress concentrations and layer [40]. This technology has been applied to SiC/SiC the development of triaxial stresses, the in-situ creep composites. The SiC/SiC modified in this way is referred behavior of the constituents in a composite can differ to as an enhanced Sic/Sic composite [40] significantly from the creep behavior measured during Since matrix microcracking may occur during the unconstrained uniaxial loading [31, 32 initial application of p load, fiber bridging of he incorporation of SiC fiber into Si3 N4 results in matrix cracks operates during the creep of standard substantial improvements in creep resistance [23-25]. SiC/SiC at high stresses, although the creep resistance of Multiple fiber fracture rather than multiple matrix SiC fibers is lower than that of the Sic matrix [41-491 cracking was observed and the creep mechanism was a This is undesirable for environmental resistance of the repetitive matrix stress relaxation-fiber rupture- load composites if they are exposed to air. Because creep of transfer and distribution scheme [23-25. Moreover, a the fibers controls matrix crack growth, increasing the threshold stress was found for the tensile creep of a creep resistance of the fibers should improve the creep unidirectional SiC/HPSN(hot-pressed silicon nitride) behavior of the composite. Hi-Nicalon'M fiber is one of composite, which was much higher than the propor- the improved SiC fibers [46, 49), which is used to rein- tional limit [23-25 force a SiC matrix(Hi-Nicalon/ SiC composite Different mechanisms were found in creep of SiC/ In this paper, monotonic tensile behavior, fracture lass-ceramics at 1200C at different tensile stress levels toughness and thermal shock resistance fatigue and [26]. At low stresses, cavities formed in the matrix with creep behavior of standard SiC/SiC, enhanced SiC/SiC no significant fiber or matrix damage [26]. At moderate and Hi-Nicalon TM SiC composites is reviewed stresses,periodic fiber rupture occurred, and at high stresses matrix fracture and rupture of the highly stres sed bridging fibers limited creep life [26]. Since grain 2. Monotonic Tension growth in Nicalon fibers enhanced creep resistance creep deformation was found to be transient in nature 2. 1. Monotonic tensile behavior in standard SiC/Sic at1200°CD27 Chemical vapor infiltration (CVI) is an important A characteristic in the stress versus strain curves of technique for manufacturing long-fiber-reinforced cera- SiC/SiC is the existence of inelastic deformation [50-60 mic-matrix composite, in which a porous preform of Fig. I shows the stress versus strain of a plain-weave fibers is infiltrated by a gaseous precursor which then 2DSiC/SiC composite at both room temperature and deposits a ceramic matrix. Although the feasibility of 1000oC in argon [55]. The room temperature curve CVI process has already been established for a number indicates a linear elastic behavior up to the proportiona of ceramic matrices including carbides(SiC, B4C, TiC), limit of 80 MPa, and this stress is about 40% of the nitrides(Bn, Si3 N4) and oxides(Al,O3, ZrO,), only Sic ultimate tensile strength (UTS). The modulus calculated matrix CVI composites are currently produced on an from the linear portion of the curve is 250 GPa. The industrial scale [2]. SiC (NicalonTM) fiber is one of the average values of UTS are 209 MPa at room tempera- most successful of the fine ceramic fibers, commercially ture and 251 MPa at 1000 C. It is noted that the UTS produced by Nippon Carbon. Therefore, NicalonTM. and the strain at UTS at 1000 C are higher than those fiber-reinforced silicon-carbide composites (SiC/SiC) at room temperature. The proportional limit and the
stress [19,20]. The gradual damage growth is accompanied by a modulus decrease in CMCs under fatigue loading, which has been studied in detail [12,13,16,19,22]. At high temperature, the fatigue limit coincides with the proportional limit for unidirectional SiCf/Si3N4, tested under the conditions for which creep occurs [14]. For creep behavior of CMCs, a classi®cation was proposed [31,32] in terms of creep mismatch ratio (CMR), de®ned as a ratio of the creep rate of the ®ber to that of the matrix. When CMR 1, matrix microcracking is the dominant mode of damage and creep behavior is governed by bridging ®bers. However, the creep mismatch ratio provides only a starting point for considering creep behavior and damage mechanisms, since it only depends on the uniaxial creep behavior of the constituents. Because of local stress concentrations and the development of triaxial stresses, the in-situ creep behavior of the constituents in a composite can dier signi®cantly from the creep behavior measured during unconstrained uniaxial loading [31,32]. The incorporation of SiC ®ber into Si3N4 results in substantial improvements in creep resistance [23±25]. Multiple ®ber fracture rather than multiple matrix cracking was observed and the creep mechanism was a repetitive matrix stress relaxation!®ber rupture! load transfer and distribution scheme [23±25]. Moreover, a threshold stress was found for the tensile creep of a unidirectional SiCf/HPSN (hot-pressed silicon nitride) composite, which was much higher than the proportional limit [23±25]. Dierent mechanisms were found in creep of SiCf/ glass-ceramics at 1200C at dierent tensile stress levels [26]. At low stresses, cavities formed in the matrix with no signi®cant ®ber or matrix damage [26]. At moderate stresses, periodic ®ber rupture occurred, and at high stresses matrix fracture and rupture of the highly stressed bridging ®bers limited creep life [26]. Since grain growth in NicalonTM ®bers enhanced creep resistance, creep deformation was found to be transient in nature at 1200C [27]. Chemical vapor in®ltration (CVI) is an important technique for manufacturing long-®ber-reinforced ceramic-matrix composite, in which a porous preform of ®bers is in®ltrated by a gaseous precursor which then deposits a ceramic matrix. Although the feasibility of CVI process has already been established for a number of ceramic matrices including carbides (SiC, B4C, TiC), nitrides (BN, Si3N4) and oxides (Al2O3, ZrO2), only SiC matrix CVI composites are currently produced on an industrial scale [2]. SiC (NicalonTM) ®ber is one of the most successful of the ®ne ceramic ®bers, commercially produced by Nippon Carbon. Therefore, NicalonTM- ®ber-reinforced silicon-carbide composites (SiC/SiC) have become attractive ceramic-matrix composites and have already been applied in some ®elds. TEM observation and electron diraction exhibited a high degree of polycrystallinity, a relatively small crystal size (3±4 nm) and a lack of preferred orientation for the SiC component of the ®bers in SiC/SiC [33]. The microstructure of SiC ®bers does not change signi®cantly during fabrication of the composite. There are two kinds of SiC grains in the matrix, polyhedral ones near to the ®bers with a size ranging from 10 to 100 nm, and columnar grains, further away from the ®bers and up to 500 nm in length. The interfacial layer between the matrix and ®bers is a carbon layer. The carbon layer in SiC/SiC leads to low oxidation resistance at intermediate and high temperatures [34±39]. A glass-forming, boron-based particulate material can be added to the matrix that reacts with oxygen to produce a sealant glass that inhibits oxidation of the carbon layer [40]. This technology has been applied to SiC/SiC composites. The SiC/SiC modi®ed in this way is referred to as an enhanced SiC/SiC composite [40]. Since matrix microcracking may occur during the initial application of a creep load, ®ber bridging of matrix cracks operates during the creep of standard SiC/SiC at high stresses, although the creep resistance of SiC ®bers is lower than that of the SiC matrix [41±49]. This is undesirable for environmental resistance of the composites if they are exposed to air. Because creep of the ®bers controls matrix crack growth, increasing the creep resistance of the ®bers should improve the creep behavior of the composite. Hi-NicalonTM ®ber is one of the improved SiC ®bers [46,49], which is used to reinforce a SiC matrix (Hi-NicalonTM/SiC composite). In this paper, monotonic tensile behavior, fracture toughness and thermal shock resistance, fatigue and creep behavior of standard SiC/SiC, enhanced SiC/SiC and Hi-NicalonTM/SiC composites is reviewed. 2. Monotonic Tension 2.1. Monotonic tensile behavior in standard SiC/SiC A characteristic in the stress versus strain curves of SiC/SiC is the existence of inelastic deformation [50±60]. Fig. 1 shows the stress versus strain of a plain-weave 2DSiC/SiC composite at both room temperature and 1000C in argon [55]. The room temperature curve indicates a linear elastic behavior up to the proportional limit of 80 MPa, and this stress is about 40% of the ultimate tensile strength (UTS). The modulus calculated from the linear portion of the curve is 250 GPa. The average values of UTS are 209 MPa at room temperature and 251 MPa at 1000C. It is noted that the UTS and the strain at UTS at 1000C are higher than those at room temperature. The proportional limit and the 834 S. Zhu et al. / Composites Science and Technology 59 (1999) 833±851
S. Zhu et al. / Composites Science and Technology 59(1999)833-851 835 Experimental Curve 250 RT 0000 (Ec=Ep+Er) 150 1000°c Estimated 0.0010.002 0.003 Unloading line i Strain Fig 1. Monotonic tensile stress-strain curves of 2D Sic/SiC compo- site at room temperature and 1000 C in argon with a dis rate of 0.5 mm/min. Fig. 2. Schematic diagram showing the permanent strain(Ep)and recoverable strain(Er) and total strain(es) modulus calculated from the linear portion of the curve at 1000C are 100 MPa and 260 GPa, respectively, slightly higher than those at room temperature. It was reported that the room-temperature tensile behavior is rate-dependent [51]: higher strain rates lead to a lower Youngs modulus, higher proportional limit and higher 700 UTS [51] 1000C 600 Evans et al [10] found that the tensile stress/strain curve for 2D composites was quite closely matched by simply scaling down the stress for the ID curve by 1/2 The matrix cracks formed in 90 bundles evolve at lower stresses than cracks in ID composites. These crack 200 extend laterally into the 0 bundles and, finally, the fibers carry the load prior to composite failure. If we use the fracture load at room temperature, the 6080100120140160180200220 stress on"dry"fibers in 0 bundles can be calculated by counting the number of oo bundles in the cross section Stress (MPa) of a specimen and using 500 fibers in every bundle. The Fig. 3. Permanent strain measured by partial unloading versus stress term"dry"means that the contribution of the matrix at room temperature and at 1000 C in argon and 90. bundles is not considered. The stress calculated on dry fibers in 0 bundles is 1. 1 GPa, which is lower than the strength(3.5 GPa) of original SiC fiber. This temperature, although the threshold stress for producing implies that there is either processing damage or non- a permanent strain (as with the proportional limit)at uniform stress and strain states on fibers in the speci- 1000C is higher. Since the permanent strain is caused mens, which reduces the fiber strength. By comparing primarily by damage in the composite, the degradation the fiber strength and the Weibull modulus with those of the composite at 1000 C is more sensitive to the stress prior to incorporation into SiC/SiC composites, Eckel than that at room temperature and Bradt [60] found that fiber damage could occur The monotonic tensile fracture surfaces of the com- either during the weaving or during another stage of posites at ambient and high temperature revealed that composite manufacture. The pores and 2D woven tensile fracture occurred both in 0 and in 90 bundles architecture might certainly lead to non-uniform stress at both ambient and high temperature. The fiber pull-out and strain fields under an applied load [19, 54 length in 0 bundles at high temperature is greater than The permanent strain, Ep, i. e unrecoverable strain as that at room temperature. The greater fiber pull-out defined in Fig. 2 and measured by a repeated loading- length at 1000 C may be the reason for the higher UTS unloading method, as a function of the stress is shown and strain at UTS than those at room temperature in Fig 3. It can be seen that the increase in permanent The increase in fiber pull-out length with temperature strain with stress at 1000C is faster than that at room in argon is the reverse of what is shown by the results in
modulus calculated from the linear portion of the curve at 1000C are 100 MPa and 260 GPa, respectively, slightly higher than those at room temperature. It was reported that the room-temperature tensile behavior is rate-dependent [51]; higher strain rates lead to a lower Young's modulus, higher proportional limit and higher UTS [51]. Evans et al [10] found that the tensile stress/strain curve for 2D composites was quite closely matched by simply scaling down the stress for the 1D curve by 1/2. The matrix cracks formed in 90 bundles evolve at lower stresses than cracks in 1D composites. These cracks extend laterally into the 0 bundles and, ®nally, the ®bers carry the load prior to composite failure. If we use the fracture load at room temperature, the stress on ``dry'' ®bers in 0 bundles can be calculated by counting the number of 0 bundles in the cross section of a specimen and using 500 ®bers in every bundle. The term ``dry'' means that the contribution of the matrix and 90 bundles is not considered. The stress calculated on dry ®bers in 0 bundles is 1.1 GPa, which is lower than the strength (3.5 GPa) of original SiC ®ber. This implies that there is either processing damage or nonuniform stress and strain states on ®bers in the specimens, which reduces the ®ber strength. By comparing the ®ber strength and the Weibull modulus with those prior to incorporation into SiC/SiC composites, Eckel and Bradt [60] found that ®ber damage could occur either during the weaving or during another stage of composite manufacture. The pores and 2D woven architecture might certainly lead to non-uniform stress and strain ®elds under an applied load [19,54]. The permanent strain, "P, i.e. unrecoverable strain as de®ned in Fig. 2 and measured by a repeated loading± unloading method, as a function of the stress is shown in Fig. 3. It can be seen that the increase in permanent strain with stress at 1000C is faster than that at room temperature, although the threshold stress for producing a permanent strain (as with the proportional limit) at 1000C is higher. Since the permanent strain is caused primarily by damage in the composite, the degradation of the composite at 1000C is more sensitive to the stress than that at room temperature. The monotonic tensile fracture surfaces of the composites at ambient and high temperature revealed that tensile fracture occurred both in 0 and in 90 bundles at both ambient and high temperature. The ®ber pull-out length in 0 bundles at high temperature is greater than that at room temperature. The greater ®ber pull-out length at 1000C may be the reason for the higher UTS and strain at UTS than those at room temperature. The increase in ®ber pull-out length with temperature in argon is the reverse of what is shown by the results in Fig. 1. Monotonic tensile stress±strain curves of 2D SiC/SiC composite at room temperature and 1000C in argon with a displacement rate of 0.5 mm/min. Fig. 2. Schematic diagram showing the permanent strain ("p) and recoverable strain ("r) and total strain ("c). Fig. 3. Permanent strain measured by partial unloading versus stress at room temperature and at 1000C in argon. S. Zhu et al. / Composites Science and Technology 59 (1999) 833±851 835
S. Zhu et al./ Composites Science and Technology 59(1999)833-857 air [35]. The fiber pull-out length decreases in air with UTS for SiC/SiC. However, high failure strains do not increasing temperature owing to a strong bonding at the always give high UTS for all CMCs. For example, some interface by oxidation at high temperature [34, 35]. The NicalonM/CAS specimens exhibit 1. 2%strain with shorter fiber pull-out in air at high temperature leads to strengths of 250 MPa. The composite strain after matrix decreased strength and ductility compared with room cracking ntially determined by the strain in the temperature intact fibers since very little stress is carried by the matrix. Thus the average composite strain is the average 2. 2. Interface and interphase effects strain of the intact fibers trix crack plane. The failure strains of nicalon and carbon/Sic are The interface between fiber and matrix in reality close to the failure strain fibers. However. the includes two interfaces. One is between the fiber and the failure strain of Sic/Sic is much lower than that of interphase, and the other is between the interphase and Nicalon fiber. This demonstrates that weak interfaces the matrix. The interphase in SiC/Sic is the carbon are not achieved in SiC/Sic during processing. There- layer between fiber and matrix, which plays a large role fore, in such a case an improvement of failure strain in tailoring the interface properties to optimise tough- may increase the UTS. a possible route toward ness, strength, and cyclic fatigue. The fracture resistance increasing the failure strain is that of further decreasing of SiC/Sic is currently promising, but the strength is the interface bond strength by means of improved not satisfactory, considering that it is only half the coating materials or coating processing, since debond- strength of the matrix and less than one tenth the ing is a prerequisite for interfacial sliding and the lower strength of the Sic fiber. sliding resistance leads to larger fiber pull-out length Although the effect of architecture and pores on Another route is to make finer SiC fibers so that more strength was emphasized in the last section, this cannot interfaces can be used to increase the fiber-bridging explain why the strength of 2D CVD carbon /Sic [19] is force behind the crack tip and fiber pull-out length wice that of 2D Sic/SiC with a similar volume percen- Certainly, increasing the strength of the fiber is also a tage of fibers. The strength and strain to failure of car- good way so that fiber pull-out occurs at a higher stress. bon fibers are 2.7 GPa and 1. 2%[1], respectively, which are similar to those of NicalonTM fibers(2.7 GPa and 23. Tensile behavior in enhanced SiC/Sic and 1.4%[ID. Although NicalonM fibers slowly degrade at Hi-Nicalon/M SiC 1000C over a period of several days during processing. the difference in strength and ductility between carbon The stress versus strain curves of Hi-Nicalon/SiC, and NicalonM fibers is not large enough to explain the enhanced SiC/SiC and standard SiC/SiC at 1300 C are different strength and ductility of carbon /SiC and Sic/ shown in Fig. 4. The enhanced SiC/SiC consists of SiC. The diameter of carbon fibers is 7-8 um and there Nicalon M fibers and the enhanced SiC matrix, which is are 1000 fibers per bundle in 2D carbon/ SiC composite. the same as in Hi-Nicalon M/SiC. The curve in Hi- Therefore, there is a larger interface area in carbon/ Sic NicalonM/SiC indicates linear elastic behavior up to than in SiC/SiC for a given volume fraction of fibers. the proportional limit of 70 MPa, and this stress Moreover, carbon /SiC has a residual stress state than about 30% of the ultimate tensile strength (UTS). The which is very different from that in SiC/SiC, with the carbon/Sic having an extensively microcracked matrix Moreover the carbon fiber does not bond to the sic as a Nicalon fiber will. As a result, the larger interface area, weaker interface bonding and more extensive matrix cracking lead to higher strength(400-500 MPa) ind strain to failure(0.8-1. 1%)in carbon/SiC [19] than hose in SiC/Sic 1300°Cc The UTS and failure strain in argon at 1000C are higher than those at RT in SiC/Sic, but the UTS and failure strain in air at high temperatures [35] are lower Enhanced SiC/SiC, Air than those at rt. this means that interface debonding is important to fiber bridging. If interface bonding is yery strong, stress concentration can cause fiber fracture and there are then no bridging fibers. Conversely, 0.002 0.006 0.008 decreasing the interface sliding resistance promotes Tensile Strain interface debonding, leading to long fiber pull-out and high strength and ductility. Therefore, we can see that a enhanced Sic/Sic in air and standard Sic/SiC composites in argon at large failure strain is generally accompanied by a high 1300.C
air [35]. The ®ber pull-out length decreases in air with increasing temperature owing to a strong bonding at the interface by oxidation at high temperature [34,35]. The shorter ®ber pull-out in air at high temperature leads to decreased strength and ductility compared with room temperature. 2.2. Interface and interphase eects The `interface' between ®ber and matrix in reality includes two interfaces. One is between the ®ber and the interphase, and the other is between the interphase and the matrix. The interphase in SiC/SiC is the carbon layer between ®ber and matrix, which plays a large role in tailoring the interface properties to optimise toughness, strength, and cyclic fatigue. The fracture resistance of SiC/SiC is currently promising, but the strength is not satisfactory, considering that it is only half the strength of the matrix and less than one tenth the strength of the SiC ®ber. Although the eect of architecture and pores on strength was emphasized in the last section, this cannot explain why the strength of 2D CVD carbon/SiC [19] is twice that of 2D SiC/SiC with a similar volume percentage of ®bers. The strength and strain to failure of carbon ®bers are 2.7 GPa and 1.2% [1], respectively, which are similar to those of NicalonTM ®bers (2.7 GPa and 1.4% [1]). Although NicalonTM ®bers slowly degrade at 1000C over a period of several days during processing, the dierence in strength and ductility between carbon and NicalonTM ®bers is not large enough to explain the dierent strength and ductility of carbon/SiC and SiC/ SiC. The diameter of carbon ®bers is 7±8 mm and there are 1000 ®bers per bundle in 2D carbon/SiC composite. Therefore, there is a larger interface area in carbon/SiC than in SiC/SiC for a given volume fraction of ®bers. Moreover, carbon/SiC has a residual stress state than which is very dierent from that in SiC/SiC, with the carbon/SiC having an extensively microcracked matrix. Moreover, the carbon ®ber does not bond to the SiC as a NicalonTM ®ber will. As a result, the larger interface area, weaker interface bonding and more extensive matrix cracking lead to higher strength (400±500 MPa) and strain to failure (0.8±1.1%) in carbon/SiC [19] than those in SiC/SiC. The UTS and failure strain in argon at 1000C are higher than those at RT in SiC/SiC, but the UTS and failure strain in air at high temperatures [35] are lower than those at RT. This means that interface debonding is important to ®ber bridging. If interface bonding is very strong, stress concentration can cause ®ber fracture and there are then no bridging ®bers. Conversely, decreasing the interface sliding resistance promotes interface debonding, leading to long ®ber pull-out and high strength and ductility. Therefore, we can see that a large failure strain is generally accompanied by a high UTS for SiC/SiC. However, high failure strains do not always give high UTS for all CMCs. For example, some NicalonTM/CAS specimens exhibit 1.2% strain with strengths of 250 MPa. The composite strain after matrix cracking is essentially determined by the strain in the intact ®bers since very little stress is carried by the matrix. Thus the average composite strain is the average strain of the intact ®bers at the matrix crack plane. The failure strains of NicalonTM/CAS and carbon/SiC are close to the failure strains of the ®bers. However, the failure strain of SiC/SiC is much lower than that of Nicalon ®ber. This demonstrates that weak interfaces are not achieved in SiC/SiC during processing. Therefore, in such a case an improvement of failure strain may increase the UTS. A possible route toward increasing the failure strain is that of further decreasing the interface bond strength by means of improved coating materials or coating processing, since debonding is a prerequisite for interfacial sliding and the lower sliding resistance leads to larger ®ber pull-out length. Another route is to make ®ner SiC ®bers so that more interfaces can be used to increase the ®ber-bridging force behind the crack tip and ®ber pull-out length. Certainly, increasing the strength of the ®ber is also a good way so that ®ber pull-out occurs at a higher stress. 2.3. Tensile behavior in enhanced SiC/SiC and Hi-NicalonTM/SiC The stress versus strain curves of Hi-NicalonTM/SiC, enhanced SiC/SiC and standard SiC/SiC at 1300C are shown in Fig. 4. The enhanced SiC/SiC consists of NicalonTM ®bers and the enhanced SiC matrix, which is the same as in Hi-NicalonTM/SiC. The curve in HiNicalonTM/SiC indicates linear elastic behavior up to the proportional limit of 70 MPa, and this stress is about 30% of the ultimate tensile strength (UTS). The Fig. 4. Tensile stress versus strain in Hi-NicalonTM/SiC in air, enhanced SiC/SiC in air and standard SiC/SiC composites in argon at 1300C. 836 S. Zhu et al. / Composites Science and Technology 59 (1999) 833±851
tes science and te 1999)833-851 UTS of Hi-NicalonTM/SiC is similar to that of standard point is important to designers, since they cannot design SiC/SiC and enhanced SiC/SiC, but the strains at UTs components for high-temperature use on basis only of of Hi-Nicalon TM/SiC and enhanced SiC/SiC are much monotonic strength higher than that of standard SiC/SiC. The fatigue limit of the composite at room tempera The modulus calculated from the linear portion of the ture is much higher than the stress for the matrix urve is about 140 GPa, which is higher than that (90 cracking. This means that the composites can avoid the GPa)of enhanced SiC/SiC and lower than that (200 unsteady propagation of the matrix cracks induced by GPa)of standard SiC/SiC at 1300oC the first loading during cyclic fatigue at stresses up to the fatigue limit. In other words, the matrix cracks formed on first loading may propagate and also micro- 3. Cyclic fatigue cracking of the matrix may continue over several thou- sand cycles [69], but they are finally arrested and remain 3.1. Fatigue behavior of standard SiC/Sic stationary during subsequent cyclic loading. Fiber brid- ging is commonly thought to be the main reason for thi Fatigue behavior of Sic/SiC composites has been phenomenon the bridging force decreases the investigated over a decade [55-68]. The cyclic fatigue life stress intensity at crack tip. At stresses above the fatigue at room temperature and 1000oC in argon is shown in limit, cyclic fatigue fracture occurs accompanied by a Fig. 5. The stress-life curve at 1000oC can be divided modulus reduction with cycles [56, 62]. This reduction of into three regimes. One is the low cycle regime(10 cycles )at 1000 C mainly are higher than those at room temperature(Fig. 1), the occurred in 0 bundles, although one or two 90 bundle fatigue limit at 1000 C is much lower than that at room fail. This fracture morphology is associated with 2D temperature. This means that the fatigue limit is not woven structures with a large amount of pores proportional to the monotonic tensile strength. This There are three kinds of fracture modes of o bundles s schematically shown in Fig. 7. The first is that in which fracture of 0o bundles results from crack propa- gation in 90 bundles. In the second mode fracture occurs at a crossover point of 0%/90 bundles, and in the 1000°cin ● Rt in Air third the fracture is caused by shearing in the middle of 0 bundle owing to two cracks originating at the two nearest crossover points. The monotonic tension and cyclic fatigue at high stresses are composed largely of the first kind of fracture mode and occasionally, locally of the second kind. The cyclic fatigue at low stresses at 1000C is primarily of the second and third fracture 100 types, and in a few places of the first kind. Clusters of fibers with similar pull-out length in some 0 bundles 01102103104105106107108 sometimes exist in cyclic fatigue at 1000%C, when they fracture by the third kind of fracture mode. The change Cycles to Failure from stage I to II for fatigue at 1000 C(Fig. 5)agrees Fig. 5. Maximum tensile stress versus cycles to failure of 2D Sic/Sic with the fracture mode change(Figs. 6 and 7), which erature and1000°C. seems related to the maximum stress
UTS of Hi-NicalonTM/SiC is similar to that of standard SiC/SiC and enhanced SiC/SiC, but the strains at UTS of Hi-NicalonTM/SiC and enhanced SiC/SiC are much higher than that of standard SiC/SiC. The modulus calculated from the linear portion of the curve is about 140 GPa, which is higher than that (90 GPa) of enhanced SiC/SiC and lower than that (200 GPa) of standard SiC/SiC at 1300C. 3. Cyclic fatigue 3.1. Fatigue behavior of standard SiC/SiC Fatigue behavior of SiC/SiC composites has been investigated over a decade [55±68]. The cyclic fatigue life at room temperature and 1000C in argon is shown in Fig. 5. The stress±life curve at 1000C can be divided into three reÂgimes. One is the low cycle reÂgime ( 104 cycles) at 1000C mainly occurred in 0 bundles, although one or two 90 bundles fail. This fracture morphology is associated with 2D woven structures with a large amount of pores. There are three kinds of fracture modes of 0 bundles, as schematically shown in Fig. 7. The ®rst is that in which fracture of 0 bundles results from crack propagation in 90 bundles. In the second mode fracture occurs at a crossover point of 0/90 bundles, and in the third the fracture is caused by shearing in the middle of 0 bundle owing to two cracks originating at the two nearest crossover points. The monotonic tension and cyclic fatigue at high stresses are composed largely of the ®rst kind of fracture mode and occasionally, locally, of the second kind. The cyclic fatigue at low stresses at 1000C is primarily of the second and third fracture types, and in a few places of the ®rst kind. Clusters of ®bers with similar pull-out length in some 0 bundles sometimes exist in cyclic fatigue at 1000C, when they fracture by the third kind of fracture mode. The change from stage I to II for fatigue at 1000C (Fig. 5) agrees with the fracture mode change (Figs. 6 and 7), which seems related to the maximum stress. Fig. 5. Maximum tensile stress versus cycles to failure of 2D SiC/SiC composite at room temperature and 1000C. S. Zhu et al. / Composites Science and Technology 59 (1999) 833±851 837
S. Zhu et al. /Composites Science and Technology 59(1999)833-851 750um 750pm 200pm Fig. 6. Fracture surfaces of cyclic fatigue specimens of Sic/SiC composites(4.3x10- cycles) at room temperat 10 Hz and (2.7x105 cycles)at 1000 C at 93.7 MPa in argon at 20 Hz. The load ratio is 0.1 for both temperatures. (a)RT, (b)1000oC, (c)RT, ( 1000°C. 3. The fiber pull-out engin or cyclic fatigue is higher length of the broken fiber. The fiber pull-out length is than that of monotonic tension at the same tempera- assumed equal to Lc/2. The sliding resistance of the ture. The sliding resistance of interface(Ti), can be cal- interface decreases with cyclic fatigue or increasing culated by the equation [55] temperature, when the fiber strength is assumed to be ti=ord/2L (1) constant (1. 1 GPa for Nicalon fiber). This is consistent with the experimental results on effects of temperature on where ar is the fiber strength, d is the diameter of the cyclic fatigue of Sic/SiC [63]. However, Eq (1)can only fiber(14um in the present SiC/SiC)and Lc is the shortest give a qualitative indication of the interfacial sliding
The ®ber pull-out length of cyclic fatigue is higher than that of monotonic tension at the same temperature. The sliding resistance of interface i, can be calculated by the equation [55] i fd=2Lc 1 where f is the ®ber strength, d is the diameter of the ®ber (14mm in the present SiC/SiC) and Lc is the shortest length of the broken ®ber. The ®ber pull-out length is assumed equal to Lc=2. The sliding resistance of the interface decreases with cyclic fatigue or increasing temperature, when the ®ber strength is assumed to be constant (1.1 GPa for NicalonTM ®ber). This is consistent with the experimental results on eects of temperature on cyclic fatigue of SiC/SiC [63]. However, Eq. (1) can only give a qualitative indication of the interfacial sliding Fig. 6. Fracture surfaces of cyclic fatigue specimens of SiC/SiC composites (4.3102 cycles) at room temperature at 180 MPa in air at a frequency of 10 Hz and (2.7105 cycles) at 1000C at 93.7 MPa in argon at 20 Hz. The load ratio is 0.1 for both temperatures. (a) RT, (b) 1000C, (c) RT, (d) 1000C. 838 S. Zhu et al. / Composites Science and Technology 59 (1999) 833±851
S. Zhu et al. /Composites Science and Te 9(1999)833-851 (C) Fig. 7. Schematic diag f fracture modes of oo bundles resistance, because the strength of the fibers depends on the gage length and the pull-out length is not exactly jual to Lc/2 Cyclic loading and unloading result in cyclic opening and closing of the matrix cracks, which leads to repe- ated shear stress and slip of interfaces between fiber and matrix in 0 bundles. This process promotes debonding nd increases the length of the debonded interface. The long debonded interface leads to long fiber pull-out Moreover, the repeated slipping of interface may cause the interphase (carbon coating layer)to fail and even the surface of fibers to wear Crushed fibers and debris were found on cyclic fatigue fracture surfaces at both room and high temperatures [55]. Two kinds of debris consist of micrometer-sized Sic debris particles and submicrometer-sized graphite dusts [20]. Similar phe- nomena were found in the present experiments. This is evidence of interface damage during cyclic fatigue 3. 2. Crack initiation and distribution Both observation on interrupted test specimens tested at stresses slightly higher than the proportional limit and in situ observation show that most of the cracks initiate at he sharp corners of large pores at the crossover points of the fiber bundle weave as shown in Fig. 8. The cracks are of three kinds cracks in o bundles cracks in 90 bun dles and cracks at crossover points of the weave. The percentage of cracks is greatest in 90 bundles followed by crossover points of 00/90 bundles for monotonic tension and cyclic fatigue at high stresses at both room and high temperatures. However, most cracks form in 0 bundles and then at crossover points of 0/90 bun dles for cyclic fatigue at low stresses at 1000C. There are few matrix cracks in the o fiber bundles in the rt tensile tested specimen The dominant damage mode changes from cracks in loading. (a)at 1000@ C and 118.8 MPa (6.9x10 cycles); (b) at room 90 bundles for cyclic fatigue at high stresses to cracking temperature and 170 MPa(3.3x 106 cycles)
resistance, because the strength of the ®bers depends on the gage length and the pull-out length is not exactly equal to Lc=2. Cyclic loading and unloading result in cyclic opening and closing of the matrix cracks, which leads to repeated shear stress and slip of interfaces between ®ber and matrix in 0 bundles. This process promotes debonding and increases the length of the debonded interface. The long debonded interface leads to long ®ber pull-out. Moreover, the repeated slipping of interface may cause the interphase (carbon coating layer) to fail and even the surface of ®bers to wear. Crushed ®bers and debris were found on cyclic fatigue fracture surfaces at both room and high temperatures [55]. Two kinds of debris consist of micrometer-sized SiC debris particles and submicrometer-sized graphite dusts [20]. Similar phenomena were found in the present experiments. This is evidence of interface damage during cyclic fatigue. 3.2. Crack initiation and distribution Both observation on interrupted test specimens tested at stresses slightly higher than the proportional limit and in situ observation show that most of the cracks initiate at the sharp corners of large pores at the crossover points of the ®ber bundle weave, as shown in Fig. 8. The cracks are of three kinds: cracks in 0 bundles, cracks in 90 bundles and cracks at crossover points of the weave. The percentage of cracks is greatest in 90 bundles followed by crossover points of 0/90 bundles for monotonic tension and cyclic fatigue at high stresses at both room and high temperatures. However, most cracks form in 0 bundles and then at crossover points of 0/90 bundles for cyclic fatigue at low stresses at 1000C. There are few matrix cracks in the 0 ®ber bundles in the RT tensile tested specimen. The dominant damage mode changes from cracks in 90 bundles for cyclic fatigue at high stresses to cracking Fig. 7. Schematic diagram of fracture modes of 0 bundles. Fig. 8. Crack initiation sites and crack growth paths under cyclic loading. (a) at 1000C and 118.8 MPa (6.9104 cycles); (b) at room temperature and 170 MPa (3.3106 cycles). S. Zhu et al. / Composites Science and Technology 59 (1999) 833±851 839
S. Zhu et al. /Composites Science and Technology 59(1999)833-851 in 0 bundles at crossover points initiated from large 3.3. Crack propagation and connection pores for cyclic fatigue at low stresses at 1000%C.This implies that there may exist a critical stress for the Cracks first propagate in the matrix or along transition of the dominant damage mode. At high faces between the matrix and the fibers in the 90 bun- stresses,matrix cracking is considered to be saturated dles and then interact with predominantly intact fibers during first loading. Since extensive matrix cracks and in the 0o bundles, as shown in Fig 8. The debonding of debonding of interfaces decrease stress concentrations, interfaces occurs in 0 bundles. Both the pores and the he 90 bundles become easy paths for crack propagation. interfaces cause cracks to deflect. The propagating However, at low stress, matrix cracking is unsaturated cracks sometimes stop at the pores If there is another and therefo trix cracking continues during cyclic crack initiated at the pore in the direction of crack pro- loading. Stress concentration points at pores take an pagation, crack interconnection occurs. The interfaces important role for fracture path orientation. More- in 90 bundles promote crack propagation. However, over, creep of the 0o fibers and the reduced interfacial the interfaces in 0o bundles can deflect cracks by sliding resistance also promote fracture in the 0 bun debonding along the interfaces between fibers and dles at high temperatures. Therefore, the saturation matrix. Debonding is a prerequisite for fiber-bridging stress for matrix cracking may be the critical stress for and fiber-pull-out. When the transverse cracks propa the transition of fracture mode. This will be studied in gate into 0 bundles, they may occasionally cut some the future fibers which are strongly bonded with the matrix so that Both the pores and the architecture are important for no debonding occurs, and then propagate along the crack initiation and propagation. As a consequence of weak, debonded interfaces in the middle of 0 bundles the woven 2D fiber architecture. non-uniform stresses At room temperature, the first loading leads to and strains develop in the axial, width and thickness cracking and at the same time leaves intact 0 fibers in the directions of the specimen. Because of the non-uni- wake of the cracks. If the stress remains constant, the formity in stress and strain near the crossover points of bridging force keeps the cracks stationary. Therefore, no the fiber bundles, there is extensive fracture and spalla- stress rupture occurs at room temperature [64]. How on of the SiC matrix at these locations [19]. In the ever, under cyclic loading the debonded interface sliding specimens, cracks commonly initiated at the sharp cor- resistance decreases, which in turn promotes interface ners of large pores at the crossover point debonding further. The average fiber strength depends The non-uniform stress and strain distributions on the length tested: longer fibers are more likely to develop at different scales along three orthogonal direc- have more detrimental, strength-limiting defects and so tions [ 19]. At the macroscopic scale, the 2D fiber bundle are weaker. Therefore, the decrease of the interface configuration produces a non-uniform stress distribu- sliding resistance can reduce the fiber bridging stress tion along the axial and width directions within indivi- and is a dominant cyclic fatigue damage mechanism in dual plies; the random stacking of the plies produces a SiC/SiC [62] non-uniform stress distribution in the thickness direc t high temperature, creep of fibers can relax the fiber tion of the specimen. Non-uniform stretching of plies bridging forces and promote subcritical crack propaga produces a shear stress and contact pressure at the con- tion. Moreover, relaxation of the residual stress at the act points between them. This leads to delamination interface between fibers and matrix at 1000C can between plies. Within individual fiber bundles, cracking decrease frictional stress of the interface. The time- of the matrix occurs as the 0 bundles attempt to align dependent decrease of the fiber bridging stress promotes with the tensile loading direction. Non-aligned fibers crack propagation. It was found that the creep fracture suppress pull-out by virtue of bending strains. The life also sharply decreased at 1000@C like cyclic fatigue longest pull-out length in 2D composite is the width at the stresses below 160 MPa [73] between the two 90 bundles A random distribution of 0 fiber failure under fati Evans et al. [67] reported that the strength of 2D gue loading and bundle rupture near the crossover SiC/SiC was half of the strength of unidirectional points as found at room temperature in a carbon/SiC NicalonM/SiC and pointed out that hence the beha- composite [19]. This differs from the present results on vior of 2D materials must be dominated by the 0o bun- cyclic fatigue at room temperature. However, fracture dles. In fact, it is not only the effect of 90 bundles to by cyclic fatigue at 1000C shows a similar distribution reduce the strength by half, since the architecture and of 0 bundle failure. The delamination cracking was he pores strongly decrease the strength too. As a result, severe under room temperature fatigue in carbon/ SiC the stress on 0% fibers, calculated from the maximum composite [19]. Evident delamination was also found load over the cross section area of 0 fibers without under high temperature cyclic fatigue in SiC/SiC com considering the contribution of 90 fibers and the posite. Considering carbon fibers do not bond to the matrix, is lower than the strength of the fibers(as men- SiC matrix like NicalonTM fibers, the interface sliding tioned in Section 2.1) resistance of carbon/Sic is expected to be lower than
in 0 bundles at crossover points initiated from large pores for cyclic fatigue at low stresses at 1000C. This implies that there may exist a critical stress for the transition of the dominant damage mode. At high stresses, matrix cracking is considered to be saturated during ®rst loading. Since extensive matrix cracks and debonding of interfaces decrease stress concentrations, the 90 bundles become easy paths for crack propagation. However, at low stress, matrix cracking is unsaturated and therefore matrix cracking continues during cyclic loading. Stress concentration points at pores take an important role for fracture path orientation. Moreover, creep of the 0 ®bers and the reduced interfacial sliding resistance also promote fracture in the 0 bundles at high temperatures. Therefore, the saturation stress for matrix cracking may be the critical stress for the transition of fracture mode. This will be studied in the future. Both the pores and the architecture are important for crack initiation and propagation. As a consequence of the woven 2D ®ber architecture, non-uniform stresses and strains develop in the axial, width and thickness directions of the specimen. Because of the non-uniformity in stress and strain near the crossover points of the ®ber bundles, there is extensive fracture and spallation of the SiC matrix at these locations [19]. In the specimens, cracks commonly initiated at the sharp corners of large pores at the crossover points. The non-uniform stress and strain distributions develop at dierent scales along three orthogonal directions [19]. At the macroscopic scale, the 2D ®ber bundle con®guration produces a non-uniform stress distribution along the axial and width directions within individual plies; the random stacking of the plies produces a non-uniform stress distribution in the thickness direction of the specimen. Non-uniform stretching of plies produces a shear stress and contact pressure at the contact points between them. This leads to delamination between plies. Within individual ®ber bundles, cracking of the matrix occurs as the 0 bundles attempt to align with the tensile loading direction. Non-aligned ®bers suppress pull-out by virtue of bending strains. The longest pull-out length in 2D composite is the width between the two 90 bundles. Evans et al. [67] reported that the strength of 2D SiC/SiC was half of the strength of unidirectional NicalonTM/SiC and pointed out that hence the behavior of 2D materials must be dominated by the 0 bundles. In fact, it is not only the eect of 90 bundles to reduce the strength by half, since the architecture and the pores strongly decrease the strength too. As a result, the stress on 0 ®bers, calculated from the maximum load over the cross section area of 0 ®bers without considering the contribution of 90 ®bers and the matrix, is lower than the strength of the ®bers (as mentioned in Section 2.1). 3.3. Crack propagation and connection Cracks ®rst propagate in the matrix or along interfaces between the matrix and the ®bers in the 90 bundles and then interact with predominantly intact ®bers in the 0 bundles, as shown in Fig. 8. The debonding of interfaces occurs in 0 bundles. Both the pores and the interfaces cause cracks to de¯ect. The propagating cracks sometimes stop at the pores. If there is another crack initiated at the pore in the direction of crack propagation, crack interconnection occurs. The interfaces in 90 bundles promote crack propagation. However, the interfaces in 0 bundles can de¯ect cracks by debonding along the interfaces between ®bers and matrix. Debonding is a prerequisite for ®ber-bridging and ®ber-pull-out. When the transverse cracks propagate into 0 bundles, they may occasionally cut some ®bers which are strongly bonded with the matrix so that no debonding occurs, and then propagate along the weak, debonded interfaces in the middle of 0 bundles. At room temperature, the ®rst loading leads to cracking and at the same time leaves intact 0 ®bers in the wake of the cracks. If the stress remains constant, the bridging force keeps the cracks stationary. Therefore, no stress rupture occurs at room temperature [64]. However, under cyclic loading the debonded interface sliding resistance decreases, which in turn promotes interface debonding further. The average ®ber strength depends on the length tested: longer ®bers are more likely to have more detrimental, strength-limiting defects and so are weaker. Therefore, the decrease of the interface sliding resistance can reduce the ®ber bridging stress and is a dominant cyclic fatigue damage mechanism in SiC/SiC [62]. At high temperature, creep of ®bers can relax the ®ber bridging forces and promote subcritical crack propagation. Moreover, relaxation of the residual stress at the interface between ®bers and matrix at 1000C can decrease frictional stress of the interface. The timedependent decrease of the ®ber bridging stress promotes crack propagation. It was found that the creep fracture life also sharply decreased at 1000C like cyclic fatigue at the stresses below 160 MPa [73]. A random distribution of 0 ®ber failure under fatigue loading and bundle rupture near the crossover points as found at room temperature in a carbon/SiC composite [19]. This diers from the present results on cyclic fatigue at room temperature. However, fracture by cyclic fatigue at 1000C shows a similar distribution of 0 bundle failure. The delamination cracking was severe under room temperature fatigue in carbon/SiC composite [19]. Evident delamination was also found under high temperature cyclic fatigue in SiC/SiC composite. Considering carbon ®bers do not bond to the SiC matrix like NicalonTM ®bers, the interface sliding resistance of carbon/SiC is expected to be lower than 840 S. Zhu et al. / Composites Science and Technology 59 (1999) 833±851
S. Zhu et al./ Composites Science and Technology 59(1999)833-857 that of SiC/SiC at room temperature. The interface reduce the thermally induced stress. Since CVI proces- liding resistance decreases with increasing temperature. sing temperature is about 1000C, the thermally induced Therefore, the fatigue fracture morphology of Sic/Sic stress at 1000oC will be greatly decreased at 1000 C is similar to that of carbon/SiC at room tem- The observation of cracks on two interrupted speci perature. This demonstrates that the interface sliding mens(at 96 and at 187 MPa in tension at room tem- resistance of SiC/SiC markedly influences the evolution perature)showed that only small cracks initiated at the of fatigue damage at high temperature large pores at 96 MPa and long cracks formed in 90 bundles at 187 MPa, which is above the fatigue limit at 3.4. Cyclic fatigue mechanisms room temperature. Therefore, at room temperature cyclic fatigue is controlled by crack propagation. At the Recently, Evans et al. [67] reviewed fatigue of ceramic- stress below the fatigue limit, the driving force is not matrix composites at room temperature. Two possible enough to cause crack propagation. However, at mechanisms for CMCs were proposed [67]:(1)changes 1000 C cracks can propagate since the driving force is in the interface sliding resistance during cycling;(2) increased by reduction of fiber-bridging stress degradation of the strength of the fibers by cyclic sliding DiCarlo [40] pointed out that creep onset of Nica along the interface by means of an abrasion mechanism, lonM fibers was at 900 C. Therefore, fiber creep is an which introduces flaws in the fibers important factor for considering fatigue mechanism at The specimens fatigued for 107 cycles at the peak 1000oC. Henager and Jones[75-78] have shown that the ress of 160 MPa at room temperature have the same creep-induced stress relaxation of crack-bridging Nica- strength as the original specimens. The reason is that lonM fibers can account for the crack growth rate in the development of many matrix cracks and fiber/ CVI SiC reinforced with Nicalon TM fibers at 1100 C matrix debonding effectively reduce the number of stress However, creep data of Nicalon fibers at 1000C are particular crack extension. Fron gy associated with any very limited, not enough to be used for a prediction of concentration points and the ener this point, the 2D SiC/ creep relaxation of fibers in SiC/SiC Sic composite is considered to be a fatigue-resistant At room temperature, the fatigue resistance of SiC/ material at room temperature Sic or other CMCs is good with a fatigue limit of about At 1000 C, however, the residual strength of the spe- 80% UTS. However, at high temperatures, even if the cimens fatigued for 10 cycles at the peak stress of 75 temperature is not high enough to cause creep of fibers, MPa is lower than that of original specimens, although fatigue resistance of the composites was decreased [20] the difference is not great. This means that damage The present results show a similar phenomenon: fatigue mechanism during cyclic fatigue at 1000C is different resistance of SiC/Sic was much decreased at 1000C. As rom that at room temperature. At high temperature, discussed in the last section, fiber creep and decreased effects of oxidation, creep and sliding resistance of the sliding resistance of interface were causing the decreased interface should be considered. The creep rate of the Sic fatigue resistance. The lower sliding resistance of inter matrix at 1000C is very low and little oxidation occurs face is good for the failure strain and UTS under in argon. Therefore, the cyclic fatigue mechanisms of monotonic loading, but inferior to cyclic fatigue resis Sic/SiC at 1000 C will be explored only by the effects of tance. This is why UTS and failure strain at 1000C are fiber creep and the interfacial sliding resistance in the higher those at room temperature and the fatigue limit at 1000.c is lower than that at room temperature The friction stress at the interface is given by [55] Therefore, a balance should be made for monotonic properties and for cyclic fatigue resistance when ≈ designing or modifying 2D-woven CMCs where u is the coefficient of sliding friction, or is the 3.5. Fatigue of enhanced SiC/Sic and Hi-Nicalon/ radial thermally induced stress, og is the radial stress caused by roughness of interface, op is the radial stress arising from the difference in Poissons ratios between The cyclic fatigue life versus the maximum stress of fiber and matrix Hi-Nicalon/SiC at 1300.C in air is almost the same The thermally induced stress caused by thermal as that of enhanced Sic/SiC in air [57, 58], but longer expansion mismatch has an important role in both the than that of standard SiC/Sic in air and argon at formation and the effectiveness of fiber bridging. The 1300C [56](Fig 9) thermal expansion coefficient (3. 1x10-6 K-)of Fig. 10 shows the evolution of the stress strain fiber is lower than that (4.8x10-6))of the esis loops. The slope decreases and the width of the matrix [74]. As a result, a compressive residual increases with cycles. The former indicates the de ( fiber-clamping stress) is produced after processing at of the modulus and the latter means the decrease of the high temperature Increasing temperature is expected to interfacial sliding resistance. The hysteresis loops move
that of SiC/SiC at room temperature. The interface sliding resistance decreases with increasing temperature. Therefore, the fatigue fracture morphology of SiC/SiC at 1000C is similar to that of carbon/SiC at room temperature. This demonstrates that the interface sliding resistance of SiC/SiC markedly in¯uences the evolution of fatigue damage at high temperature. 3.4. Cyclic fatigue mechanisms Recently, Evans et al. [67] reviewed fatigue of ceramicmatrix composites at room temperature. Two possible mechanisms for CMCs were proposed [67]: (1) changes in the interface sliding resistance during cycling; (2) degradation of the strength of the ®bers by cyclic sliding along the interface by means of an abrasion mechanism, which introduces ¯aws in the ®bers. The specimens fatigued for 107 cycles at the peak stress of 160 MPa at room temperature have the same strength as the original specimens. The reason is that the development of many matrix cracks and ®ber/ matrix debonding eectively reduce the number of stress concentration points and the energy associated with any particular crack extension. From this point, the 2D SiC/ SiC composite is considered to be a fatigue-resistant material at room temperature. At 1000C, however, the residual strength of the specimens fatigued for 107 cycles at the peak stress of 75 MPa is lower than that of original specimens, although the dierence is not great. This means that damage mechanism during cyclic fatigue at 1000C is dierent from that at room temperature. At high temperature, eects of oxidation, creep and sliding resistance of the interface should be considered. The creep rate of the SiC matrix at 1000C is very low and little oxidation occurs in argon. Therefore, the cyclic fatigue mechanisms of SiC/SiC at 1000C will be explored only by the eects of ®ber creep and the interfacial sliding resistance in the following. The friction stress at the interface is given by [55] i T R P 2 where is the coecient of sliding friction, T is the radial thermally induced stress, R is the radial stress caused by roughness of interface, P is the radial stress arising from the dierence in Poisson's ratios between ®ber and matrix. The thermally induced stress caused by thermal expansion mismatch has an important role in both the formation and the eectiveness of ®ber bridging. The thermal expansion coecient (3.110ÿ6 Kÿ1 ) of SiC ®ber is lower than that (4.810ÿ6 Kÿ1 ) of the SiC matrix [74]. As a result, a compressive residual stress (®ber-clamping stress) is produced after processing at high temperature. Increasing temperature is expected to reduce the thermally induced stress. Since CVI processing temperature is about 1000C, the thermally induced stress at 1000C will be greatly decreased. The observation of cracks on two interrupted specimens (at 96 and at 187 MPa in tension at room temperature) showed that only small cracks initiated at the large pores at 96 MPa and long cracks formed in 90 bundles at 187 MPa, which is above the fatigue limit at room temperature. Therefore, at room temperature cyclic fatigue is controlled by crack propagation. At the stress below the fatigue limit, the driving force is not enough to cause crack propagation. However, at 1000C cracks can propagate since the driving force is increased by reduction of ®ber-bridging stress. DiCarlo [40] pointed out that creep onset of NicalonTM ®bers was at 900C. Therefore, ®ber creep is an important factor for considering fatigue mechanism at 1000C. Henager and Jones [75±78] have shown that the creep-induced stress relaxation of crack-bridging NicalonTM ®bers can account for the crack growth rate in CVI SiC reinforced with NicalonTM ®bers at 1100C. However, creep data of NicalonTM ®bers at 1000C are very limited, not enough to be used for a prediction of creep relaxation of ®bers in SiC/SiC. At room temperature, the fatigue resistance of SiC/ SiC or other CMCs is good with a fatigue limit of about 80% UTS. However, at high temperatures, even if the temperature is not high enough to cause creep of ®bers, fatigue resistance of the composites was decreased [20]. The present results show a similar phenomenon: fatigue resistance of SiC/SiC was much decreased at 1000C. As discussed in the last section, ®ber creep and decreased sliding resistance of interface were causing the decreased fatigue resistance. The lower sliding resistance of interface is good for the failure strain and UTS under monotonic loading, but inferior to cyclic fatigue resistance. This is why UTS and failure strain at 1000C are higher those at room temperature and the fatigue limit at 1000C is lower than that at room temperature. Therefore, a balance should be made for monotonic properties and for cyclic fatigue resistance when designing or modifying 2D-woven CMCs. 3.5. Fatigue of enhanced SiC/SiC and Hi-NicalonTM/ SiC The cyclic fatigue life versus the maximum stress of Hi-NicalonTM/SiC at 1300C in air is almost the same as that of enhanced SiC/SiC in air [57,58], but longer than that of standard SiC/SiC in air and argon at 1300C [56] (Fig. 9). Fig. 10 shows the evolution of the stress strain hysteresis loops. The slope decreases and the width of the loops increases with cycles. The former indicates the decrease of the modulus and the latter means the decrease of the interfacial sliding resistance. The hysteresis loops move S. Zhu et al. / Composites Science and Technology 59 (1999) 833±851 841
S. Zhu et al. /Composites Science and Technology 59(1999)833-851 Fatigue of Hi-Nicalon/SIC, 1300 C, 120 MPa, Air Hi-Nicalon, Air 1201E10yy8x0121014 100 680 0000 0 101102103104105106107108 00020.004 Cycles to Failure Fig9. Maximum stress versus cycles to failure of Hi-Nicalon M/Sic Fig. 10. Evolution of the hysteresis loops during fatigue of Hi- Nica- n air, enhanced Sic/Sic in air, and standard Sic/SiC in air and argon lonM SiC at 1300 C under a maximum stress of MPa in air at1300°C. to the right along the strain axis, which is known as 180MP 120 MPa 90 MPa ratchetting due to time-dependent deformation(creep) H105 MPa The modulus normalized with respect to the value from the linear part during the first loading versus cycles is shown in Fig. 1l(a). At stresses higher than 120 MPa the modulus decreases rapidly within 10 cycles, and then goes to gradual decrease stage, and finally drops fast up to fracture. At stresses lower than 105 MPa. the mod ulus first keeps constant up to 10 cycles and then monotonously decreases. At 75 MPa, the modulus keeps constant up to 10cycles, at which the test was stopped. When the modulus decreases to 20-40% of the 0.2 original value, the specimens fracture 10°101102103104105106107 4. Creep and stress rupture Cycle 4.1. Creep behavio Fig. Il. Elastic modulus normalized by the value of the module Creep and stress rupture tests in SiC/Sic were con- ducted recently [73, 79-85]. A constant tensile load pro- The minimum creep strain rate as a function of stress duces an instantaneous strain response followed by a is shown in Fig. 13. The minimum creep strain rate, E time dependent strain(Fig. 12). The instantaneous can be described by power law strain consists of recoverable (elastic) strain at low stresses and nonrecoverable strain at high stresses E=Ao" exp(-Q/RT) which can be determined in Fig. 12. The time dependent (creep)strain is transient, and a continuously decreasing where A is a constant, n is the apparent stress exponent strain rate(primary stage)appears at first. Then it goes for creep, o is the apparent activation energy for creep to a steady state (constant strain rate, secondary) stage, R is the gas constant and T is the absolute temperature at last accelerating(tertiary stage) to rupture. The exis- Fig. 13 reveals that the apparent stress exponent for ence of one, two or three stages depends on the stress creep increases with decreasing stress. The minimum and temperature conditions. At high stresses, there is no apparent stress exponent is 5 and the maximum is 25 tertiary stage or even no secondary stage. The accel- Moreover, an apparent threshold stress exists at a given erating creep stage appears after the steady state creep temperature, below which the creep strain rate falls at low stresses. Abbe et al. [84, 85] also found steady below the detectable level. The apparent threshold stress state creep in flexure tests of SiC/Sic in vacuum at is 75 MPa at 1000 and 1100oC, 60 MPa at 1200C, and temperatures of 1100 to 1400C. 0 MPa at1300°C
to the right along the strain axis, which is known as ratchetting due to time-dependent deformation (creep). The modulus normalized with respect to the value from the linear part during the ®rst loading versus cycles is shown in Fig. 11(a). At stresses higher than 120 MPa, the modulus decreases rapidly within 10 cycles, and then goes to gradual decrease stage, and ®nally drops fast up to fracture. At stresses lower than 105 MPa, the modulus ®rst keeps constant up to 104 cycles and then monotonously decreases. At 75 MPa, the modulus keeps constant up to 107 cycles, at which the test was stopped. When the modulus decreases to 20±40% of the original value, the specimens fracture. 4. Creep and stress rupture 4.1. Creep behavior Creep and stress rupture tests in SiC/SiC were conducted recently [73,79±85]. A constant tensile load produces an instantaneous strain response followed by a time dependent strain (Fig. 12). The instantaneous strain consists of recoverable (elastic) strain at low stresses and nonrecoverable strain at high stresses, which can be determined in Fig. 12. The time dependent (creep) strain is transient, and a continuously decreasing strain rate (primary stage) appears at ®rst. Then it goes to a steady state (constant strain rate, secondary) stage, at last accelerating (tertiary stage) to rupture. The existence of one, two or three stages depends on the stress and temperature conditions. At high stresses, there is no tertiary stage or even no secondary stage. The accelerating creep stage appears after the steady state creep at low stresses. Abbe et al. [84,85] also found steady state creep in ¯exure tests of SiC/SiC in vacuum at temperatures of 1100 to 1400C. The minimum creep strain rate as a function of stress is shown in Fig. 13. The minimum creep strain rate, "_ can be described by power law "_ A:n : exp ÿQ=RT 3 where A is a constant, n is the apparent stress exponent for creep, Q is the apparent activation energy for creep, R is the gas constant and T is the absolute temperature. Fig. 13 reveals that the apparent stress exponent for creep increases with decreasing stress. The minimum apparent stress exponent is 5 and the maximum is 25. Moreover, an apparent threshold stress exists at a given temperature, below which the creep strain rate falls below the detectable level. The apparent threshold stress is 75 MPa at 1000 and 1100C, 60 MPa at 1200C, and 30 MPa at 1300C. Fig. 9. Maximum stress versus cycles to failure of Hi-NicalonTM/SiC in air, enhanced SiC/SiC in air, and standard SiC/SiC in air and argon at 1300C. Fig. 10. Evolution of the hysteresis loops during fatigue of Hi-NicalonTM/SiC at 1300C under a maximum stress of 120 MPa in air. Fig. 11. Elastic modulus normalized by the value of the modulus under the ®rst loading (E=Eo) versus fatigue cycles of Hi-NicalonTM/ SiC in air at 1300C under dierent maximum stresses. 842 S. Zhu et al. / Composites Science and Technology 59 (1999) 833±851
