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, the main damage mechanism occurring during creep is periodic ®ber fracture and the creep behavior is governed by embedded ®bers. 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(< 104 modulus can be caused by the increase of either crack cycles), in which the stress exponent for fatigue life is number or crack length. Therefore, any factor(cyclic high and there is no evident difference in fatigue life loading, high temperature creep, oxidation, etc. )can between room temperature and high temperature, lead to a decrease of the modulus (life) of a specimen if although the slope of the curve at the latter temperature it increases crack number or crack length seems higher. The second regime is the rapid decrease in ruments on CMCs are fatigue life with stress at 1000oC for stresses lower than mostly applicable to unidirectional fiber reinforced 180 MPa. There is no second regime at room tempera- composites [70, 71]. For 2D woven SiC/SiC composite, ture. This means that the second regime is dependent on the basic elements are 0 bundles, 90 bundles and high-temperature effects. This will be discussed later. pores. Therefore, their interaction must be taken into The third regime exhibits a fatigue limit defined by the consideration to explain mechanical properties [72] specimens below it having a life over 10 cycles. The The cyclic fatigue fracture surfaces at room tempera fatigue limit at 1000 C is only 75 MPa, which is about ture and at high stresses(fatigue life <10 cycles )at 30% of UTS. The fatigue limit at room temperature is 000C are similar to the monotonic tensile fracture 160 MPa(about 80% of UTS) surfaces, shown in Fig. 6. However, cyclic fracture at Although the UTS and proportional limit at 1000c low stresses( fatigue life >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), in which the stress exponent for fatigue life is high and there is no evident dierence in fatigue life between room temperature and high temperature, although the slope of the curve at the latter temperature seems higher. The second reÂgime is the rapid decrease in fatigue life with stress at 1000C for stresses lower than 180 MPa. There is no second reÂgime at room temperature. This means that the second reÂgime is dependent on high-temperature eects. This will be discussed later. The third reÂgime exhibits a fatigue limit de®ned by the specimens below it having a life over 107 cycles. The fatigue limit at 1000C is only 75 MPa, which is about 30% of UTS. The fatigue limit at room temperature is 160 MPa (about 80% of UTS). Although the UTS and proportional limit at 1000C are higher than those at room temperature (Fig. 1), the fatigue limit at 1000C is much lower than that at room temperature. This means that the fatigue limit is not proportional to the monotonic tensile strength. This point is important to designers, since they cannot design components for high-temperature use on basis only of monotonic strength. The fatigue limit of the composite at room temperature is much higher than the stress for the matrix cracking. This means that the composites can avoid the unsteady propagation of the matrix cracks induced by the ®rst loading during cyclic fatigue at stresses up to the fatigue limit. In other words, the matrix cracks formed on ®rst loading may propagate and also microcracking of the matrix may continue over several thousand cycles [69], but they are ®nally arrested and remain stationary during subsequent cyclic loading. Fiber bridging is commonly thought to be the main reason for this phenomenon, since the bridging force decreases the stress intensity at crack tip. At stresses above the fatigue limit, cyclic fatigue fracture occurs accompanied by a modulus reduction with cycles [56,62]. This reduction of modulus can be caused by the increase of either crack number or crack length. Therefore, any factor (cyclic loading, high temperature creep, oxidation, etc.) can lead to a decrease of the modulus (life) of a specimen if it increases crack number or crack length. The available theory and experiments on CMCs are mostly applicable to unidirectional ®ber reinforced composites [70,71]. For 2D woven SiC/SiC composite, the basic elements are 0 bundles, 90 bundles and pores. Therefore, their interaction must be taken into consideration to explain mechanical properties [72]. The cyclic fatigue fracture surfaces at room temperature and at high stresses (fatigue life <104 cycles) at 1000C are similar to the monotonic tensile fracture surfaces, shown in Fig. 6. However, cyclic fracture at low stresses (fatigue life > 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