《复合材料 Composites》课程教学资源（学习资料）第五章陶瓷基复合材料_Fracture mechanisms for SiC fibers and SiC/SiC composites under stress-rupture conditions at high temperatures.pdf

J.A. DiCarlo et al. Appl Math. Comput. 152(2004)473-481 fiber-reinforced SiC matrix composites [1, 2]. Since these SiC/SiC composites are still in their infancy in terms of selecting and demonstrating the optimum fiber, interphase, and matrix constituents, there currently exists a strong need for studies that mechanistically analyze and predict the fracture- limited enve lope of thermostructual capability provided by currently available constituents Thus the objective of this paper is to present mechanistic models concerning the high-temperature stress-rupture behavior of SiC fibers and Sic/Sic com posites of current technical interest. Since the time-dependent fracture of these materials is controlled by creep-induced flaw growth, these models are based on key experimental observations made on the creep-rupture behavior for various SiC fiber types. These observations are important not only because the fibers are the primary structural constituents controlling ultimate composite rupture, but also because SiC fibers display microstructures and fracture be- havior representative of SiC matrices, which can also carry structural loads Sic/SiC composites 2. SiC fiber creep-rupture In terms of providing good thermomechanical reinforcement capability for SiC/SiC composites, small-diameter SiC fiber types with low oxygen content, such as the non-stoichiometric (C/Si>)Hi-Nicalon fiber from Nippon Car- bon and the stoichiometric(C/SiN I)Sylramic fiber from Dow Corning, have many desirable physical and chemical properties [3, 4]. For this reason, con- siderable creep-rupture property data exist for these two fiber types not only as single fibers [3-8], multifilament tows[9, 10], and woven fabric pieces [ll],but so as reinforcement for Sic/Sic composites with various types of Sic ma trices [12-15]. For example, Fig. I shows typical creep strain versus time curves ④F140c cREEPsTRAIN TIME. HRS Fig. 1. Typical creep curves for Sic fibers tested under high-temperature stress-rupture conditions

fiber-reinforced SiC matrix composites [1,2]. Since these SiC/SiC composites are still in their infancy in terms of selecting and demonstrating the optimum fiber, interphase, and matrix constituents, there currently exists a strong need for studies that mechanistically analyze and predict the fracture-limited envelope of thermostructual capability provided by currently available constituents. Thus the objective of this paper is to present mechanistic models concerning the high-temperature stress-rupture behavior of SiC fibers and SiC/SiC composites of current technical interest. Since the time-dependent fracture of these materials is controlled by creep-induced flaw growth, these models are based on key experimental observations made on the creep-rupture behavior for various SiC fiber types. These observations are important not only because the fibers are the primary structural constituents controlling ultimate composite rupture, but also because SiC fibers display microstructures and fracture behavior representative of SiC matrices, which can also carry structural loads in SiC/SiC composites. 2. SiC fiber creep-rupture In terms of providing good thermomechanical reinforcement capability for SiC/SiC composites, small-diameter SiC fiber types with low oxygen content, such as the non-stoichiometric (C/Si >1) Hi-Nicalon fiber from Nippon Carbon and the stoichiometric (C/Si 1) Sylramic fiber from Dow Corning, have many desirable physical and chemical properties [3,4]. For this reason, considerable creep-rupture property data exist for these two fiber types not only as single fibers [3–8], multifilament tows [9,10], and woven fabric pieces [11], but also as reinforcement for SiC/SiC composites with various types of SiC matrices [12–15]. For example, Fig. 1 shows typical creep strain versus time curves 0 10 20 30 40 50 0.0 0.5 1.0 1.5 2.0 2.5 TIME, HRS 1400C 270 MPa AIR : Sylramic : Hi-Nicalon. : Ultra SCS : Nicalon : SCS-6 C R E E P S T R A I N, % Fr R. Fr A D D B B E C A C E Fig. 1. Typical creep curves for SiC fibers tested under high-temperature stress-rupture conditions. 474 J.A. DiCarlo et al. / Appl. Math. Comput. 152 (2004) 473–481

J.A. DiCarlo et al. /Appl. Math. Comput. 152(2004)473-481 for single Hi-Nicalon and Sylramic fibers as measured at 1400C in air under a constant stress of 270 MPa Rupture points are indicated by the symbol Fr Also shown are typical creep curves measured under the same conditions for the oxygen-containing Nicalon SiC fiber and for the oxygen-free SCS-6 and Ultra-SCS SiC fibers from Textron Specialty Materials [5, 16]. As will be dis- cussed, although the large diameters of the SCS fibers limit their formability for complex-shaped Sic/Sic components, their microstructures are similar those of SiC matrices formed by chemical vapor infiltration, a common matrix formation process for SiC/SiC composites One important observation from the Fig. I data is that all Sic fibers display a primary or transient creep stage that varies in magnitude depending on fiber type. This effect is typically observed in Si-based ceramics in which there exists an initial"" or recoverable creep component controlled by grain- boundary sliding [17]. Along with the anelastic component, there also exists a viscoelastic""or steady-state creep component, which is typically accompanied by creep-induced micro-crack formation, resulting in critical flaw growth and non-recoverable strain. Both components are enhanced by small grain size and by the high diffusion of impurity second phases in the grain boundaries. Thus the high creep behavior of the Nicalon fiber can be explained both by its ox- ygen content and its very fine grain size(5 nm); while the high creep of the H Nicalon fiber is primarily due to its fine grain size of 10 nm. The rather large transient stage for these fibers may also be caused in part by concurrent grain growth during the high-temperature stress-rupture test [7]. Impurity effects also appear to be responsible for the creep behavior of the coarse-grained (100 nm) Sylramic and SCS-6 fibers, which contain boron sintering phases and free silicon, respectively. This conclusion is supported by the observation that the Sylramic fiber transient stage decreases significantly by post-treatments that remove boron from the fiber bulk [8], and by the Fig. I creep data for the UItra-SCS fiber, which has the same grain size as the SCs-6 fiber, but no boron or excess silicon. Thus microstructure plays a strong role in the creep behavior of Sic fibers in particular and Sic-based ceramics in general Another point not shown in Fig. I is that creep rates and total creep strains for a given time period are typically increased for some fiber types by testing in an argon environment, with the effect being the largest for the Sylramic fiber [6, 8]. Although the exact mechanism responsible for this behavior is still un clear, the source appears to be related to decomposition effects. That is, testing in an argon or inert environments should enhance the gaseous volatilization of atomic species(such as silicon or its oxides) from the fiber surface and from within the fiber bulk, either along grain boundaries or through open porosity On the other hand, testing in air can form a silica coating on the fiber surface, which will effectively seal off gaseous volatilization from the bulk. As for possible volatizing species for the Sylramic fiber, this type does contain second phase impurities, such as boron and TiB. However, their removal should only

for single Hi-Nicalon and Sylramic fibers as measured at 1400 C in air under a constant stress of 270 MPa. Rupture points are indicated by the symbol Fr. Also shown are typical creep curves measured under the same conditions for the oxygen-containing Nicalon SiC fiber and for the oxygen-free SCS-6 and Ultra-SCS SiC fibers from Textron Specialty Materials [5,16]. As will be discussed, although the large diameters of the SCS fibers limit their formability for complex-shaped SiC/SiC components, their microstructures are similar to those of SiC matrices formed by chemical vapor infiltration, a common matrix formation process for SiC/SiC composites. One important observation from the Fig. 1 data is that all SiC fibers display a primary or transient creep stage that varies in magnitude depending on fiber type. This effect is typically observed in Si-based ceramics in which there exists an initial ‘‘anelastic’’ or recoverable creep component controlled by grainboundary sliding [17]. Along with the anelastic component, there also exists a ‘‘viscoelastic’’ or steady-state creep component, which is typically accompanied by creep-induced micro-crack formation, resulting in critical flaw growth and non-recoverable strain. Both components are enhanced by small grain size and by the high diffusion of impurity second phases in the grain boundaries. Thus the high creep behavior of the Nicalon fiber can be explained both by its oxygen content and its very fine grain size (5 nm); while the high creep of the HiNicalon fiber is primarily due to its fine grain size of 10 nm. The rather large transient stage for these fibers may also be caused in part by concurrent grain growth during the high-temperature stress-rupture test [7]. Impurity effects also appear to be responsible for the creep behavior of the coarse-grained (100 nm) Sylramic and SCS-6 fibers, which contain boron sintering phases and free silicon, respectively. This conclusion is supported by the observation that the Sylramic fiber transient stage decreases significantly by post-treatments that remove boron from the fiber bulk [8], and by the Fig. 1 creep data for the Ultra-SCS fiber, which has the same grain size as the SCS-6 fiber, but no boron or excess silicon. Thus microstructure plays a strong role in the creep behavior of SiC fibers in particular and SiC-based ceramics in general. Another point not shown in Fig. 1 is that creep rates and total creep strains for a given time period are typically increased for some fiber types by testing in an argon environment, with the effect being the largest for the Sylramic fiber [6,8]. Although the exact mechanism responsible for this behavior is still unclear, the source appears to be related to decomposition effects. That is, testing in an argon or inert environments should enhance the gaseous volatilization of atomic species (such as silicon or its oxides) from the fiber surface and from within the fiber bulk, either along grain boundaries or through open porosity. On the other hand, testing in air can form a silica coating on the fiber surface, which will effectively seal off gaseous volatilization from the bulk. As for possible volatizing species for the Sylramic fiber, this type does contain secondphase impurities, such as boron and TiB2. However, their removal should only J.A. DiCarlo et al. / Appl. Math. Comput. 152 (2004) 473–481 475

476.A.DiCarloetal./Appl.Mat.Comput.152(2004)473-481 reduce rather than enhance creep. Therefore, at the present time, it would appear that the most likely species for volatilization is silicon from the Sic grains, which in turn can reduce grain size and enhance fiber creep rate. Indeed this effect should be especially observable in the Sylramic fiber, which has the smallest diameter(10 um)and therefore the largest surface to volume ratio of all the Fig. I fibers. Thus, from a technical point-of-view, one might expect differences in fiber creep behavior within a SiC/SiC composite, depending on whether the fiber surfaces are exposed to inert or to oxidizing environments However, for most high-temperature applications of current interest, the typ- ical service environments for the composites are oxidizing; so that even with porous interfacial zones, silica protective layers should eventually form on the Sic/SiC outer surfaces and effectively inhibit further volatilization of gaseous species from the fiber. It follows then that for mechanistic modeling of Sic/SiC creep-rupture behavior, the fiber in situ environment should be considered; but for most applications, the single fiber data in air should be appropriate A final point of interest for Fig. I is the rupture life behavior of the various SiC fiber types. In this regard, some important general observations have been made in previous studies [8, 18]. These include: (1) the activation energy con- trolling high-temperature rupture of the SiC fibers is nearly constant and equal to the activation energy controlling fiber creep;(2)fiber rupture strains for a given test temperature and test environment are fairly independent of the minimum creep rate at rupture, but are dependent on fiber type; (3)when the minimum creep rate is enhanced by argon testing, rupture times are shorter rupture strains are typically 100% higher in air than in argon; and(5)some of the fiber types rupture before clear attainment of steady-state creep behavior (see Fig. 1). The activation energy observation confirms that the most im- portant mechanism controlling the high-temperature fracture of SiC fibers(and monolithic ceramics)is creep-induced micro-crack growth. The observa of rupture strain independence on creep rate and shorter rupture times higher creep rates in argon indicate that for each fiber type, a certain amount of micro-crack growth is required to create the critical flaw size. Thus lower creep rates are conducive to longer fiber rupture lives. The observation of an approximate doubling of fiber rupture strain in air suggests that the con trolling micro-crack growth is primarily on the fiber surface, where silica for mation can blunt the stress concentration arising from this micro-crack growth. Finally, the observation of rupture before obvious steady-state creep behavior indicates that viscoelastic creep and micro-crack growth are on going The high-temperature creep curves of monolithic Si-based ceramics, such as pressureless sintered Si3N3[19], are very similar to those of the Sic fibers. For the monolithic materials, a convenient empirical approach for modeling rup- ture time is by use of Monkman-Grant(MG) diagrams that plot the log of average rupture time versus the log of minimum creep rate at various tem-

reduce rather than enhance creep. Therefore, at the present time, it would appear that the most likely species for volatilization is silicon from the SiC grains, which in turn can reduce grain size and enhance fiber creep rate. Indeed, this effect should be especially observable in the Sylramic fiber, which has the smallest diameter (10 lm) and therefore the largest surface to volume ratio of all the Fig. 1 fibers. Thus, from a technical point-of-view, one might expect differences in fiber creep behavior within a SiC/SiC composite, depending on whether the fiber surfaces are exposed to inert or to oxidizing environments. However, for most high-temperature applications of current interest, the typical service environments for the composites are oxidizing; so that even with porous interfacial zones, silica protective layers should eventually form on the SiC/SiC outer surfaces and effectively inhibit further volatilization of gaseous species from the fiber. It follows then that for mechanistic modeling of SiC/SiC creep-rupture behavior, the fiber in situ environment should be considered; but for most applications, the single fiber data in air should be appropriate. A final point of interest for Fig. 1 is the rupture life behavior of the various SiC fiber types. In this regard, some important general observations have been made in previous studies [8,18]. These include: (1) the activation energy controlling high-temperature rupture of the SiC fibers is nearly constant and equal to the activation energy controlling fiber creep; (2) fiber rupture strains for a given test temperature and test environment are fairly independent of the minimum creep rate at rupture, but are dependent on fiber type; (3) when the minimum creep rate is enhanced by argon testing, rupture times are shorter; (4) rupture strains are typically 100% higher in air than in argon; and (5) some of the fiber types rupture before clear attainment of steady-state creep behavior (see Fig. 1). The activation energy observation confirms that the most important mechanism controlling the high-temperature fracture of SiC fibers (and SiC monolithic ceramics) is creep-induced micro-crack growth. The observations of rupture strain independence on creep rate and shorter rupture times under higher creep rates in argon indicate that for each fiber type, a certain amount of micro-crack growth is required to create the critical flaw size. Thus lower creep rates are conducive to longer fiber rupture lives. The observation of an approximate doubling of fiber rupture strain in air suggests that the controlling micro-crack growth is primarily on the fiber surface, where silica formation can blunt the stress concentration arising from this micro-crack growth. Finally, the observation of rupture before obvious steady-state creep behavior indicates that viscoelastic creep and micro-crack growth are on going during the transient stage. The high-temperature creep curves of monolithic Si-based ceramics, such as pressureless sintered Si3N3 [19], are very similar to those of the SiC fibers. For the monolithic materials, a convenient empirical approach for modeling rupture time is by use of Monkman–Grant (MG) diagrams that plot the log of average rupture time versus the log of minimum creep rate at various tem- 476 J.A. DiCarlo et al. / Appl. Math. Comput. 152 (2004) 473–481

peratures [20]. On MG diagrams, the log–log results at various temperatures typically fall on a set of parallel straight lines that allow average rupture time t to be described by the simple equation, tem ¼ C: ð1Þ Here e is the steady state (or minimum) creep rate, and m and C are empirically determined parameters. Using this approach, Fig. 2 shows that the rupture behavior for the five SiC fiber types of Fig. 1 can indeed be best-fit to straight MG lines using data measured at 1200 C in air and at 1400 C in air for the Hi-Nicalon fiber. For these lines, fiber creep rate is equivalent to minimum creep rate or to the instantaneous creep rate at fiber rupture since no tertiary stages were ever observed. Fig. 2 also shows that, except for the Nicalon and Hi-Nicalon fibers, the rupture lines of the different fiber types at 1200 C in air do not fall on one master line as has been observed for various alumina-based fibers [21]. This can probably be attributed in part to the measured creep rates being larger than the viscoelastic creep rates due to the presence of variable anelastic creep components at fiber rupture. However, in relation to Eq. (1), all the lines display approximately the same slope with an m exponent of 1.2. The fact that the m value is greater than unity indicates that the average fiber rupture strain from the viscoelastic component ( e t) increased with increasing rupture time or decreasing stress on the fibers. Also, by increasing test temperature from 1200 to 1400 C, all fiber lines shifted upward (>fivefold increase in rupture time for a given creep rate), like that of the Hi-Nicalon fiber in Fig. 2. However, the HiNicalon line at 1400 C displayed an m value less than unity, implying reduced 1E-10 1E-09 1E-08 1E-07 1E-06 1E-05 0.0001 0.01 0.1 1 10 100 1,000 RUPTURE TIME, hours CREEP RATE, sec -1 Sylramic (1200C) Ultra-SCS (1200C) Hi-Nicalon and Nicalon CVI SiC Matrix (1300C) (1200C) Hi-Nicalon (1400C) :: + + + D D SiC/SiC Composites + Zhu et al., 1997 (1300C) O Zhu et al., 1999 (1300C) D Hurst et al., 2001 (1315C) Fig. 2. Monkman–Grant lines measured in air for SiC fibers, SiC matrices, and SiC/SiC composites. J.A. DiCarlo et al. / Appl. Math. Comput. 152 (2004) 473–481 477

《复合材料 Composites》课程教学资源（学习资料）第五章 陶瓷基复合材料_Fracture mechanisms for SiC fibers and SiC/SiC composites under stress-rupture conditions at high temperatures

《复合材料 Composites》课程教学资源（学习资料）第五章陶瓷基复合材料_Fracture mechanisms for SiC fibers and SiC/SiC composites under stress-rupture conditions at high temperatures