Int./. Appl Ceram. Technol., 2/2/75-84(2005) pplied Ceramic Tech ceramic Product Development and Commercialization SiC-Matrix Composites: Nonbrittle Ceramics for Thermo- Structural application oger asain Laboratory for Thermostructural Composites, UMR-5801(CNRS-SNECMA-CEA-UBI), University Bordeaux 1, 33600 Pessac, france C/SiC and SiC/SiC composites are tough ceramics when the fiber-matrix bonding is properly optimized, usually through a thin layer of an interfacial material referred to as the interphase. These composites can be fabricated by a variety of techniques that are briefly described and compared. The design of the interphase, matrix, and coating at the nanometer scale, in order to promote microcrack deflection and to enhance the oxidation resistance is discussed. Selected properties of the composites are presented and discussed. Examples of application in engines, heat shields, braking systems, and high-temperature nuclear factors are shown to illustrate the potential of these materials and the key points that still require research and development Introduction ceramic material, potentially allows their use as struc tural materials for HT application in corrosiv e atmos- The ceramic matrix composites(CMCs)considered pheres and explains the tremendous effort of research here consist of ceramic fibers(mainly carbon-or SiC- and development in this field based fibers, generally arranged in multidirectional pre- Although SiC-matrix composite forms)embedded in a SiC-matrix. They have been first materials for application in severe environments, their imagined to replace the carbon/carbon(C/C)compos- development raises a number of issues that will be dis- ites in long-term application at high temperature(Ht) cussed in terms of processing, material design,main hen the atmosphere is oxidizing. It was further d properties, and actual or potential applications in var- covered in a fortuitous manner, and then confirmed by ious fields. a more detailed analysis has been recently theoretical considerations, that CMCs could display a presented elsewhere nonbrittle behavior if the fiber-matrix(FM) bonding was lowered enough, e.g., through the in situ formation of a suitable interfacial reaction zone or the use of a so- Processing called interphase deposited on the fber before the infil tration of the matrix. 4.The damage-tolerant character CMCs are fabricated according to gas phase routes of these CMCs, which is an outstanding property for a CVI: chemical vapor infiltration), liquid phase routes either from polymers(PIP: polymer impregnatic pyrolysis), or molten elements reacting with the the atmosphere (RMi: reactive melt
SiC-Matrix Composites: Nonbrittle Ceramics for ThermoStructural Application Roger R. Naslain* Laboratory for Thermostructural Composites, UMR-5801 (CNRS-SNECMA-CEA-UB1), University Bordeaux 1, 33600 Pessac, France C/SiC and SiC/SiC composites are tough ceramics when the fiber–matrix bonding is properly optimized, usually through a thin layer of an interfacial material referred to as the interphase. These composites can be fabricated by a variety of techniques that are briefly described and compared. The design of the interphase, matrix, and coating at the nanometer scale, in order to promote microcrack deflection and to enhance the oxidation resistance is discussed. Selected properties of the composites are presented and discussed. Examples of application in engines, heat shields, braking systems, and high-temperature nuclear reactors are shown to illustrate the potential of these materials and the key points that still require research and development. Introduction The ceramic matrix composites (CMCs) considered here consist of ceramic fibers (mainly carbon- or SiCbased fibers, generally arranged in multidirectional preforms) embedded in a SiC-matrix. They have been first imagined to replace the carbon/carbon (C/C) composites in long-term application at high temperature (HT) when the atmosphere is oxidizing.1,2 It was further discovered in a fortuitous manner, and then confirmed by theoretical considerations, that CMCs could display a nonbrittle behavior if the fiber–matrix (FM) bonding was lowered enough, e.g., through the in situ formation of a suitable interfacial reaction zone3 or the use of a socalled interphase deposited on the fiber before the infiltration of the matrix.4,5 The damage-tolerant character of these CMCs, which is an outstanding property for a ceramic material, potentially allows their use as structural materials for HT application in corrosive atmospheres and explains the tremendous effort of research and development in this field. Although SiC-matrix composites are promising materials for application in severe environments, their development raises a number of issues that will be discussed in terms of processing, material design, main properties, and actual or potential applications in various fields. A more detailed analysis has been recently presented elsewhere.6 Processing CMCs are fabricated according to gas phase routes (CVI: chemical vapor infiltration), liquid phase routes either from polymers (PIP: polymer impregnation and pyrolysis), or molten elements reacting with the preforms or the atmosphere (RMI: reactive melt Int. J. Appl. Ceram. Technol., 2 [2] 75–84 (2005) Ceramic Product Development and Commercialization *naslain@lcts.u-bordeaux1.fr
International y ournal of Applied Ceramic TechnologyNaslain Vol.2,No.2,2005 filtration), or finally the so-called ceramic or slurry diffusion barrier. Further, the matrix is rarely pure SiC routes(SI-HP: slurry infiltration and hot processing) but a mixture of Sic and free silicon(free silicon low each displaying advantages and drawbacks. generally ering its refractoriness and creep resistance), however, speaking, the matrix should be homogeneously distrib- the content of the latter can be limited if liquid silicon is uted in the preform with limited residual porosity and replaced by a suitable silicon alloy. On the other hand, the FM-bonding well controlled with no significant fb- RMI is a fast densification technique and the corre- er degradation. Further, the process should be flexible sponding composites are near net shape with low resid with limited handling and yield near net shape com- ual porosity(Vp <5%) posites, in order to lower production cost I-CVI and RMi are the that display, from In the CVI-process, the interphase, the matrix, and our viewpoint, the best potential in terms of cost and the seal-coating(used to seal the open residual porosity volume production. Further, they are complementary, and enhance the oxidation resistance)are successively de- i.e., the residual porosity of CVI-composites, at a suit posited from gaseous precursors. In conventional CVI able state of densification, can be filled via an RMI-step (referred to as I-CVI, I standing for isothermal/isobaric), Conversely, the PIP-process, which is also a low-temper there are no temperature/pressure gradients in the fiber ature technique, is lengthy since several time-consuming low-temperature (typically, PI/P sequences( from 6 to 10) to achieve 900-1100C), low-pressure(<100 kPa) process, yielding an able densification near net shape composites with limited fiber degradar significant residual porosity and implies considerable and materials of high microstructural quality. It is also a handling. It can also be combined with RMI, as previ- highly flexible process, a large number of preforms(whic ously mentioned. Fin HP is both could be different in size and shapes) being treated l800° for SiC) and a high-pressure(≈25MPa) multaneously with limited handling, in large infiltration process, which is only compatible with fibers of hi furnaces. All these features justify that I-CVI has been thermal stability(carbon or stoichiometric SiC fibers rapidly transferred from the laboratory to the plant levels. with a risk of fiber degradation. It has been improved Conversely, in I-CVI, the densification rate is relatively through the use of nanometric SiC particles slurry and slow and the residual porosity is significant(typically, 10- additives(Al2O3, Y2O3)forming a liquid phase at 15%). The densification rate can be actually improved by sintering temperature(see, e.g., the NITE-process) pplying to the preform a temperature gradient TG of the main advantages of SI-HP lies in CVD), a pressure gradient(P-CVI), or both(as in forced that it is a fast densification process, yielding composites or F-CVD), but it is at the expense of fexibility(some with almost no residual porosity and hence a high ther fixturing being necessary for each preform to create the mal conductivity. However, its extension to large multi gradient(s). It can also be improved by performing in- directional fiber preforms seems to be problematic termediate surface machining(to re-open the porosity) From this brief analysis, it appears that none of the but that requires additional handling and raises the fab- existing processes is perfect and that hybrid techniques rication cost. Residual porosity(which is detrimental to combining two approaches, such as PIP/RMI or CVI hermal conductivity and oxidation resistance) is usually RMI, might presently be the most appropriate choice; sealed by depositing on the external surface of the com- each step could still be improved in order to gain in posites a suitable coating at the end of the process reproducibility and cost, at plant level In the RMi (or more simply, MI)process, the fiber form is first consolidated with carbon( deposited on the coated fibers, e. g, by PIp)and then impregnated with liquid silicon(or an Si alloy), silicon reacting exo- thermally with the carbon to form in situ the SiC-based The choice of a suitable reinforcement, for a given matrix. RMI is a hT (1400-1600.C)and liquid matrix, is dictated by several considerations including silicon is a highly reactive medium. Hence, it can be FM compatibility, mechanical or/and thermal proper used only with fibers of high thermal stability(carbon or ties, chemical compatibility with the high service tem- oxygen-free SiC-based fibers)protected with a suitable perature, density, and cost. Covalent nonoxide fibers interphase, e. g, dual pyrocarbon/SiC or boron nitride(carbon and oxygen-free SiC fibers) display the best HT (BN)/SiC interphases where the SiC-sublayer acts as a mechanical properties and can be good heat conductors
infiltration), or finally the so-called ceramic or slurry routes (SI–HP: slurry infiltration and hot processing), each displaying advantages and drawbacks. Generally speaking, the matrix should be homogeneously distributed in the preform with limited residual porosity and the FM-bonding well controlled with no significant fiber degradation. Further, the process should be flexible with limited handling and yield near net shape composites, in order to lower production cost. In the CVI-process, the interphase, the matrix, and the seal-coating (used to seal the open residual porosity and enhance the oxidation resistance) are successively deposited from gaseous precursors. In conventional CVI (referred to as I-CVI, I standing for isothermal/isobaric), there are no temperature/pressure gradients in the fiber preform.1,2,7 I-CVI is a low-temperature (typically, 900–11001C), low-pressure (o100 kPa) process, yielding near net shape composites with limited fiber degradation and materials of high microstructural quality. It is also a highly flexible process, a large number of preforms (which could be different in size and shapes) being treated simultaneously with limited handling, in large infiltration furnaces. All these features justify that I-CVI has been rapidly transferred from the laboratory to the plant levels. Conversely, in I-CVI, the densification rate is relatively slow and the residual porosity is significant (typically, 10– 15%). The densification rate can be actually improved by applying to the preform a temperature gradient (TGCVI), a pressure gradient (P-CVI), or both (as in forced or F-CVI), but it is at the expense of flexibility (some fixturing being necessary for each preform to create the gradient(s).8 It can also be improved by performing intermediate surface machining (to re-open the porosity) but that requires additional handling and raises the fabrication cost. Residual porosity (which is detrimental to thermal conductivity and oxidation resistance) is usually sealed by depositing on the external surface of the composites a suitable coating at the end of the process. In the RMI (or more simply, MI) process, the fiber preform is first consolidated with carbon (deposited on the coated fibers, e.g., by PIP) and then impregnated with liquid silicon (or an Si alloy), silicon reacting exothermally with the carbon to form in situ the SiC-based matrix. RMI is a HT process (1400–16001C) and liquid silicon is a highly reactive medium. Hence, it can be used only with fibers of high thermal stability (carbon or oxygen-free SiC-based fibers) protected with a suitable interphase, e.g., dual pyrocarbon/SiC or boron nitride (BN)/SiC interphases where the SiC-sublayer acts as a diffusion barrier.9 Further, the matrix is rarely pure SiC but a mixture of SiC and free silicon (free silicon lowering its refractoriness and creep resistance), however, the content of the latter can be limited if liquid silicon is replaced by a suitable silicon alloy. On the other hand, RMI is a fast densification technique and the corresponding composites are near net shape with low residual porosity (Vpo5%). I-CVI and RMI are the processes that display, from our viewpoint, the best potential in terms of cost and volume production. Further, they are complementary, i.e., the residual porosity of CVI-composites, at a suitable state of densification, can be filled via an RMI-step. Conversely, the PIP-process, which is also a low-temperature technique, is lengthy since several time-consuming PI/P sequences (from 6 to 10) are necessary to achieve an acceptable densification. It yields composites with a significant residual porosity and implies considerable handling. It can also be combined with RMI, as previously mentioned. Finally, SI–HP is both a HT (1700– 18001C for SiC) and a high-pressure ( 25 MPa) process, which is only compatible with fibers of high thermal stability (carbon or stoichiometric SiC fibers) with a risk of fiber degradation.10 It has been improved through the use of nanometric SiC particles slurry and additives (Al2O3, Y2O3) forming a liquid phase at sintering temperature (see, e.g., the NITE-process).11 One of the main advantages of SI–HP lies in the fact that it is a fast densification process, yielding composites with almost no residual porosity and hence a high thermal conductivity. However, its extension to large multidirectional fiber preforms seems to be problematic. From this brief analysis, it appears that none of the existing processes is perfect and that hybrid techniques combining two approaches, such as PIP/RMI or CVI/ RMI, might presently be the most appropriate choice; each step could still be improved in order to gain in reproducibility and cost, at plant level. Material Design The choice of a suitable reinforcement, for a given matrix, is dictated by several considerations including FM compatibility, mechanical or/and thermal properties, chemical compatibility with the high service temperature, density, and cost. Covalent nonoxide fibers (carbon and oxygen-free SiC fibers) display the best HT mechanical properties and can be good heat conductors 76 International Journal of Applied Ceramic Technology—Naslain Vol. 2, No. 2, 2005
wwceramics. org/ACT (depending on their microstructure). Further, they are plex and contradictory functions. First, it should arrest ght and some of them are available in large quantity at and deflect the ma atrix microcr a relatively low cost(carbon fibers). Obviously, they are the interphase debonding energy, I, is low relative to the the reinforcement of choice for nonoxide matrices, e.g., failure energy of the fiber, Ib a generally accepted crite- since amic equilibrium with rion being I /T <1/4. This is the so-called mechanical carbon at high temperature. Also, there exist well-iden- fuse function(the interphase protecting the fiber from an tified nonoxide interphases (pyrocarbon and boron ni- early failure). Second, the interphase may act as a diffu tride)compatible with both carbon and SiC within a sion barrier(as previously mentioned for the rMi proc- wide temperature range. Unfortunately, nonoxide CMCs ess) and relax partly thermal residual stresses. It has been re oxidation prone and their long exposure to oxidizing recently postulated that the best interphase materials res efficient against oxida- might be those with a layered crystal structure or micro- tion(PAO). Oxide-based CMCs are by essence inert in structure, the layers being deposited, parallel to the fiber oxidizing atmospheres and migh ght appear more attrac- surface, weakly bonded to one another(for low T but tive. However, most oxide-based fibers(containing a- strongly adherent to the fiber surface(to avoid debonding alumina, mullite, or zirconia) display poor HT mecha at the fiber surface). In SiC-matrix composites, the best al properties(they suffer from grain growth and creep interphase material from a mechanical standpoint beyond about 1000-1100C). Further, there on(Fig. 1a). 1 Unfor ly no stable oxide interphase formally equivalent to tunately, pyrocarbon is intrisically oxidation-prone at pyrocarbon or BN, i. e, a layered oxide with a low shear temperatures as low as 500%C. BN is an interesting al- at col deposited on fibers ternative since it has a similar layered crystal structure and though there are few oxide interphase materials, such a better oxidation resistance, its oxidation starting at monazite or hibonite, that could deflect matrix cracks in about 800C and yielding a Auid B2O3 oxide known oxide-oxide composites but not as easily as their not for its healing properties. However, its formation on a oxide counterparts). Also, oxide- based CMCs are in- SiC fiber is not straightforward. When deposited at low sulating materials and their density can be slightly temperature by CVD/CVI, it is amorphous or poorly higher than that of C-or SiC-based composites. Final- crystallized and hence sensitive to moisture. Its crystalli ly, fibers should exhibit a good weavability, which sup- zation by heat treatment is often limited by the thermal poses a low enough diameter(typically, 10 um or less) stability of the fibers and the bonding with the fibers is when their stiffness is high(this is the case for most poor. An interesting alternative might be to form a arbon fibers but not for all stoichiometric SiC fibers) more adherent bn coating by annealing a SiC fiber con- and preferably a high failure strain. To conclude, non- taining some boron(used as a sintering aid) in a nitriding oxide CMCs(C/SiC or SiC/SiC)are presently preferred atmosphere at high temperature. SiC/SiC composites for most structural applications even though their use with such an in situ formed BN interphase have been in oxidizing atmospheres raises a difficult problem of reported to be more oxidation resistant than those with durability. BN interphase deposited by CVD/CVl. The choice of a concept of damage tolerance is a k Since the number of thermally stable materials with step in the design of CMCs In SiC-matrix composites, layered structures is limited, the concept of layered inter damage tolerance is achieved through a weakening of the phase has been further extended to materials with a lay FM-bonding(controlled by an interphase), which allows ered microstructure at the nanometer scale, i.e., to (X-Y) the matrix microcracks to be deflected by the FM-inter- multilayers. Such interphases offer a much higher design based on the use of a highly porous matrix (and no in- pulsed CVI (or P-CVT) being the ers by, e.g,pressure- terphase), is known. Its use might be appropriate in ox overall thickness of the interphase, the thicknesses of the X ide/oxide composites since both constituents are inert and Y sublayers, the number of X-Y sequences, n, and Conversely, it might be the X/Bonding. As an example, in(PyC-SiC)n or(BN- lematic in nonoxide CMCs since a porous matrix will SiC)m, the amount of oxidation-prone mechanical favor fiber oxidation and lower thermal conductivity. (-= PyC or BN) can be strongly reduced(the thickness The design of the interphase in SiC-matrix compos- of X-layers being a few nanometers, typically 3-20 nm) ites is not straightforward since with the result that the durability of the composites
(depending on their microstructure). Further, they are light and some of them are available in large quantity at a relatively low cost (carbon fibers). Obviously, they are the reinforcement of choice for nonoxide matrices, e.g., SiC, since SiC is in thermodynamic equilibrium with carbon at high temperature. Also, there exist well-identified nonoxide interphases (pyrocarbon and boron nitride) compatible with both carbon and SiC within a wide temperature range. Unfortunately, nonoxide CMCs are oxidation prone and their long exposure to oxidizing atmospheres requires efficient protection against oxidation (PAO). Oxide-based CMCs are by essence inert in oxidizing atmospheres and might appear more attractive. However, most oxide-based fibers (containing aalumina, mullite, or zirconia) display poor HT mechanical properties (they suffer from grain growth and creep beyond about 1000–11001C). Further, there is presently no stable oxide interphase formally equivalent to pyrocarbon or BN, i.e., a layered oxide with a low shear strength that could be easily deposited on fibers (although there are few oxide interphase materials, such as monazite or hibonite, that could deflect matrix cracks in oxide–oxide composites but not as easily as their nonoxide counterparts).12 Also, oxide-based CMCs are insulating materials and their density can be slightly higher than that of C- or SiC-based composites. Finally, fibers should exhibit a good weavability, which supposes a low enough diameter (typically, 10 mm or less) when their stiffness is high (this is the case for most carbon fibers but not for all stoichiometric SiC fibers) and preferably a high failure strain. To conclude, nonoxide CMCs (C/SiC or SiC/SiC) are presently preferred for most structural applications even though their use in oxidizing atmospheres raises a difficult problem of durability. The choice of a concept of damage tolerance is a key step in the design of CMCs. In SiC-matrix composites, damage tolerance is achieved through a weakening of the FM-bonding (controlled by an interphase), which allows the matrix microcracks to be deflected by the FM-interfaces. However, another concept of damage tolerance, based on the use of a highly porous matrix (and no interphase), is known. Its use might be appropriate in oxide/oxide composites since both constituents are inert in oxidizing atmospheres.13,14 Conversely, it might be problematic in nonoxide CMCs since a porous matrix will favor fiber oxidation and lower thermal conductivity. The design of the interphase in SiC-matrix composites is not straightforward since the interphase has complex and contradictory functions.4 First, it should arrest and deflect the matrix microcracks, which supposes that the interphase debonding energy, Gi , is low relative to the failure energy of the fiber, Gf, a generally accepted criterion being Gi /Gf o1/4.5 This is the so-called mechanical fuse function (the interphase protecting the fiber from an early failure). Second, the interphase may act as a diffusion barrier (as previously mentioned for the RMI process) and relax partly thermal residual stresses. It has been recently postulated that the best interphase materials might be those with a layered crystal structure or microstructure, the layers being deposited, parallel to the fiber surface, weakly bonded to one another (for low Gi ) but strongly adherent to the fiber surface (to avoid debonding at the fiber surface).4 In SiC-matrix composites, the best interphase material from a mechanical standpoint is probably an anisotropic pyrocarbon (Fig. 1a).4,15 Unfortunately, pyrocarbon is intrisically oxidation-prone at temperatures as low as 5001C. BN is an interesting alternative since it has a similar layered crystal structure and a better oxidation resistance, its oxidation starting at about 8001C and yielding a fluid B2O3 oxide known for its healing properties. However, its formation on a SiC fiber is not straightforward. When deposited at low temperature by CVD/CVI, it is amorphous or poorly crystallized and hence sensitive to moisture. Its crystallization by heat treatment is often limited by the thermal stability of the fibers and the bonding with the fibers is poor. 16 An interesting alternative might be to form a more adherent BN coating by annealing a SiC fiber containing some boron (used as a sintering aid) in a nitriding atmosphere at high temperature. SiC/SiC composites with such an in situ formed BN interphase have been reported to be more oxidation resistant than those with a BN interphase deposited by CVD/CVI.17 Since the number of thermally stable materials with layered structures is limited, the concept of layered interphase has been further extended to materials with a layered microstructure at the nanometer scale, i.e., to (X–Y)n multilayers. 4 Such interphases offer a much higher design flexibility, the adjustable parameters by, e.g., pressurepulsed CVI (or P-CVI) being the nature of X and Y, the overall thickness of the interphase, the thicknesses of the X and Y sublayers, the number of X–Y sequences, n, and the X/Y bonding. As an example, in (PyC–SiC)n or (BN– SiC)n, the amount of oxidation-prone mechanical fuse (X 5 PyC or BN) can be strongly reduced (the thickness of X-layers being a few nanometers, typically 3–20 nm) with the result that the durability of the composites in www.ceramics.org/ACT SiC-Matrix Composites: Application 77
International y ournal of Applied Ceramic TechnologyNaslain Vol.2,No.2,2005 Fiber Matrix 500 Fig 1. Interphases for SiC/SiC composites with layered crystal structure or microstructure: (a)anisotropic pyrocarbon single-layer ase and (b)(pyC-SiCIo multilayered interphase. oxidizing atmospheres is improved by self-healing phe- of damaging phenomena, mainly including multiple nomena(silica or SiO2-B2O3 scales formed by oxidation matrix microcracking and FM-debonding. As a result, healing the narrow annular pore created around each fiber their stiffness progressively decreases as the applied load by oxidation)(Fig. 1b). Another interphase concept is raised beyond the proportional limit(SiC/SiC com- that has been less explored is the use of a porous SiC layer, posites), with little permanent deformation upon un- a porous solid displaying a lower failure energy than its loading (at least for well-processed materials). Hence, dense counterpart. However, such a porous interface they are often referred to as damageable elastic materials would favor the oxidation of the fibers as mentioned pre- TI he extent of the nonlinear domain in which the ma- ously for porous matrices terials are damage-tolerant is related to the ultimate fail- Finally, a seal-coating is usually deposited on the ex- ure strain of the fibers (e. g, the latter becoming low, ternal surface of C/SiC and SiC/SiC composites, mainly typically 0.6-0.7% for stoichiometric SiC fibers). Fur- to seal the residual open porosity(composites fabricated ther, the damage features are strongly related to the in- by the Pip or CVI processes)or/and to improve their tensity of the FM-bonding, a point that is often resistance to corrosive environments. Dense single layer underestimated. When the FM-bonding is too weak ceramic coatings(such as SiC or Si3N4) displaying a the matrix microcrack density is low, the microcracks tendency to microcracking(as a result of CTE-mismatch are widely open under load, and debonding occurs over or mechanical loading) multilayered coatings are prefer a long distance(and sometimes over the whole fiber able, as it will be discussed in the next section. Such ngth, exposing the oxidation-prone fibers to the am- coatings are deposited by PVd or P-CVD bient environment). By contrast, when the FM-bonding is stronger and the interphase is strongly adherent to the Selected Properties fiber. it is the reverse situation that is observed. the composite displaying a higher failure stress(Fig. 2)and Mechanical Bebavior a better oxidation resistance. SiC-matrix ce tough when properly designed and fabricated SiC-matrix composites display a nonlinear stress- toughness, expressed in terms of critical energy release ain behavior when tensile loaded in one of the fiber rate of the order of 10 kJ/m2, whereas that of monolithic ections. This nonlinearity is related to the occurrence SiC-ceramics is of the order of a few 100
oxidizing atmospheres is improved by self-healing phenomena (silica or SiO2–B2O3 scales formed by oxidation healing the narrow annular pore created around each fiber by oxidation) (Fig. 1b).4,18 Another interphase concept that has been less explored is the use of a porous SiC layer, a porous solid displaying a lower failure energy than its dense counterpart. However, such a porous interface would favor the oxidation of the fibers as mentioned previously for porous matrices. Finally, a seal-coating is usually deposited on the external surface of C/SiC and SiC/SiC composites, mainly to seal the residual open porosity (composites fabricated by the PIP or CVI processes) or/and to improve their resistance to corrosive environments. Dense single layer ceramic coatings (such as SiC or Si3N4) displaying a tendency to microcracking (as a result of CTE-mismatch or mechanical loading) multilayered coatings are preferable, as it will be discussed in the next section.19 Such coatings are deposited by PVD or P-CVD. Selected Properties Mechanical Behavior SiC-matrix composites display a nonlinear stress– strain behavior when tensile loaded in one of the fiber directions. This nonlinearity is related to the occurrence of damaging phenomena, mainly including multiple matrix microcracking and FM-debonding. As a result, their stiffness progressively decreases as the applied load is raised beyond the proportional limit (SiC/SiC composites), with little permanent deformation upon unloading (at least for well-processed materials). Hence, they are often referred to as damageable elastic materials. The extent of the nonlinear domain in which the materials are damage-tolerant is related to the ultimate failure strain of the fibers (e.g., the latter becoming low, typically 0.6–0.7% for stoichiometric SiC fibers). Further, the damage features are strongly related to the intensity of the FM-bonding, a point that is often underestimated.15 When the FM-bonding is too weak, the matrix microcrack density is low, the microcracks are widely open under load, and debonding occurs over a long distance (and sometimes over the whole fiber length, exposing the oxidation-prone fibers to the ambient environment). By contrast, when the FM-bonding is stronger and the interphase is strongly adherent to the fiber, it is the reverse situation that is observed, the composite displaying a higher failure stress (Fig. 2) and a better oxidation resistance. SiC-matrix composites are tough when properly designed and fabricated with toughness, expressed in terms of critical energy release rate of the order of 10 kJ/m2 , whereas that of monolithic SiC-ceramics is of the order of a few 100 J/m2 . 15 Fig. 1. Interphases for SiC/SiC composites with layered crystal structure or microstructure: (a) anisotropic pyrocarbon single-layer interphase15 and (b) (PyC–SiC)10 multilayered interphase.18 78 International Journal of Applied Ceramic Technology—Naslain Vol. 2, No. 2, 2005
wwceramics. org/ACT matri I matrix pyc 02 LONGITUDINAL TENSILE STRAIN (o) Fig. 2. Typical stress-strain tensile curves of 2D-SiC(Nicalon )/PyC/SiC composites with weak FM-bonding(material D)and stronger aterial刀 Finally, they are less fatigue-prone than metals and al- fibers fabricated from mesophase pitch and heat treated loys with stress threshold below which no fatigue failure beyond 2500C(P55-P130 series)but low(10 W/m. K occurs, of the order of 75% the ultimate failure stress and less) for poorly organized fibers(ex-PAN T300 fib- ers).In a similar manner, nearly stoichiometric SiC The tensile stress-strain behavior of SiC-matrix fibers fabricated at high temperatures(e. g, Tyranno S. composites does not change markedly up to N 1100C. fibers, Ube Industrial, Japan) display a much better However, some change may be observed either at higher conductivity than the quasi-amorphous Si-C-O fibers temperatures if the fibers are limited in thermal stability prepared at low temperatures, typically 65 and 10 W/ (case of the unstable Si-C-O fibers) or even at lower m K at room temperature, respectively. Equally im- temperatures when an oxidizing atmosphere has access portant is the effect of the residual porosity, a composite to the fibers and the interphase(case of insufficiently produced by RMI or hot pressing(Vp s5%)exhibiting protected materials). Further, SiC-matrix composites a higher conductivity than a composite fabricated by creep at high temperatures with a creep rate depending PIP or CVI (VP N 10-15%). Hence, a SiC/SiC com on the nature of the fibers(stoichiometric microcrystal- posite is expected to show a thermal conductivity of the line SiC fibers prepared or treated at high temperatures order of 30 W/m K at 1000.C when prepared from being more creep-resistant than their Si-C-O nano- nearly stoichiometric SiC fibers with almost no residual rystalline counterparts)and that of the matrix and possibly higher if the reinforcement con sists of graphitized carbon fibers(with, however, in this Thermal Conductivity case a risk related to the occurrence of microcracking due to CTE-mismatch that will lower the conductivity) Thermal conductivity is a key property in many HT applications of CMCs. Generally speaking, SiC Oxidation resistance matrix composites are relatively good conductors of heat but their thermal conductivity depends on the crystal- In most thermostructural applications, SiC-matrix linity of their constituents, the FM-bonding, and resid- composites are exposed to oxidizing atmospheres. Since al porosity. The thermal conductivity of carbon fibers their constituents are intrinsically oxidation-prone, their can be very high (100 W/m K and more) for the behavior under such environments is of key importand
Finally, they are less fatigue-prone than metals and alloys with stress threshold below which no fatigue failure occurs, of the order of 75% the ultimate failure stress under static loading.20 The tensile stress–strain behavior of SiC-matrix composites does not change markedly up to 11001C. However, some change may be observed either at higher temperatures if the fibers are limited in thermal stability (case of the unstable Si–C–O fibers) or even at lower temperatures when an oxidizing atmosphere has access to the fibers and the interphase (case of insufficiently protected materials). Further, SiC-matrix composites creep at high temperatures with a creep rate depending on the nature of the fibers (stoichiometric microcrystalline SiC fibers prepared or treated at high temperatures being more creep-resistant than their Si–C–O nanocrystalline counterparts) and that of the matrix.21 Thermal Conductivity Thermal conductivity is a key property in many HT applications of CMCs. Generally speaking, SiCmatrix composites are relatively good conductors of heat but their thermal conductivity depends on the crystallinity of their constituents, the FM-bonding, and residual porosity. The thermal conductivity of carbon fibers can be very high (100 W/m K and more) for those fibers fabricated from mesophase pitch and heat treated beyond 25001C (P55–P130 series) but low (10 W/m K and less) for poorly organized fibers (ex-PAN T300 fibers).22 In a similar manner, nearly stoichiometric SiC fibers fabricated at high temperatures Q2 (e.g., Tyranno SA fibers, Ube Industrial, Japan) display a much better conductivity than the quasi-amorphous Si–C–O fibers prepared at low temperatures, typically 65 and 10 W/ m K at room temperature, respectively.23 Equally important is the effect of the residual porosity, a composite produced by RMI or hot pressing (Vpr5%) exhibiting a higher conductivity than a composite fabricated by PIP or CVI (Vp 10–15%). Hence, a SiC/SiC composite is expected to show a thermal conductivity of the order of 30 W/m K at 10001C when prepared from nearly stoichiometric SiC fibers with almost no residual porosity, and possibly higher if the reinforcement consists of graphitized carbon fibers (with, however, in this case a risk related to the occurrence of microcracking due to CTE-mismatch that will lower the conductivity). Oxidation Resistance In most thermostructural applications, SiC-matrix composites are exposed to oxidizing atmospheres. Since their constituents are intrinsically oxidation-prone, their behavior under such environments is of key importance Fig. 2. Typical stress–strain tensile curves of 2D-SiC(Nicalon)/PyC/SiC composites with weak FM-bonding (material I) and stronger FM-bonding (material J), corresponding to different matrix crack deflection schemes (according to Droillard15). www.ceramics.org/ACT SiC-Matrix Composites: Application 79