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
enational yournal of Applied Ceramic Technology-Naslain Vol.2,No.2,2005 barrier top coat B:B,C: SiB,: Si-B-c functional layer(s) Si or sic CMC-substrate (a) 017K0X1,79810ymWD35 Fig 3. Nonoxide CMCs with improued oxidation resistance through the use of multilayered seal-coating (a)and multilayered selfhealing nara (b). Adapted from C nd vandenbulckeg and lamouroux for durability. When a SiC/SiC (or a C/SiC)composite the behavior of C/SiC and SiC/SiC composites in ox- with a pyrocarbon(or BN) interphase is heated in an idizing atmospheres is usually better at relatively high oxidizing atmosphere, active or passive oxidation phe- temperatures(1200C) than at lower temperatures nomena are observed depending on whether all the re- However, this protection due to condensed oxides is action products are gaseous( CO or CO2 for carbon, insufficient for long exposures under load and even dis- CO and Sio for SiC)or at least one reaction product is appears in wet atmospheres(volatilization of silica) ondensed as a covering protective scale(silica for Sic Under such conditions, some protection against oxida- and boria for BN), respectively. Fortunately, in many tion is necessary cases the oxidation regime is passive. Under this as- There are two ways to improve the oxidation resist sumption, the effect of oxidation on the microstructure, ance of Sic-matrix composites, which are based on mechanical, and thermal properties depends on the multilayered seal-coatings or self-healing matrices. Ho- oxidation conditions(temperature, oxygen partial pres- mogeneous single-layer coatings, such as dense SiC sure)and material parameters(interphase thickness). coating, provide an insufficient oxidation protection At low temperatures, 500 tional layer containing species(such as B, B4C, SiB,,or Conversely, at high temperatures(1000-1200oC)the Si-B-C mixture)that can form Auid oxides(B2O3 or kinetics of formation of silica (and boria when a B2O3-SiO2) when exposed to an oxidizing atmosphere BN-interphase is used) is fast and the condensed oxide in a given range of temperature, and (ii)a barrier top scale(which is covering for both B2O3 on BN and silica coat that can be a dense SiC-layer, the overall thickness on SiC) is protective and tends to seal or/and fill the of the coating being of the order of 150-200 um(Fig residual pores and microcracks, stopping(or at least 3a). When the coating goes microcrack slowing down) the in-depth diffusion of oxygen. Hence, cyclic loading, the microcracks are being filled by the
for durability. When a SiC/SiC (or a C/SiC) composite with a pyrocarbon (or BN) interphase is heated in an oxidizing atmosphere, active or passive oxidation phenomena are observed depending on whether all the reaction products are gaseous (CO or CO2 for carbon, CO and SiO for SiC) or at least one reaction product is condensed as a covering protective scale (silica for SiC and boria for BN), respectively. Fortunately, in many cases the oxidation regime is passive. Under this assumption, the effect of oxidation on the microstructure, mechanical, and thermal properties depends on the oxidation conditions (temperature, oxygen partial pressure) and material parameters (interphase thickness).24 At low temperatures, 500oTo9001C, the kinetics of oxidation of the pyrocarbon interphase in a SiC/PyC/ SiC composite is already fast whereas that of SiC is almost negligible. As a result, oxidation is an in-depth phenomenon that progressively consumes the interphase, destroys the FM-bonding, degrades the mechanical behavior, and alters the thermal conductivity. The effect is still more significant if the composite is reinforced with carbon fibers and heavily microcracked as a result of an applied load or CTE-mismatch.25 Conversely, at high temperatures (1000–12001C) the kinetics of formation of silica (and boria when a BN-interphase is used) is fast and the condensed oxide scale (which is covering for both B2O3 on BN and silica on SiC) is protective and tends to seal or/and fill the residual pores and microcracks, stopping (or at least slowing down) the in-depth diffusion of oxygen. Hence, the behavior of C/SiC and SiC/SiC composites in oxidizing atmospheres is usually better at relatively high temperatures (12001C) than at lower temperatures. However, this protection due to condensed oxides is insufficient for long exposures under load and even disappears in wet atmospheres (volatilization of silica).26 Under such conditions, some protection against oxidation is necessary. There are two ways to improve the oxidation resistance of SiC-matrix composites, which are based on multilayered seal-coatings or self-healing matrices. Homogeneous single-layer coatings, such as dense SiCcoating, provide an insufficient oxidation protection for C/SiC and SiC/SiC composites submitted to thermal shocks or/and mechanical cyclic loading. In both cases, microcracks are formed in the coating that favor the in-depth diffusion of oxygen. A first strategy is to use a multilayered seal-coating that usually consists of the following: (i) a bond coat, such as a dense layer of SiC (for SiC CVI-matrix) or silicon (for SiC1Si RMI-matrix) deposited on the external surface of the composite at the end of the fiber preform densification, (ii) a functional layer containing species (such as B, B4C, SiB6, or Si–B–C mixture) that can form fluid oxides (B2O3 or B2O3–SiO2) when exposed to an oxidizing atmosphere in a given range of temperature, and (iii) a barrier topcoat that can be a dense SiC-layer, the overall thickness of the coating being of the order of 150–200 mm (Fig. 3a). When the coating undergoes microcracking upon cyclic loading, the microcracks are being filled by the Fig. 3. Nonoxide CMCs with improved oxidation resistance through the use of multilayered seal-coating (a) and multilayered self-healing matrix (b). Adapted from Goujard and Vandenbulcke19 and Lamouroux et al.,27 respectively. 80 International Journal of Applied Ceramic Technology—Naslain Vol. 2, No. 2, 2005
wwceramics. org/ACT Auid oxides and the in-depth diffusion of oxygen is established and joining techniques under development. slowed down or stopped. An even more efficient strat- Similar materials could also be used for the shielding of egy consists in replacing the homogeneous SiC-matrix satellites against the impact of meteorites or foreign ob- itself by an engineered multilayered matrix based on a jects owing to the high toughness and hardness of these similar principle, referred to as a self healing matrix and materials that could be infiltrated by P-CVI at the laborator Another promising field of application is that of the scale.Here, the matrix is deposited as the repetition of hot structures of aerojet engines and related gas turbines a given sequence S comprising thin layers X acting as which are presently made of heavy, low melting point mechanical fuses(X being C, C B), BN, BN (Si), or any nickel-based superalloys that require complex coolin suitable fuse) and layers Y of species forming fluid ox- systems. Replacing superalloys by SiC-based composites ides, as mentioned above(Fig. 3b). Durability of the would permit to raise the gas temperature, suppress,or 1100@C has been reported for SiC fiber composites fab- ciency of the engine, and reduce both the weight anda) ricated with such self-healing matrices. Finally, noise/pollution level(Fig. 4b). However, it will proba ific multilayered coatings containing oxide layers bly take some time(these materials are still your as mullite or/and baryum strontium aluminosilicate, their fabrication costly) and be limited in a first step to BSAS) have been proposed and tested for SiC/SiC com- nonrotating parts, i. e,, the combustors and the after- posites exposed to wet oxidizing atmospheres to reduce burner parts such as the flaps of the exhaust nozzles. The the recession rate of the materials, e.g., in hot combu main concern here is durability, that should be of the tion gas rich in water vapor order of several thousands of hours. The outer(diver gent) Aaps of exhaust nozzles experience a temperature R applic that is relatively low(T<700oC). Hence, they can be fabricated with carbon fiber-reinforced SiC-matrix, with Space and Aeronautic Field a weight gain of 50%. Tested in Hight as early as 1989, they are now in volume production(M88 Snecn ma en- SiC-matrix composites are potential material can- gines of the Rafale fighter, Paris, France). 00, The didates for the fabrication of hot structures of spacecraft inner(convergent) Aaps of the exhaust nozzles are ex- as demonstrated at the prototype part level some years posed to higher temperatures(up to 1100C). Durabil go within the scope of the Hermes European space iry of the order of 1000 h has been demonstrated for shuttle project. Here, the maximum temperature 3D-composites with a self-healing multilayered matrix ranges from 800%C to 1600%C, during the ascent and on the basis of bench combustion tests. The next step is re-entry phases of a fight, the structures being submit- the combustion chamber or combustor, whose fabrica- ted to thermal shocks and cyclic mechanical loading tion with SiC-based composites is in progress. No re- under ablative or passive oxidizing atmospheres, with an sults of tests are presently available for military aerojet expected durability of a few tens of hours( Fig. 4a). Such engine combustors in the open literature, as far as we conditions are compatible with modern composites fab- know. However, the use of SiC-matrix combustor in ricated with carbon fibers(to reduce weight and achieve power plant gas turbine of cogeneration is well docu good mechanical properties at the highest temp mented, with similar (not to say more severe)service and engineered multilayered self-healing matrix. How- condition. 2.33 Combustors of large size comprising ever, some environmental barrier coating(EBC)might concentric cylindrical CMC liners have been fabricated undaccessary to limit the recession rate of the material by CVI or RMI with SiC(Hi-Nicalon)/BN/SiC(Si) under active oxidation/ablation regime(hT and veloc- composites. Durability of several 10,000 h has been es- ity combined with low P(O2). Such a material ap tablished under real service conditions, for composites proach will benefit from the high refractoriness of C/ with a BSAS-EBC.' To conclude, the use of SiC-ma- SiC composites(a 2500%C)relative to the low melting trix composites in the hot nonrotating parts of gas tur- point of aluminum(a 650%C) in the metallic option bines(aerojet engines and cogeneration gas turbines combined with a thermal insulation. Further, the CVI- appears to be promising. It is now a matter of engi process is well suited to the fabrication of large size neering, reliability, and cost(that of performant SiC structures(two meters or more), its feasibility already fibers still remaining relatively dissuasive
fluid oxides and the in-depth diffusion of oxygen is slowed down or stopped.19 An even more efficient strategy consists in replacing the homogeneous SiC-matrix itself by an engineered multilayered matrix based on a similar principle, referred to as a self-healing matrix and that could be infiltrated by P-CVI at the laboratory scale.27 Here, the matrix is deposited as the repetition of a given sequence S comprising thin layers X acting as mechanical fuses (X being C, C (B), BN, BN (Si), or any suitable fuse) and layers Y of species forming fluid oxides, as mentioned above (Fig. 3b). Durability of the order of 1000 h under cyclic loading in air up to 11001C has been reported for SiC fiber composites fabricated with such self-healing matrices.28 Finally, specific multilayered coatings containing oxide layers (such as mullite or/and baryum strontium aluminosilicate, BSAS) have been proposed and tested for SiC/SiC composites exposed to wet oxidizing atmospheres to reduce the recession rate of the materials, e.g., in hot combustion gas rich in water vapor.29 Representative Applications Space and Aeronautic Field SiC-matrix composites are potential material candidates for the fabrication of hot structures of spacecraft, as demonstrated at the prototype part level some years ago within the scope of the Hermes European space shuttle project.30 Here, the maximum temperature ranges from 8001C to 16001C, during the ascent and re-entry phases of a flight, the structures being submitted to thermal shocks and cyclic mechanical loading under ablative or passive oxidizing atmospheres, with an expected durability of a few tens of hours (Fig. 4a). Such conditions are compatible with modern composites fabricated with carbon fibers (to reduce weight and achieve good mechanical properties at the highest temperatures) and engineered multilayered self-healing matrix. However, some environmental barrier coating (EBC) might be necessary to limit the recession rate of the material under active oxidation/ablation regime (HT and velocity combined with low P(O2)). Such a material approach will benefit from the high refractoriness of C/ SiC composites ( 25001C) relative to the low melting point of aluminum ( 6501C) in the metallic option combined with a thermal insulation. Further, the CVIprocess is well suited to the fabrication of large size structures (two meters or more), its feasibility already established and joining techniques under development. Similar materials could also be used for the shielding of satellites against the impact of meteorites or foreign objects owing to the high toughness and hardness of these materials. Another promising field of application is that of the hot structures of aerojet engines and related gas turbines, which are presently made of heavy, low melting point nickel-based superalloys that require complex cooling systems. Replacing superalloys by SiC-based composites would permit to raise the gas temperature, suppress, or at least limit the cooling requirement, increase the effi- ciency of the engine, and reduce both the weight and the noise/pollution level (Fig. 4b). However, it will probably take some time (these materials are still young and their fabrication costly) and be limited in a first step to nonrotating parts, i.e., the combustors and the afterburner parts such as the flaps of the exhaust nozzles. The main concern here is durability, that should be of the order of several thousands of hours. The outer (divergent) flaps of exhaust nozzles experience a temperature that is relatively low (To7001C). Hence, they can be fabricated with carbon fiber-reinforced SiC-matrix, with a weight gain of 50%. Tested in flight as early as 1989, they are now in volume production Q3 (M88 Snecma engines of the Rafale fighter, Paris, France).28,30,31 The inner (convergent) flaps of the exhaust nozzles are exposed to higher temperatures (up to 11001C). Durability of the order of 1000 h has been demonstrated for 3D-composites with a self-healing multilayered matrix on the basis of bench combustion tests. The next step is the combustion chamber or combustor, whose fabrication with SiC-based composites is in progress. No results of tests are presently available for military aerojet engine combustors in the open literature, as far as we know. However, the use of SiC-matrix combustor in power plant gas turbine of cogeneration is well documented, with similar (not to say more severe) service condition.32,33 Combustors of large size comprising concentric cylindrical CMC liners have been fabricated by CVI or RMI with SiC (Hi-Nicalon)/BN/SiC (Si) composites. Durability of several 10,000 h has been established under real service conditions, for composites with a BSAS–EBC.29 To conclude, the use of SiC-matrix composites in the hot nonrotating parts of gas turbines (aerojet engines and cogeneration gas turbines) appears to be promising. It is now a matter of engineering, reliability, and cost (that of performant SiC fibers still remaining relatively dissuasive). www.ceramics.org/ACT SiC-Matrix Composites: Application 81
iternational Journal of Applied Ceramic Technology-Naslain Vol.2,No.2,2005 C-SiC shingle antenna RF recevltrans nstrated 1100C C-Sic wing shingle TPs shingle TPS plasma fest 1350C C-SIC box ass C-SiC elevons body nap Fig 4. Potential applications of Sic-matrix composite materials: (a)in spacecraft hot structure and(b) in aerojet engine. Adapted from Christin se Braking systems is another field of applications. Up near the ambient temperature are altered in wet atmos- until recently, aircraft and racing car disk brakes were phere and their wear is significant. It has been proposed fabricated with C/C composites sliding against them- recently to replace part of the C-matrix by a SiC-based selves(disk/disk or disk/pad configuration). C/C brakes matrix,the new matrix being formed by a combination display, relative to conventional steel disk brakes, a low of carbon-PIP and SiC (Si)-RMI techniques, as depicted er density, a higher service temperature(and hence im- previously. 4.35 Coated C/SiC (Si)brake disks, sliding proved security), good friction properties at HT, and against themselves or pads(organic or metallic), show a longer lifetime. Conversely, their friction properties higher coefficient of friction less depending on moisture
Braking systems is another field of applications. Up until recently, aircraft and racing car disk brakes were fabricated with C/C composites sliding against themselves (disk/disk or disk/pad configuration). C/C brakes display, relative to conventional steel disk brakes, a lower density, a higher service temperature (and hence improved security), good friction properties at HT, and longer lifetime. Conversely, their friction properties near the ambient temperature are altered in wet atmosphere and their wear is significant. It has been proposed recently to replace part of the C-matrix by a SiC-based matrix, the new matrix being formed by a combination of carbon-PIP and SiC (Si)-RMI techniques, as depicted previously.34,35 Coated C/SiC (Si) brake disks, sliding against themselves or pads (organic or metallic), show a higher coefficient of friction less depending on moisture Fig. 4. Potential applications of SiC-matrix composite materials: (a) in spacecraft hot structure and (b) in aerojet engine. Adapted from Christin.30 82 International Journal of Applied Ceramic Technology—Naslain Vol. 2, No. 2, 2005
wwceramics. org/ACT and a very low wear. They are now proposed as an op- microstructure scale, a toughness of the order of 10 kj/ tion for cars in Europe and could be extended to other m, and a good fatigue resistance. They creep at high fields(trains, lifts, etc. temperature but with a creep rate that is lower than that Finally, Sic-matrix composites might become key of metals. Their thermal conductivity is relatively high structural materials in high temperature nuclear reactors when their constituents are well crystallized and their of the future(e.g, the tokamak fusion reactors in which residual porosity is low. Their oxidation resistance is heat generated by the fusion reaction is extracted better at 1000-1200oC than at lower temperatures due through the wall of the toroidal plasma chamber with to the formation of protective oxide scales(SiO2, a cooling agent such as gaseous helium, at a temperature B2O3), in the so-called passive regime. It can be im of 800-1000C), on the basis of their refractoriness, proved with multilayered self-healing matrices and coat HT-mechanical properties, high thermal conductivity, ings, based on the formation of Auid healing oxides. The and more importantly low activation under radia- oxidation resistance is degraded in wet atmospheres tion.36-39 This domain is becoming an active field of where specific EBCs needed to be used research, the main concerns being: the combined effect SiC-matrix composites are matured enough to be of temperature and radiation on the structure(with ei- utilized in a variety of applications including the hot ther a shrinkage or swelling) and mechanical properties structures of spacecraft, aerojet engines, and gas turbines of both SiC fiber and matrix, the nature of the inter- of cogeneration, some parts being already in volume (thick layers of carbon or BN being inappropriate production, with a weight gain of a 50% versus su under radiation), the effect of activable impurities in the peralloys and a durability that can be several thousands composites, the thermal conductivity, and the corrosion of hours. They are also promising materials for braking by residual or in situ formed gaseous species(oxygen or systems and HT nuclear reactors of the future helium in fusion reactors). It appears from preliminary data that the composites should be better fabricated with fibers and matrix consisting of crystalline p-SiC Acknowledgments with a low impurity content (additives introduced as The author acknowledges the collaboration during sintering aids in fibers or/and matrix being a subject of many years of the senior researchers from LCTS and concern), with either porous SiC,(PyC-SiC)m, or(BN- engineers from Snecma and CEA when he was in charge SiC)n multilayered interphases of low C or BN content of the research program on SiC-matrix composites at and displaying a low residual porosity(for hermeticity LCTS and the support of CNRS, Snecma, CEA, and thermal conductivity considerations). Further and Bordeaux University, and the Aquitaine Regional mentioned previously, large size structures could be Authority. CVI/RMI combined ick ing technologies(PIP/RMI or bricated by alread chniques and joinin References 1. F. Christin, R. Naslain, an SiC-matrix composites, i.e., C/SiC and SiC/SiC, d H. Lydrin. The Electrochemical can be fabricated from different carbon or sic fibers playing a variety of properties, usually by single(such 2. R Naslain, J. Y. Rossignol, P. Hagenmuller, F. Christin, L. Heraud, and J. ). lry. "Synthesis and Properties of New Composite Materials for H as Cvi)or combined processes(e.g, PIP/RMI or CVI RMD). They are damage-tolerant when the fibers and 3. I.J. Brennan, "Interfacial Characterization of Glass and Glass-ceras matrix are bonded together with an interphase that can atrix/Nicalon SiC Fiber Composites, Tailoring Multiphase and Comp be a single relatively thick layer of pyrocarbon (or BN) Ceramie. eds. R. E. Tressler, G. L Messing, C.G. Pantano, and R.E. wham. Plenum Press, New York, 549-570. 1986 a porous single layer of SiC, or an engineered multi- 4. R Naslain, The Design of the Fibre- Marix Interfacial Zone in Ceramic layered (X-Y)n interphase (with X= PyC or BN and y= SiC)offering more design Flexibility. The compos- 6.R Nalain, "Design, Preparation and Properies of Non-Oxide CMCs for latrix Composites, Actd Metal, 37[10]2567-2583(1989) ites display a nonlinear stress-strain behavior under Application in Engines and Nuclear Reactors: An Overview, Composite Sci ensile loading related to damaging phenomena at the Technol.,6155-170(2004)
and a very low wear. They are now proposed as an option for cars in Europe and could be extended to other fields (trains, lifts, etc.). Finally, SiC-matrix composites might become key structural materials in high temperature nuclear reactors of the future (e.g., the tokamak fusion reactors in which heat generated by the fusion reaction is extracted through the wall of the toroidal plasma chamber with a cooling agent such as gaseous helium, at a temperature of 800–10001C), on the basis of their refractoriness, HT-mechanical properties, high thermal conductivity, and more importantly low activation under radiation.36–39 This domain is becoming an active field of research, the main concerns being: the combined effect of temperature and radiation on the structure (with either a shrinkage or swelling) and mechanical properties of both SiC fiber and matrix, the nature of the interphase (thick layers of carbon or BN being inappropriate under radiation), the effect of activable impurities in the composites, the thermal conductivity, and the corrosion by residual or in situ formed gaseous species (oxygen or helium in fusion reactors). It appears from preliminary data that the composites should be better fabricated with fibers and matrix consisting of crystalline b-SiC with a low impurity content (additives introduced as sintering aids in fibers or/and matrix being a subject of concern), with either porous SiC, (PyC–SiC)n, or (BN– SiC)n multilayered interphases of low C or BN content, and displaying a low residual porosity (for hermeticity and thermal conductivity considerations). Further and as mentioned previously, large size structures could be fabricated by already existing technologies (PIP/RMI or CVI/RMI combined techniques and joining). Conclusion SiC-matrix composites, i.e., C/SiC and SiC/SiC, can be fabricated from different carbon or SiC fibers displaying a variety of properties, usually by single (such as CVI) or combined processes (e.g., PIP/RMI or CVI/ RMI). They are damage-tolerant when the fibers and matrix are bonded together with an interphase that can be a single relatively thick layer of pyrocarbon (or BN), a porous single layer of SiC, or an engineered multilayered (X–Y )n interphase (with X 5 PyC or BN and Y 5 SiC) offering more design flexibility. The composites display a nonlinear stress–strain behavior under tensile loading related to damaging phenomena at the microstructure scale, a toughness of the order of 10 kJ/ m2 , and a good fatigue resistance. They creep at high temperature but with a creep rate that is lower than that of metals. Their thermal conductivity is relatively high when their constituents are well crystallized and their residual porosity is low. Their oxidation resistance is better at 1000–12001C than at lower temperatures due to the formation of protective oxide scales (SiO2, B2O3), in the so-called passive regime. It can be improved with multilayered self-healing matrices and coatings, based on the formation of fluid healing oxides. The oxidation resistance is degraded in wet atmospheres where specific EBCs needed to be used. SiC-matrix composites are matured enough to be utilized in a variety of applications including the hot structures of spacecraft, aerojet engines, and gas turbines of cogeneration, some parts being already in volume production, with a weight gain of 50% versus superalloys and a durability that can be several thousands of hours. They are also promising materials for braking systems and HT nuclear reactors of the future. Acknowledgments The author acknowledges the collaboration during many years of the senior researchers from LCTS and engineers from Snecma and CEA when he was in charge of the research program on SiC-matrix composites at LCTS and the support of CNRS, Snecma, CEA, Bordeaux University, and the Aquitaine Regional Authority. References 1. F. Christin, R. Naslain, and C. Bernard, ‘‘A Thermodynamic and Experimental Approach of Silicon Carbide CVD. Application to the CVD-Infiltration of Porous Carbon Composites,’’ Proceedings of the 7th International Conference on CVD. eds. T. O. Sedwick and H. Lydtin. The Electrochemical Society, Princeton, 499–514, 1979. 2. R. Naslain, J. Y. Rossignol, P. Hagenmuller, F. Christin, L. Heraud, and J. J. Choury, ‘‘Synthesis and Properties of New Composite Materials for High Temperature Applications Based on Carbon Fibers and C–SiC or C–TiC Hybrid Matrices,’’ Rev. Chim. Mine´rale, 18 544–564 (1981). 3. J. J. Brennan, ‘‘Interfacial Characterization of Glass and Glass–Ceramic Matrix/Nicalon SiC Fiber Composites,’’ Tailoring Multiphase and Composite Ceramics. eds. R. E. Tressler, G. L. Messing, C. G. Pantano, and R. E. Newnham. Plenum Press, New York, 549–570, 1986. 4. R. Naslain, ‘‘The Design of the Fibre–Matrix Interfacial Zone in Ceramic Matrix Composites,’’ Composites Part A, 29A 1145–1155 (1998). 5. A. G. Evans and D. B. Marshall, ‘‘The Mechanical Behavior of Ceramic Matrix Composites,’’ Acta Metall., 37 [10] 2567–2583 (1989). 6. R. Naslain, ‘‘Design, Preparation and Properties of Non-Oxide CMCs for Application in Engines and Nuclear Reactors: An Overview,’’ Composite Sci. Technol., 64 155–170 (2004). www.ceramics.org/ACT SiC-Matrix Composites: Application 83
International y ournal of Applied Ceramic TechnologyNaslain Vol.2,No.2,2005 7. R. Naslain, "CVI Com 25. F. Lamouroux, R. Naslain, and J M. Jouin, " Kinetics and Mechanisms C/SiC Composites: 2-Theorerical Approach 8. T. M. Besmann, " CVI Processing of Ceramic Matrix Composites, "Ceram. Am. Ceram Soc., 77 [8 2058-2 9. K L Luthra, R N Singh, and M.K. Brun, " Toughened Silcomp Com- Sac,868]1238-1248(2003) posites-Process and Preliminary Properties, "Am. Ceram. Sac. Bull, 72[7] 27. F Lamouroux S Bertrand, R. Paillet, R. Naslain, and M. Cataldi, Oxida- on-Resistant Carbon Fiber Reinforced Ceramic-Matrix Composites, " Com- O.K. Nakano, K. Suzuki, M. Drissi-Habti, and Y. Kanno, "Processing and posites Sci. Technol, 59 1073-1085(1999) Characterization of 3 D-Carbon Fiber Reinforced Silicon Carbide and Silicon 28. F. Lamouroux, E. Bouillon, J. C. Cavalier, P. 11. Y Katoh, S-M. Dong, and A Kohyama,A Novel Processing Technique of e. eds Silicon Carbide Based Ceramic Composites for High Temperature Applica m,783-78 477-86(2002). 12. R. J. Kerans, R.S. Hay, T. A Parthasarathy, and M. K. Cinibulk, "Interface 29. K. N. Lee, D. S. Fox. ]. I. Eldridge, D-M. Zhu, R C. Robinson, N P. Bansal, sign for Oxidation Resistant Ceramic Composites, ". Amer. Ceram. Soc, and R. A. Miller, "Upper Temperature Limir of Envi Coatings Based on Mullite and BSAS, "J. Am. Ceram Soc., 86[8]1299-1309 13. W-C. Tu, F. F. Lange, and A G. Evans, Concept of Damage-Tolerant Ceramic Composite with"Strong" Interfaces, " / Am. Ceram. Soc., 79[2] 30. F. Christin, "Design, Fabrication and Application of C /C, C/SiC and Sic/ SiC Composites, High Temperature Ceramic Matrix Composite. eds. W 4. F. F. Lange, T. C. Radsick, and M. Holmquist, Oxide/Oxide Compo. Krenkel, R Naslain, and H. Schneider. Wiley-VCH, Weinheim, 732-74 Control of Microstructure and Properties, High Temperature Cenamic ix Copasites. eds. W. Krenkel, R. Naslain, and H. Schneider. Wiley-V 31. F. Christin, "Design, Fabrication and Applications of Thermostructural Like C/C, C/SiC C/SiC Composites, Ade. En 15. C. Droillard, Processing and Characterization of Meter,4[12]903912(2002) Multilayered C/SiC Interphase, "PhD Thesis, 32. K. L Luthra and G.S. red (Mi) SiC/SiC Composites for Gas Turbine Applications, "Hi 16. S Le Galler, F. Rebillar, A. Guerte, and R. Naslain in Air at An ds. W. Krenkel, R. Naslain, and H. Schneider. Wiley-VCH, Weinheim, bient Temperature of BN Coatings Processed from BCla-NH3-H3 744-753,2001 nures, High Temperature Ceramic Matrix Composites. eds. W. Krenkel, 33. W.D. Brentnall, M. van Roode, P. F. Norton, S Gates, J.L Rice, O. Jime- R. Naslain, and H Schneider. Wiley-VCH, Weinheim, 187-192. 2001 nez, and N. Miriyala, Ceramic Gas Turbine Development 17. H. M. Yun, J.Z. G corporated, "Ceramic Gas Turbine Deign and Test Experience, Vol. I. eds Tensile Behavior of SiC/SiC Composites Reinforced by Treated Sylramic M van Roode, M. K. Ferber, and D. w. Richerson ASME Press, New York, 155-192,2002. 18. S. Bertrand, "Improvement of the Durability of SiC/SiC Composites with 34. W. Krenkel, C/C-SiC Composites for Hot Structures and PyC-SiC)n or(BN-SiC) Interphases, " PhD Thesis, Friction Systems, Ceram. Eng. Sci. Proc., 24(4] 483-492(2003) 19. S. Goujard and L Vandenbulcke, "Deposition of Si-B-C Maerials from the ation for Aircraft Brake Applications, High Temperature Ceram aterial Evak. Vapor Phase for A Composites, eds. W. Krenkel, R. Naslain, and H. Schneider. Wiley-VCH, U. Ramamurty, J. C. McNulty, and M. Steen, " Fatigue in Ceramic Matrix 36. R H. Jones, D. Steiner, H. L Heinish, G.A. Newsome, and H M. Herc Composites, "Comprelensive Compesite Materials, Vol. 4. ed. R Warren, Radiation Resistant Ceramic Matrix Composites, J. Nuc. Mater, 245 21. G. Fantozzi, J. Chevalier, C. OLagnon, and J. L Chermant, "Creep of 37. L Giancarli, J. P. Bonal, A. Caso, G. Le Marois, N. B Morley, and J. F. ceramic Matrix Composites, Comprehensire Composite Material, VoL. 4. Salavy. "Design Requirements for SiC/SiC Composites Structural Material 22. R. Taylor, "Carbon Matrix Composites, "Comprehensie Composite Materials, Vol. 4. ed. R. Warren. Elsevier, Amsterdam, 387-426, 20 38.R H. Jones, LL. Snead, A Kohyama, and P Fenici, "Recent Advances in 23. R. Yamada, N. Igawa, and T. Taguchi, "A Finite-Element Analysis the Development of SiC/SiC as a Fusion Structural Material, Fusion Eng Thermal Diffusivity/Conductivity of Sic/SiC Composites, Ceram. Tre Design,4115-24(1998) 144289-299(200 39.R H. Jones, L. Giancarli, A. Hasegawa, Y. Katoh, A. Kohyama. B. Riccardi, d, and w.J. Weber, "Promise an iC/C/SiC Composite Materials, 2-Modelling, "/. m. Ceram. Soc., 77 [81 osites for Fusion Energy Applications, " /.Nucl. Mater, 307-311 1057- 1072(2002)
7. R. Naslain, ‘‘CVI Composites,’’ Ceramic Matrix Composites. ed. R. Warren. Blackie, Glasgow, 199–244, 1992. 8. T. M. Besmann, ‘‘CVI Processing of Ceramic Matrix Composites,’’ Ceram. Trans., 58 1–12 (1995). 9. K. L. Luthra, R. N. Singh, and M. K. Brun, ‘‘Toughened Silcomp Composites—Process and Preliminary Properties,’’ Am. Ceram. Soc. Bull., 72 [7] 79–85 (1993). 10. K. Nakano, K. Suzuki, M. Drissi-Habti, and Y. Kanno, ‘‘Processing and Characterization of 3D-Carbon Fiber Reinforced Silicon Carbide and Silicon Nitride Matrix Composites,’’ Ceram. Trans., 99 157–166 (1998). 11. Y. Katoh, S-M. Dong, and A. Kohyama, ‘‘A Novel Processing Technique of Silicon Carbide Based Ceramic Composites for High Temperature Applications,’’ Ceram. Trans., 144 77–86 (2002). 12. R. J. Kerans, R. S. Hay, T. A. Parthasarathy, and M. K. Cinibulk, ‘‘Interface Design for Oxidation Resistant Ceramic Composites,’’ J. Amer. Ceram. Soc., 85 [11] 2599–2632 (2002). 13. W-C. Tu, F. F. Lange, and A. G. Evans, ‘‘Concept of Damage-Tolerant Ceramic Composite with ‘‘Strong’’ Interfaces,’’ J. Am. Ceram. Soc., 79 [2] 417–424 (1996). 14. F. F. Lange, T. C. Radsick, and M. Holmquist, ‘‘Oxide/Oxide Composites: Control of Microstructure and Properties,’’ High Temperature Ceramic Matrix Composites. eds. W. Krenkel, R. Naslain, and H. Schneider. Wiley-VCH, Weinheim, 587–599, 2001. 15. C. Droillard, ‘‘Processing and Characterization of SiC-Matrix Composites with Multilayered C/SiC Interphase,’’ PhD Thesis, No. 913, University of Bordeaux 1, June 19, 1993. 16. S. Le Gallet, F. Rebillat, A. Guette, and R. Naslain, ‘‘Stability in Air at Ambient Temperature of BN Coatings Processed from BCl3–NH3–H2 Gas Mixtures,’’ High Temperature Ceramic Matrix Composites. eds. W. Krenkel, R. Naslain, and H. Schneider. Wiley-VCH, Weinheim, 187–192, 2001. 17. H. M. Yun, J. Z. Gyekenyesi, Y. L. Chen, D. R. Wheeler, and J. A. DiCarlo, ‘‘Tensile Behavior of SiC/SiC Composites Reinforced by Treated Sylramic SiC Fibers,’’ Ceram. Eng. Sci. Proc., 22 [3] 521–531 (2001). 18. S. Bertrand, ‘‘Improvement of the Durability of SiC/SiC Composites with Multilayered (PyC–SiC)n or (BN–SiC)n Interphases,’’ PhD Thesis, no. 1927, University of Bordeaux 1, September 29, 1998 19. S. Goujard and L. Vandenbulcke, ‘‘Deposition of Si–B–C Maerials from the Vapor Phase for Applications in Ceramic Matrix Composites,’’ Ceram. Trans., 46 925–935 (1994). 20. U. Ramamurty, J. C. McNulty, and M. Steen, ‘‘Fatigue in Ceramic Matrix Composites,’’ Comprehensive Composite Materials, Vol. 4. ed. R. Warren, Elsevier, Amsterdam, 163–219, 2000. 21. G. Fantozzi, J. Chevalier, C. Olagnon, and J. L. Chermant, ‘‘Creep of Ceramic Matrix Composites,’’ Comprehensive Composite Materials, Vol. 4. ed. R. Warren. Elsevier, Amsterdam, 115–162, 2000. 22. R. Taylor, ‘‘Carbon Matrix Composites,’’ Comprehensive Composite Materials, Vol. 4. ed. R. Warren. Elsevier, Amsterdam, 387–426, 2000. 23. R. Yamada, N. Igawa, and T. Taguchi, ‘‘A Finite-Element Analysis of the Thermal Diffusivity/Conductivity of SiC/SiC Composites,’’ Ceram. Trans., 144 289–299 (2002). 24. L. Filipuzzi, and R. Naslain, ‘‘Oxidation Mechanisms and Kinetics of 1DSiC/C/SiC Composite Materials, 2-Modelling,’’ J. Am. Ceram. Soc., 77 [8] 467–480 (1994). 25. F. Lamouroux, R. Naslain, and J. M. Jouin, ‘‘Kinetics and Mechanisms of Oxidation of 2D Woven C/SiC Composites: 2—Theoretical Approach,’’ J. Am. Ceram. Soc., 77 [8] 2058–2068 (1994). 26. E. J. Opila, ‘‘Oxidation and Volatilization of Silica Formers in Water Vapor,’’ J. Am. Ceram. Soc., 86 [8] 1238–1248 (2003). 27. F. Lamouroux, S. Bertrand, R. Pailler, R. Naslain, and M. Cataldi, ‘‘Oxidation-Resistant Carbon Fiber Reinforced Ceramic-Matrix Composites,’’ Composites Sci. Technol., 59 1073–1085 (1999). 28. F. Lamouroux, E. Bouillon, J. C. Cavalier, P. Spriet, and G. Habarou, ‘‘An improved Long Life Duration CMC for Jet Aircraft Engine Applications,’’ High Temperature Ceramic Matrix Composites. eds. W. Krenkel, R. Naslain, and H. Schneider. Wiley-VCH, Weinheim, 783–788, 2001. 29. K. N. Lee, D. S. Fox, J. I. Eldridge, D-M. Zhu, R. C. Robinson, N. P. Bansal, and R. A. Miller, ‘‘Upper Temperature Limit of Environmental Barrier Coatings Based on Mullite and BSAS,’’ J. Am. Ceram. Soc., 86 [8] 1299–1309 (2003). 30. F. Christin, ‘‘Design, Fabrication and Application of C/C, C/SiC and SiC/ SiC Composites,’’ High Temperature Ceramic Matrix Composites. eds. W. Krenkel, R. Naslain, and H. Schneider. Wiley-VCH, Weinheim, 732–743, 2001. 31. F. Christin, ‘‘Design, Fabrication and Applications of Thermostructural Composites (TSC) Like C/C, C/SiC and SiC/SiC Composites,’’ Adv. Eng. Mater., 4 [12] 903–912 (2002). 32. K. L. Luthra and G. S. Corman, ‘‘Melt Infiltrated (MI) SiC/SiC Composites for Gas Turbine Applications,’’ High Temperature Ceramic Matrix Composites. eds. W. Krenkel, R. Naslain, and H. Schneider. Wiley-VCH, Weinheim, 744–753, 2001. 33. W. D. Brentnall, M. van Roode, P. F. Norton, S. Gates, J. L. Rice, O. Jimenez, and N. Miriyala, ‘‘Ceramic Gas Turbine Development at Solar Turbines Incorporated,’’ Ceramic Gas Turbine Design and Test Experience, Vol. 1. eds. M. van Roode, M. K. Ferber, and D. W. Richerson. ASME Press, New York, 155–192, 2002. 34. W. Krenkel, ‘‘C/C–SiC Composites for Hot Structures and Advanced Friction Systems,’’ Ceram. Eng. Sci. Proc., 24 [4] 483–492 (2003). 35. S. Vaidyaraman, M. Purdy, T. Walker, and S. Horst, ‘‘C/SiC Material Evaluation for Aircraft Brake Applications,’’ High Temperature Ceramic Matrix Composites. eds. W. Krenkel, R. Naslain, and H. Schneider. Wiley-VCH, Weinheim, 802–808, 2001. 36. R. H. Jones, D. Steiner, H. L. Heinish, G. A. Newsome, and H. M. Herch, ‘‘Radiation Resistant Ceramic Matrix Composites,’’ J. Nucl. Mater., 245 87–107 (1997). 37. L. Giancarli, J. P. Bonal, A. Caso, G. Le Marois, N. B. Morley, and J. F. Salavy, ‘‘Design Requirements for SiC/SiC Composites Structural Material in Fusion Power Reactor Blankets,’’ Fusion Eng. Design, 41 165–171 (1998 Q5 ). 38. R. H. Jones, L. L. Snead, A. Kohyama, and P. Fenici, ‘‘Recent Advances in the Development of SiC/SiC as a Fusion Structural Material,’’ Fusion Eng. Design, 41 15–24 (1998 Q6 ). 39. R. H. Jones, L. Giancarli, A. Hasegawa, Y. Katoh, A. Kohyama, B. Riccardi, L. L. Snead, and W. J. Weber, ‘‘Promise and Challenges of SiC/SiC Composites for Fusion Energy Applications,’’ J. Nucl. Mater., 307–311 1057– 1072 (2002). 84 International Journal of Applied Ceramic Technology—Naslain Vol. 2, No. 2, 2005