Part A: applied scienc and manufacturing ELSEVIER Composites: Part A 30(1999)411-417 Fiber-reinforced composites with polymer-derived matrix: processing, matrix formation and properties G. Ziegler, I Richter, D. Suttor University of Bayreuth, Institute for Materials Research(IMA), D-95440 Bayreuth, Germany Carbon and SiC fiber-reinforced ceramic matrix composites were prepared via infiltration of fiber preforms using the polymer infiltration technique and polymer pyrolysis. Suitable silazane (SicN) precursors with appropriate thermosetting behavior, viscosity and ceramic yield were synthesized, starting from functionalized chlorosilanes. Microstructural development and fracture behavior was studied after various infiltration and pyrolysis cycles. Residual stresses induced during processing were evaluated. Mechanical and thermo-physical properties of the composites with polymer-derived matrix, i.e. 3-pt bending strength and thermal expansion coefficients(CTE), were measured dependent on reinfiltration cycles and fiber orientation. The oxidation resistance was investigated. Specific pyrolyzed samples were infiltrated via silicon melts in order to enhance corrosion and wear resistance. C 1999 Elsevier Science Ltd. All rights reserved Keywords: A Ceramic matrix composite(CMCs); Fiber reinforcement, Polymer pyrolysis; Silazane precursor 1. Introduction Furthermore, linear shrinkage values of typically 2 15% during sintering complicate the realization of dense matrices Ceramic matrix composites( CMCs)are potential candi- [2]. CVI techniques have been developed in the 1980s to dates for applications requiring damage tolerance, high form matrices in a multiple step process by infiltrating 3D strength at elevated temperatures and low specific weight preforms via the reaction of gaseous precursors. The isother- ]. However, other areas of application are now under nal CVI process, however, is very time and cost intensive investigation, e.g. pump sealings in chemical engineering while the CVi gradient technique is faster, but limited in or brake discs in transportation systems. In the future, strong shape complexity 3] emphasis will be put on suitable and thus affordable proces An alternative method is based on organometallic poly sin g techniques, enabling complex geometries without meric SiC/Si3 n4 precursors, e. g silazanes or carbosilane damaging the fiber reinforcement. Thermo-chemical and Upon pyrolysis, the organometallic polymeric precursor is physical compatibility between the matrix and fiber is converted into a ceramic material yielding a Sicn glass, necessary in order to control the interface characteristics, SiC, Si3N4 or mixtures thereof. Such polymers can be and thus the mechanical properties. In the case of carbon used for matrix formation of CMCs by infiltration of porous fiber-based composites and high in service temperatures, fiber preforms or single fiber lay-up, enabling the realization usually an external and internal oxidation protection system of very complex shapes by using standard techniques as has to be applied. SiC-based fibers are less susceptible to common in the polymer community [4]. The fabrication xidation, however, economic arguments have then to be of CMCs includes the liquid impregnation of fiber preforms considered. The property profile of these composites also (or single fibers) with suitable precursors(e.g. low viscosity, includes factors, e.g. wear and corrosion resistance, for solvent free) followed by subsequent crosslinking to ther which up to now only a few data have been presented moses(resulting in fiber-reinforced polymers) and pyroly Various processing methods are used. They are based on sis at temperatures s 1000 C in inert atmospheres to yield a ceramic slurry techniques, in situ chemical reactions, e.g. fiber-reinforced ceramic [5]. These polymers contain vinyl CVI(chemical vapor infiltration) or liquid infiltration as well as hydrogen groups, enabling thermally activated processes, e.g. polymer pyrolysis. The slurry method is setting at temperatures up to 300C. However, with respect limited with respect to shape comple and is restricted to infiltration techniques, e.g. resin transfer molding(RTM by infiltration depth and ther ity. processing requires lower setting temperatures in order to simplify the mold design. Suitable catalysts are, e.g orresponding author peroxides, especially dicumylperoxide, enabling setting 835X/99/- see front e 1999 Elsevier Science Ltd. All rights reserved 1359-835X(98)00128-6
Fiber-reinforced composites with polymer-derived matrix: processing, matrix formation and properties G. Ziegler*, I. Richter, D. Suttor University of Bayreuth, Institute for Materials Research (IMA), D-95440 Bayreuth, Germany Abstract Carbon and SiC fiber-reinforced ceramic matrix composites were prepared via infiltration of fiber preforms using the polymer infiltration technique and polymer pyrolysis. Suitable silazane (SiCN) precursors with appropriate thermosetting behavior, viscosity and ceramic yield were synthesized, starting from functionalized chlorosilanes. Microstructural development and fracture behavior was studied after various infiltration and pyrolysis cycles. Residual stresses induced during processing were evaluated. Mechanical and thermo-physical properties of the composites with polymer-derived matrix, i.e. 3-pt bending strength and thermal expansion coefficients (CTE), were measured dependent on reinfiltration cycles and fiber orientation. The oxidation resistance was investigated. Specific pyrolyzed samples were infiltrated via silicon melts in order to enhance corrosion and wear resistance. q 1999 Elsevier Science Ltd. All rights reserved. Keywords: A. Ceramic matrix composite (CMCs); Fiber reinforcement; Polymer pyrolysis; Silazane precursor 1. Introduction Ceramic matrix composites (CMCs) are potential candidates for applications requiring damage tolerance, high strength at elevated temperatures and low specific weight [1]. However, other areas of application are now under investigation, e.g. pump sealings in chemical engineering or brake discs in transportation systems. In the future, strong emphasis will be put on suitable and thus affordable processing techniques, enabling complex geometries without damaging the fiber reinforcement. Thermo-chemical and -physical compatibility between the matrix and fiber is necessary in order to control the interface characteristics, and thus the mechanical properties. In the case of carbon fiber-based composites and high in service temperatures, usually an external and internal oxidation protection system has to be applied. SiC-based fibers are less susceptible to oxidation, however, economic arguments have then to be considered. The property profile of these composites also includes factors, e.g. wear and corrosion resistance, for which up to now only a few data have been presented. Various processing methods are used. They are based on ceramic slurry techniques, in situ chemical reactions, e.g. CVI (chemical vapor infiltration) or liquid infiltration processes, e.g. polymer pyrolysis. The slurry method is limited with respect to shape complexity, and is restricted by infiltration depth and therefore homogeneity. Furthermore, linear shrinkage values of typically $ 15% during sintering complicate the realization of dense matrices [2]. CVI techniques have been developed in the 1980s to form matrices in a multiple step process by infiltrating 3D preforms via the reaction of gaseous precursors. The isothermal CVI process, however, is very time and cost intensive, while the CVI gradient technique is faster, but limited in shape complexity [3]. An alternative method is based on organometallic polymeric SiC/Si3N4 precursors, e.g. silazanes or carbosilanes. Upon pyrolysis, the organometallic polymeric precursor is converted into a ceramic material yielding a SiCN glass, SiC, Si3N4 or mixtures thereof. Such polymers can be used for matrix formation of CMCs by infiltration of porous fiber preforms or single fiber lay-up, enabling the realization of very complex shapes by using standard techniques as common in the polymer community [4]. The fabrication of CMCs includes the liquid impregnation of fiber preforms (or single fibers) with suitable precursors (e.g. low viscosity, solvent free) followed by subsequent crosslinking to thermosets (resulting in fiber-reinforced polymers) and pyrolysis at temperatures # 10008C in inert atmospheres to yield a fiber-reinforced ceramic [5]. These polymers contain vinyl as well as hydrogen groups, enabling thermally activated setting at temperatures up to 3008C. However, with respect to infiltration techniques, e.g. resin transfer molding (RTM), processing requires lower setting temperatures in order to simplify the mold design. Suitable catalysts are, e.g. peroxides, especially dicumylperoxide, enabling setting Composites: Part A 30 (1999) 411–417 1359-835X/99/$ - see front matter q 1999 Elsevier Science Ltd. All rights reserved. PII: S1359-835X(98)00128-6 * Corresponding author
G. Ziegler et al./Composites: Part A 30(1999)411-417 400 Fig. 1. Density and porosity changes of the C/siCn composite upon pyrolysis at 1000.C after three, five and seven infiltration and pyrolysis cycles temperatures as low as 130C. This procedure is, compared polymer-derived matrix are discussed. Particularly, three to the usually applied processing techniques for ceramic topics are investigated: microstructural development and matrices, a low temperature, pressureless process, enabling stress-strain behavior of C and Sic/SiCn composites matrix formation of complex shaped 2D or 3D composites, respectively, the coefficient of thermal expansion depending without damaging the fibers, neither chemically nor on fiber orientation, and oxidation behavior and siliconiza- impregnation and pyrolysis steps are necessary to achieve relatively dense matrices, which can be reduced by synthe- sizing suitable precursors showing a high ceramic yield 2. Experimental procedure A variety of aspects has to be considered if the whole Synthesis of the polymeric silazanes followed standard potential of this approach is going to be used. The polymers procedures, as described elsewhere [7. Preparation of cera- have to be processable, and thus liquid or meltable. Ther- nic matrix composites was based on multiple infiltration mosetting is a requirement in order to retain the shape after and pyrolysis of UD- and 2D(0/90%)fiber prepregs(T300 the forming step, preferably crosslinking via an addition mechanism without the evolution of gaseous products. A 6k, Toray Japan and Hi-Nicalon, Nippon Carbon).Fiber high ceramic yield is favorable in general with respect to bundles were fixed in a metal mold of known volume and the number of necessary infiltration cycles, however, not infiltrated with the liquid thermosetting polymer (fiber only the ceramic yield but actually the residual volume, volume 50%). As a curing catalyst, I wt% of dicumylper which depends on the bulk density of the polymer-derived oxide was added to the polymer, enabling setting tempera- amorphous structures, has to be considered. The density of ares of 130C. After curing, the fiber-reinforced polymers the polymer-derived structures may vary with varying poly were pyrolyzed in a tube furnace up to 1000C in argon and mer density, as has been shown in experiments, which are ambient pressure. Samples were reinfiltrated and pyrolyzed currently conducted. Furthermore, pyrolysis should result in up to seven times(P1-P7). Subsequently,composites a ceramic material with the desired chemical composition characterized by helium pcynometry (AccuPyc 1630 and microstructure. Although the chemical composition can Micromeritics, Germany), scanning electron (JSM be adjusted to some degree in the precursor, it is changing Jeol, Japan) and optical(Leica/Reichert, Polyvar 2Met, upon pyrolysis and crystallization at temperatures Austria)microscopy on polished cros-sections Thermoana- >1000C, as has been shown elsewhere [6]. Therefore, lytical methods(TG/DTA, STA409 and DIL402E, Netzsch, Germany)were used to monitor the associated mass and processing temperatures should be higher than the perspec- dimensional changes. Mechanical properties were evaluated tive application temperature, otherwise properties of the composite will change accordingly. Furthermore, even the using the 3-pt bending method with a span width of 70 mm pyrolysis behavior may change with increasing number of and a crosshead speed of 0. 1 mm/min Specific composite cycles due to increasing density of the matrix and a subse- samples were infiltrated with liquid silicon at 1800 C by quent change in ceramic yield, being dependent on the local Sintec, Buching, Germany, whereby part of the matrix partial pressures of the evolved gaseous species. With was expected to be transformed to a Si/SiC matrix respect to the mechanical properties, residual compressive stresses can develop at the fiber/matrix interface, caused by 3. Results and discussion the large shrinkage (30% linear) of the polymer upon py lysis. In addition, stresses caused by the thermal expansion 3.1 Microstructural development and stress-strain mismatch between fiber and matrix can be induced with a b C/SiCN combination being more critical than a SiC/SiCN one. with the Sic fiber and Sicn matrix having similar Density, poros hermal expansion an physical properties properties of C/Sicn composites with a fiber content of In this work, processing, mechanical and thermo-physical 50 vol% were characterized after three, five and seven infil properties of C and SiC fiber-reinforced composites with tration and pyrolysis cycles(P3, P5, P7). Fig. I shows the
temperatures as low as 1308C. This procedure is, compared to the usually applied processing techniques for ceramic matrices, a low temperature, pressureless process, enabling matrix formation of complex shaped 2D or 3D composites, without damaging the fibers, neither chemically nor mechanically. Using this process, however, multiple impregnation and pyrolysis steps are necessary to achieve relatively dense matrices, which can be reduced by synthesizing suitable precursors showing a high ceramic yield. A variety of aspects has to be considered if the whole potential of this approach is going to be used. The polymers have to be processable, and thus liquid or meltable. Thermosetting is a requirement in order to retain the shape after the forming step, preferably crosslinking via an addition mechanism without the evolution of gaseous products. A high ceramic yield is favorable in general with respect to the number of necessary infiltration cycles, however, not only the ceramic yield but actually the residual volume, which depends on the bulk density of the polymer-derived amorphous structures, has to be considered. The density of the polymer-derived structures may vary with varying polymer density, as has been shown in experiments, which are currently conducted. Furthermore, pyrolysis should result in a ceramic material with the desired chemical composition and microstructure. Although the chemical composition can be adjusted to some degree in the precursor, it is changing upon pyrolysis and crystallization at temperatures .10008C, as has been shown elsewhere [6]. Therefore, processing temperatures should be higher than the perspective application temperature, otherwise properties of the composite will change accordingly. Furthermore, even the pyrolysis behavior may change with increasing number of cycles due to increasing density of the matrix and a subsequent change in ceramic yield, being dependent on the local partial pressures of the evolved gaseous species. With respect to the mechanical properties, residual compressive stresses can develop at the fiber/matrix interface, caused by the large shrinkage (30% linear) of the polymer upon pyrolysis. In addition, stresses caused by the thermal expansion mismatch between fiber and matrix can be induced, with a C/SiCN combination being more critical than a SiC/SiCN one, with the SiC fiber and SiCN matrix having similar physical properties. In this work, processing, mechanical and thermo-physical properties of C and SiC fiber-reinforced composites with polymer-derived matrix are discussed. Particularly, three topics are investigated: microstructural development and stress–strain behavior of C and SiC/SiCN composites, respectively, the coefficient of thermal expansion depending on fiber orientation, and oxidation behavior and siliconization of C/SiC composites. 2. Experimental procedure Synthesis of the polymeric silazanes followed standard procedures, as described elsewhere [7]. Preparation of ceramic matrix composites was based on multiple infiltration and pyrolysis of UD- and 2D (08/908) fiber prepregs (T300 6k, Toray Japan and Hi-Nicalon, Nippon Carbon). Fiber bundles were fixed in a metal mold of known volume and infiltrated with the liquid thermosetting polymer (fiber volume 50%). As a curing catalyst, 1 wt% of dicumylperoxide was added to the polymer, enabling setting temperatures of 1308C. After curing, the fiber-reinforced polymers were pyrolyzed in a tube furnace up to 10008C in argon and ambient pressure. Samples were reinfiltrated and pyrolyzed up to seven times (P1–P7). Subsequently, composites were characterized by helium pcynometry (AccuPyc 1330, Micromeritics, Germany), scanning electron (JSM 6400, Jeol, Japan) and optical (Leica/Reichert, Polyvar 2Met, Austria) microscopy on polished cros-sections. Thermoanalytical methods (TG/DTA, STA409 and DIL402E, Netzsch, Germany) were used to monitor the associated mass and dimensional changes. Mechanical properties were evaluated using the 3-pt bending method with a span width of 70 mm and a crosshead speed of 0.1 mm/min. Specific composite samples were infiltrated with liquid silicon at 18008C by Sintec, Buching, Germany, whereby part of the matrix was expected to be transformed to a Si/SiC matrix. 3. Results and discussion 3.1. Microstructural development and stress–strain behavior Density, porosity, thermal expansion and mechanical properties of C/SiCN composites with a fiber content of 50 vol% were characterized after three, five and seven infiltration and pyrolysis cycles (P3, P5, P7). Fig. 1 shows the 412 G. Ziegler et al. / Composites: Part A 30 (1999) 411–417 Fig. 1. Density and porosity changes of the C/SiCN composite upon pyrolysis at 10008C after three, five and seven infiltration and pyrolysis cycles
G. Ziegler et al./Composites: Part 4 30(1999)411-417 Fig. 2. Comparison of the microstructure of the(a)C/SiCN and (b) Hi-Nicalon/SiCN composites. density change as a function of the number of infiltration the matrix has already reached its saturation crack density and pyrolysis cycles. upon cooling due to the thermal expansion mismatch of fib After three infiltration and pyrolysis cycles, a density of and matrix(a F(axial)0, aM=3.0X 10K) 1. 4 g/cm is reached in the C/SiCN system with an open The residual stresses upon cooling due to thermal expa porosity of 28%, which is further reduced to 16% and sion mismatch between fiber and matrix were calculated 12% after five and seven cycles, respectively. The final using a simple model, as described elsewhere [8], assuming geometrical density is approximately 1.7 g/em,, bulk a stress-free state at the pyrolysis temperature. The calcu- density 1.9 g/cm. A closed porosity of about 10% cannot lated stresses axial and radial between fiber and matrix(C/ be eliminated and is present already after three cycles. Fig. 2 SiCN) were determined to be 780 and 90 MPa, respectively shows the comparison of the microstruture of the C/Sicn Upon cooling, the fiber should therefore shrink away from and Hi-Nicalon/SiCN composites the matrix, which was observed. The matrix is also under Microstructural evolution and conclusions regarding the tension and with a matrix yield strength(as has been experi stress-strain behavior of the C/SiCn composites were mentally determined on monolithic material), an array of monitored on fracture surfaces of tested composite materials regular matrix cracks will develop, as has been confirmed P3-P7(F1g.3) by SEM-micrographs on polished cross-sections As is evident, with an increasing number of infiltration Different results are obtained when working with C cycles, a more continuous matrix develops, but also the coated Hi-Nicalon fibers. Fig. 5 shows the stress-strain fracture characteristics are changing. Fig. 3(b)shows the diagrams of these composites after five and seven infiltra microstructure of a composite after five cycles, clearly tion cycles. The high strength values are obviously due to ith pull-out lengths >100 um. However, with increasing matrix density, the Polished cross-sections of a Hi-Nicalon/SiCn material fracture behavior becomes more brittle, as is seen in were used to get some insight into the matrix formation Fig. 3(c). This can be attributed to the development of (Fig. 6). Various points can be made. First, different infiltra local residual compressive stresses at the fiber interface, tion cycles(marked as I and 2 in Fig. 6) can be distin- caused by the large shrinkage(30% linear) of the polymer guished, appearing as lighter and darker areas. Second, upon pyrolysis. In addition, stresses caused by the thermal looking at the crack patterns in matrix areas adjacent to expansion mismatch between fiber and matrix can develop, the fibers(lst cycle)it is obvious that the matrix may shrink although, as calculations have shown, the interface should onto the fibers (3) but also away from them(4), depending be in tension in this system upon cooling from the proces- on the local geometric conditions. The arrow marks a condi sing temperature tion, where the fiber applies a load on the surrounding The corresponding stress-strain curves are shown in Fig. matrix material, with subsequent crack formation at the 4. Comparing composite materials P5 and P7, the more center of the matrix area. Similar observations were also brittle behavior of composite P7 is evident made in the C/SiCN system, indicating that this phenomena The overall strength level, however, is low, compared to is of general meaning the strength of the C fibers. Evaluation of the fiber surface after annealing at 1400"C in N2, showed substantial kink 3.2. Coeficient of thermal expansion formation at the surface, thus degrading the strength. No further matrix cracking is observed during loading, since Further characterization of the cision material included
density change as a function of the number of infiltration and pyrolysis cycles. After three infiltration and pyrolysis cycles, a density of 1.4 g/cm3 is reached in the C/SiCN system with an open porosity of 28%, which is further reduced to 16% and 12% after five and seven cycles, respectively. The final geometrical density is approximately 1.7 g/cm3 , bulk density 1.9 g/cm3 . A closed porosity of about 10% cannot be eliminated and is present already after three cycles. Fig. 2 shows the comparison of the microstruture of the C/SiCN and Hi-Nicalon/SiCN composites. Microstructural evolution and conclusions regarding the stress–strain behavior of the C/SiCN composites were monitored on fracture surfaces of tested composite materials P3–P7 (Fig. 3). As is evident, with an increasing number of infiltration cycles, a more continuous matrix develops, but also the fracture characteristics are changing. Fig. 3(b) shows the microstructure of a composite after five cycles, clearly demonstrating pull-out behavior with pull-out lengths .100 mm. However, with increasing matrix density, the fracture behavior becomes more brittle, as is seen in Fig. 3(c). This can be attributed to the development of local residual compressive stresses at the fiber interface, caused by the large shrinkage (30% linear) of the polymer upon pyrolysis. In addition, stresses caused by the thermal expansion mismatch between fiber and matrix can develop, although, as calculations have shown, the interface should be in tension in this system upon cooling from the processing temperature. The corresponding stress–strain curves are shown in Fig. 4. Comparing composite materials P5 and P7, the more brittle behavior of composite P7 is evident. The overall strength level, however, is low, compared to the strength of the C fibers. Evaluation of the fiber surface, after annealing at 14008C in N2, showed substantial kink formation at the surface, thus degrading the strength. No further matrix cracking is observed during loading, since the matrix has already reached its saturation crack density upon cooling due to the thermal expansion mismatch of fiber and matrix (aF(axial) < 0, a M < 3.0 × 1026 K21 ). The residual stresses upon cooling due to thermal expansion mismatch between fiber and matrix were calculated using a simple model, as described elsewhere [8], assuming a stress-free state at the pyrolysis temperature. The calculated stresses axial and radial between fiber and matrix (C/ SiCN) were determined to be 780 and 90 MPa, respectively. Upon cooling, the fiber should therefore shrink away from the matrix, which was observed. The matrix is also under tension and with a matrix yield strength (as has been experimentally determined on monolithic material), an array of regular matrix cracks will develop, as has been confirmed by SEM-micrographs on polished cross-sections. Different results are obtained when working with Ccoated Hi-Nicalon fibers. Fig. 5 shows the stress–strain diagrams of these composites after five and seven infiltration cycles. The high strength values are obviously due to the undamaged SiC fibers and the coating. Polished cross-sections of a Hi-Nicalon/SiCN material were used to get some insight into the matrix formation (Fig. 6). Various points can be made. First, different infiltration cycles (marked as 1 and 2 in Fig. 6) can be distinguished, appearing as lighter and darker areas. Second, looking at the crack patterns in matrix areas adjacent to the fibers (1st cycle) it is obvious that the matrix may shrink onto the fibers (3) but also away from them (4), depending on the local geometric conditions. The arrow marks a condition, where the fiber applies a load on the surrounding matrix material, with subsequent crack formation at the center of the matrix area. Similar observations were also made in the C/SiCN system, indicating that this phenomena is of general meaning. 3.2. Coefficient of thermal expansion Further characterization of the C/SiCN material included G. Ziegler et al. / Composites: Part A 30 (1999) 411–417 413 Fig. 2. Comparison of the microstructure of the (a) C/SiCN and (b) Hi-Nicalon/SiCN composites
G. Ziegler et al./Composites: Part A 30(1999)411-417 b) K Fig 3. Fracture surfaces of a C/SICN composite after(a)three, (b)five and (c) seven infiltration and pyrolysis cycles, and schematic model for the formation of esidual stresses the determination of coefficients of thermal expansion adjusted by varying the fiber orientation. The changes after (CTE)parallel and perpendicular to the fiber axis of UD- fferent infiltration and pyrolysis cycles are small laminates(Fig. 7). As expected, the thermal expansion of In corporation with the TU Munich, which performed unidirectional reinforced C/SiCN composites perpendicular FEM calculations with respect to an optimal prepreg config to the fiber orientation is higher than parallel. This is due to uration in order to minimize the 2D anisotropy, C fiber the anisotropy of the Cte of the carbon fiber which(accord- laminates with varying configurations(0%, 30., 45, etc. ing to the manufacturers data) varies by two orders of were prepared ( Table I magnitude perpendicular versus parallel to the fiber axis As is shown in Fig 8(a), isotropic properties in the x-and C fibers embedded in a SiCN matrix, however, exhibit a y-direction were achieved already for a simple 0/90 CTE-ratio of only 2-4(perpendicular/parallel). The prepreg(configuration 0) with eight symmetrical arranged measured average CTE values of C/SiCn (20-1000C) layers. Fig. 8(b) shows configuration 6(0/90/45), with are between1.5×10-6K-l( parallel) and45×10-6K basically the same result (-direction is perpendicular to (perpendicular). Between these two bounds, the Cte can be the x-y-area)
the determination of coefficients of thermal expansion (CTE) parallel and perpendicular to the fiber axis of UDlaminates (Fig. 7). As expected, the thermal expansion of unidirectional reinforced C/SiCN composites perpendicular to the fiber orientation is higher than parallel. This is due to the anisotropy of the CTE of the carbon fiber which (according to the manufacturer’s data) varies by two orders of magnitude perpendicular versus parallel to the fiber axis. C fibers embedded in a SiCN matrix, however, exhibit a CTE-ratio of only 2–4 (perpendicular/parallel). The measured average CTE values of C/SiCN (20–10008C) are between 1.5 × 1026 K21 (parallel) and 4.5 × 1026 K21 (perpendicular). Between these two bounds, the CTE can be adjusted by varying the fiber orientation. The changes after different infiltration and pyrolysis cycles are small. In corporation with the TU Munich, which performed FEM calculations with respect to an optimal prepreg configuration in order to minimize the 2D anisotropy, C fiber laminates with varying configurations (08, 308, 458, etc.) were prepared (Table 1). As is shown in Fig. 8(a), isotropic properties in the x- and y-direction were achieved already for a simple 08/908- prepreg (configuration 0) with eight symmetrical arranged layers. Fig. 8(b) shows configuration 6 (08/908/458), with basically the same result (z-direction is perpendicular to the x–y-area). 414 G. Ziegler et al. / Composites: Part A 30 (1999) 411–417 Fig. 3. Fracture surfaces of a C/SiCN composite after (a) three, (b) five and (c) seven infiltration and pyrolysis cycles, and schematic model for the formation of residual stresses
G. Ziegler et al./Composites: Part 4 30(1999)411-417 Fig. 6. Polished cross-section of a Hi-Nicalon/SiCN composite, P5, Fig. 4. Stress-strain diagram of C/SiCN composites 000°C P Hi-Nicalon/SiCN P5, perpendicular P7, perpendicular parallel P5, parallel Fig. 7. Relative dimensional changes of a UD-C/SiCN composite parallel and perpendicular to the fiber orientation after various infiltration cycles Fig. 5. Stress-strain diagram of Hi-Nicalon/SICN composites. the individual prepreg-layers), the fibers are partially still present, showing sufficient pull-out behavior. Fig. 12 33. Oxidation behavior and siliconization demonstrates the enhanced oxidation resistance of the For further characterization, C/SICN composites as well as the individual components(matrix and fiber)[material 4 Conclusions P5] were tested with respect to oxidation stability by oxida- tion up to 1600C in flowing air(Fig 9). The matrix itself is It has been shown that C and SiC fiber-reinforced compo- very stable at least up to 1400.C, whereas the fiber is sites can be successfully prepared using the liquid infiltra- oxidized as low as 600C Fig. 10 shows a micrograph of tion/pyrolysis method. Prerequisites are suitable precursor the oxidized composite. Therefore, this matrix in combina- catalyst systems with respect to setting temperature, viScos- tion with an oxidation-stable(eventually oxidic)fiber may ity and ceramic yield. An important aspect is the matrix constitute an excellent oxidation-resistant material development, affecting the mechanical properties by indu- In order to improve the oxidation and wear resistance, cing residual stresses due to shrinkage of the matrix upon siliconization experiments were carried out with the C/ pyrolysis, and due to the marked anisotropy of the CTEs of SICN composites. Experiments with varying fiber contents fiber and matrix. Furthermore, the C fibers used are showed clearly that closed packed and regular arranged fib damaged, also high modulus/high tenacity fibers(which bundles within a Sicn matrix remain unchanged, whereas will be used in the future)should be less susceptible to randomly arranged fibers within such a matrix will be degradation Composites with favorable mechanical proper converted to SiC. It is therefore possible to prepare C/ ties were prepared with Hi-Nicalon reinforcements, but SiCN composites with siliconized areas by careful arrange should also be preparable from suitable and coated C fibe ment of the c fibers reinforcements. Coefficients of thermal expansion of Fig. 1l shows a fracture surface of such a siliconized samples prepared from UD-prepregs were found to be material, indicating that in closed packed regions(within anisotropic, but can be adjusted by varying the stacking
3.3. Oxidation behavior and siliconization For further characterization, C/SiCN composites as well as the individual components (matrix and fiber) [material P5] were tested with respect to oxidation stability by oxidation up to 16008C in flowing air (Fig. 9). The matrix itself is very stable at least up to 14008C, whereas the fiber is oxidized as low as 6008C. Fig. 10 shows a micrograph of the oxidized composite. Therefore, this matrix in combination with an oxidation-stable (eventually oxidic) fiber may constitute an excellent oxidation-resistant material. In order to improve the oxidation and wear resistance, siliconization experiments were carried out with the C/ SiCN composites. Experiments with varying fiber contents showed clearly that closed packed and regular arranged fiber bundles within a SiCN matrix remain unchanged, whereas randomly arranged fibers within such a matrix will be converted to SiC. It is therefore possible to prepare C/ SiCN composites with siliconized areas by careful arrangement of the C fibers. Fig. 11 shows a fracture surface of such a siliconized material, indicating that in closed packed regions (within the individual prepreg-layers), the fibers are partially still present, showing sufficient pull-out behavior. Fig. 12 demonstrates the enhanced oxidation resistance of the material. 4. Conclusions It has been shown that C and SiC fiber-reinforced composites can be successfully prepared using the liquid infiltration/pyrolysis method. Prerequisites are suitable precursor/ catalyst systems with respect to setting temperature, viscosity and ceramic yield. An important aspect is the matrix development, affecting the mechanical properties by inducing residual stresses due to shrinkage of the matrix upon pyrolysis, and due to the marked anisotropy of the CTEs of fiber and matrix. Furthermore, the C fibers used are damaged, also high modulus/high tenacity fibers (which will be used in the future) should be less susceptible to degradation. Composites with favorable mechanical properties were prepared with Hi-Nicalon reinforcements, but should also be preparable from suitable and coated C fiber reinforcements. Coefficients of thermal expansion of samples prepared from UD-prepregs were found to be anisotropic, but can be adjusted by varying the stacking G. Ziegler et al. / Composites: Part A 30 (1999) 411–417 415 Fig. 5. Stress–strain diagram of Hi-Nicalon/SiCN composites. Fig. 4. Stress–strain diagram of C/SiCN composites. Fig. 6. Polished cross-section of a Hi-Nicalon/SiCN composite, P5, 10008C. Fig. 7. Relative dimensional changes of a UD-C/SiCN composite parallel and perpendicular to the fiber orientation after various infiltration cycles
G. Ziegler et al./Composites: Part A 30(1999)411-417 z-direction 00 800 Temperature°Cl erature Fig 8. Thermal expansion of 2D C/SiCN composites as a function of fiber orientation(materials P5).(a)Configuration 0.(b)Configuration 6(see Table 1) Tested laminate configurations with respect to minimized x-y anisotropy of Lamination. Orientation of the individual Config. o 0°0°90°0°/90°10°190° 5-1550930456 Confi 30°0°90°60°/30°0°/30°60% Config. 3 -80 Config. 4 0°30°60°/90°/1201150°0°30 100 60°90°120%/150° Config. 5 Config. 6 Fig 9. Oxidation behavior of the C/SICN composite and the individual 5°0°90°145%-45° Config. 7 0°45°/90°-45°0°45°90一 45°0°14591907-45° Config. 8 0-90°459-45°90°0° Config. 9 0-90°145-45°/90°10° 0°0°°0°0°0°0°0°90°90° 90°90°90°/90°/90°90° 015°30°/4560°75°90°/105 8 120%/135/150/165%180°/195° 210°/225° quence, as has been shown on 0/90 and other prepregs Siliconization of the C/SiCN composites for improved prop- the s possible without damaging the fibers depending on Acknowledgements Dipl. Ing. B. Rothe is gratefully acknowledged for providing the optical microscopy work Fig. 10. Oxidized C/SiCN composite
sequence, as has been shown on 08/908 and other prepregs. Siliconization of the C/SiCN composites for improved properties is possible without damaging the fibers, depending on the fiber arrangement. Acknowledgements Dipl. Ing. B. Rothe is gratefully acknowledged for providing the optical microscopy work. 416 G. Ziegler et al. / Composites: Part A 30 (1999) 411–417 Fig. 9. Oxidation behavior of the C/SiCN composite and the individual components. Table 1 Tested laminate configurations with respect to minimized x–y anisotropy of CTEs Laminatconf. Orientation of the individual laminates Config. 0 08/908/08/908/08/908/08/908 Config. 1 08/158/ 2 158/308/ 2 308/458/ 2 458/608/ 2 608/758/ 2 758/908 Config. 2 308/08/908/608/308/08/308/608/ 908/08/308 Config. 3 08/308/608/908/08/308/608/908/08/ 308/608/908 Config. 4 08/308/608/908/1208/1508/08/308/ 608/908/1208/1508 Config. 5 08/08/08/308/308/308/608/608/608/ 908/908/908 Config. 6 08/908/458/ 2 458/08/908/458/ 2 458/08/908/458/ 2 458 Config. 7 08/458/908/ 2 458/08/458/908/ 2 458/08/458/908/ 2 458 Config. 8 08/ 2 908/458/ 2 458/908/08/ 1358/458/08/ 2 908/458/ 2 458/ 908/08/1358/458 Config. 9 08/ 2 908/458/ 2 458/908/08/ 1358/458/458/1358/08/908/ 2 458/ 458/ 2 908/08 Config. 10 08/08/08/08/08/08/08/08/908/908/ 908/908/908/908/908/908 Config. 11 08/158/308/458/608/758/908/1058/ 1208/1358/1508/1658/1808/1958/ 2108/2258 Fig. 8. Thermal expansion of 2D C/SiCN composites as a function of fiber orientation (materials P5). (a) Configuration 0. (b) Configuration 6 (see Table 1). Fig. 10. Oxidized C/SiCN composite
G. Ziegler et al./Composites: Part 4 30(1999)411-417 Fig. 11. Siliconized C/SICN composite, showing partially integer fibers [2] Guo JK, Mao ZQ, Bao CD, Wang RH, Yan DS. Carbon fiber-rein- orced silicon nitride composite J Mater Sci 1982; 17: 3611 3] Lackey WJ. Review, status and future of the chemical vapor infiltration process for fabrication of fiber-reinforced ceramic composites. Ceram Engng Sci Proc 1989: 10(7-8):577 [4] Haug T, Ostertag R, Knabe H, Ehrmann U, Woltersdorf J. Processing nd mechanical properties of CMC 's by the infiltration and pyrolysis of Si-polymers. In: Naslain R, Lamon J, Doumeingts D, editors. High Temp. Ceram. Mat. Comp. Bordeaux: Woodhead, 1993: 767. pyrolysis siliconized [5] Ziegler G, Hapke J, Lucke J. Processing of CMCs from novel organo- netallic precursors. In: Evans AG, Naslain R, editors. Proc. 2nd Intern. Conf High Temp. Ceram Matrix Compos. HT-CMC-ll St Barbara, 4008001200 [6 Suttor D, Hacker J, TraBl S, Muller H, Kleebe H-I Temperature [c] hitecture and crystallization behavior of s Ceram Engng Sci Proc 1997; 18(3): 127-131 Fig. 12. Oxidation behavior of the siliconized composites(compare Fig. 9) I Lucke J, Hacker J, Suttor D, Ziegler G. Synthesis and characte silazane-based polymers as precursors for ceramic matrix sites. Appl Organomet Chem 1996; 11: 181 References [8 Kuntz M, Meier B, Grathwohl G. Residual stresses in fiber-reinforced ceramics due to thermal expansion mismatch. J Am Ceram Sc [1] Evans AG, Zok Fw. Review. The physics and mechanics of fiber- reinforced brittle matrix composites. J Mater Sci 1994, 29:3896
References [1] Evans AG, Zok FW. Review: The physics and mechanics of fiberreinforced brittle matrix composites. J Mater Sci 1994;29:3896. [2] Guo JK, Mao ZQ, Bao CD, Wang RH, Yan DS. Carbon fiber-reinforced silicon nitride composite. J Mater Sci 1982;17:3611. [3] Lackey WJ. Review, status and future of the chemical vapor infiltration process for fabrication of fiber-reinforced ceramic composites. Ceram Engng Sci Proc 1989;10(7–8):577. [4] Haug T, Ostertag R, Knabe H, Ehrmann U, Woltersdorf J. Processing and mechanical properties of CMC’s by the infiltration and pyrolysis of Si-polymers. In: Naslain R, Lamon J, Doumeingts D, editors. High. Temp. Ceram. Mat. Comp. Bordeaux: Woodhead, 1993:767. [5] Ziegler G, Hapke J, Lu¨cke J. Processing of CMCs from novel organometallic precursors. In: Evans AG, Naslain R, editors. Proc. 2nd Intern. Conf. High Temp. Ceram. Matrix Compos. HT-CMC-II. St. Barbara, Vol. II 58, 1995:13. [6] Suttor D, Hacker J, Traßl S, Mu¨ller H, Kleebe H-J, Ziegler G. Polymer architecture and crystallization behavior of SiCN-fiber precursors. Ceram Engng Sci Proc 1997;18(3):127–131. [7] Lu¨cke J, Hacker J, Suttor D, Ziegler G. Synthesis and characterisation of silazane-based polymers as precursors for ceramic matrix composites. Appl Organomet Chem 1996;11:181. [8] Kuntz M, Meier B, Grathwohl G. Residual stresses in fiber-reinforced ceramics due to thermal expansion mismatch. J Am Ceram Soc 1993;76(10):2607. G. Ziegler et al. / Composites: Part A 30 (1999) 411–417 417 Fig. 12. Oxidation behavior of the siliconized composites (compare Fig. 9). Fig. 11. Siliconized C/SiCN composite, showing partially integer fibers