MATERIALS HIENGE& ENGIEERING ELSEVIER Materials Science and Engineering A 437(2006)268-273 www.elsevier.com/locate/msea Formation of carbon fiber-reinforced ceramic matrix composites with polysiloxane/silicon derived matrix R. M. Rocha", C.A. A. Cairo, M.L.A. Graca Aerospace Technical Center/Aeronautic and Space Institute, CTA, IAE, Pca. Marechal do Ar Eduardo Gomes, 50-S. Jose dos Campos 12228-904, SP, Brazil Received 30 January 2006: received in revised form 1I July 2006; accepted 30 August 2006 A ceramic matrix for carbon fiber-reinforced ceramic matrix composites(CMCs) has been developed from poly(methylsilsesquioxane)/silicon mixtures, using a low-cost process. In this process the space in two-dimensional carbon fiber preform was filled with a slurry composed by si powder dispersed into poly(methylsilsesquioxane )trietoxysilane solutions. Three different volume ratio of Si polymer were used to stack eigh arness plain weave of carbon fiber, forming laminates composites, which were pressed and cured up to 200C. The compact bodies were first pre-pyrolyzed at 1000C and then pyrolyzed at 1450Ch hand 1500C/1 h On pyrolysis, the polymer- filler mixture was converted to a multiphase ceramic matrix through reactions between Si, gaseous and solids products from the polymer degradation and the N2 atmosphere. Pyrolysis led to conversion of the initial matrix into silicon oxide( SiO2), silicon carbide(SiC) and silicon oxinitride(Si2ON2), though after pyrolysis at 1450"C metallic silicon was still detected. With one cycle of infiltration the composite characteristics were followed by bulk density and open porosity measurements, X-ray diffraction, microscopy and mechanical testing o 2006 Elsevier B.V. All rights reserved Keywords: Silicon oxycarbide; Polysiloxane; Polymer/filler derived ceramics; Carbon fiber; Ceramic matrix composites 1. Introduction between the matrix and the fiber are necessary to control the Ceramic matrix composites(CMC,), especially carbon fiber. the precursor rpolvmser filtration and pyrolysis (Pipi which is reinforced ceramic matrix composites(CFCMCs), are widely being actively developed because it offers many advantages such used in different areas of engineering due to their unique com- as low processing temperature, controllable ceramic composi bination of properties, in particular high mechanical strength tions, availability of complex reinforcements(three dimensional together with low density. That is why they are so used in the textiles)and near-net shape technologies [3, 7]. This is simi aeronautical industry and in the aerospace applications where lar in principle to the formation of the carbon matrix in the most of the ceramic and metallic materials cannot be used [l- preparation of C-C composites, using a preceramic polymer Therefore, CMCs are ideal candidates for demanding low-and instead of a phenolic resin. Polysiloxane(PSO), polycarbosi- high temperatures environments like those in brake disks, ther- lane(PCS), and polysilazane(PSz), which give Si-O-C, Si-C mal protective plates, heat exchanges, gas turbine and chemical and Si-C-N ceramics upon pyrolysis, respectively, are the most [1-4]. often used preceramic polymers at present [7-1l]. The PIP pro- Eforts to minimize CFCMCs processing times and costs cess includes basically three stages. In stage 1, fiber preform have led to the development of a wide range of fabrication is infiltrated with precursor solution. In stage 2, the precursor methods, which must fill some requirements: the fabrication is cured into a solid. In stage 3, the cured solid is pyrolyzed echnique should allow complex geometries without damage to in inert or reactive environment. A major drawback is the large the fiber; the composite must be protected against oxidation, shrinkage of up to 60 vol. and low ceramic yield of usually because the carbon fiber will react with oxygen at temperatures 70 wt. %. Several infiltration-cure-pyrolysis cycles are required of as low as 500 C; thermochemical and physical compatibility to achieve dense matrices what make the process both expensive and time-consuming Corresponding author. Tel: +55 12 3947 6441: fax: +55 12 3947 6405. Anew approach to overcome this latter problem is the impreg- E-mail address: rosarocha @iae. cta. br(R. M. Rocha) nation of the fiber with mixtures of organometallic polymers 0921-5093 )6 Elsevier B v. All rights reserved doi:10.1016/msea.200608.1
Materials Science and Engineering A 437 (2006) 268–273 Formation of carbon fiber-reinforced ceramic matrix composites with polysiloxane/silicon derived matrix R.M. Rocha ∗, C.A.A. Cairo, M.L.A. Grac¸a Aerospace Technical Center/Aeronautic and Space Institute, CTA, IAE, P¸ca. Marechal do Ar Eduardo Gomes, 50-S. Jos´e dos Campos 12228-904, SP, Brazil Received 30 January 2006; received in revised form 11 July 2006; accepted 30 August 2006 Abstract A ceramic matrix for carbon fiber-reinforced ceramic matrix composites (CMCs) has been developed from poly(methylsilsesquioxane)/silicon mixtures, using a low-cost process. In this process the space in two-dimensional carbon fiber preform was filled with a slurry composed by Si powder dispersed into poly(methylsilsesquioxane)/trietoxysilane solutions. Three different volume ratio of Si:polymer were used to stack eightharness plain weave of carbon fiber, forming laminates composites, which were pressed and cured up to 200 ◦C. The compact bodies were first pre-pyrolyzed at 1000 ◦C and then pyrolyzed at 1450 ◦C/2 h and 1500 ◦C/1 h. On pyrolysis, the polymer-filler mixture was converted to a multiphase ceramic matrix through reactions between Si, gaseous and solids products from the polymer degradation and the N2 atmosphere. Pyrolysis led to conversion of the initial matrix into silicon oxide (SiO2), silicon carbide (SiC) and silicon oxinitride (Si2ON2), though after pyrolysis at 1450 ◦C metallic silicon was still detected. With one cycle of infiltration the composite characteristics were followed by bulk density and open porosity measurements, X-ray diffraction, microscopy and mechanical testing. © 2006 Elsevier B.V. All rights reserved. Keywords: Silicon oxycarbide; Polysiloxane; Polymer/filler derived ceramics; Carbon fiber; Ceramic matrix composites 1. Introduction Ceramic matrix composites (CMCs), especially carbon fiberreinforced ceramic matrix composites (CFCMCs), are widely used in different areas of engineering due to their unique combination of properties, in particular high mechanical strength together with low density. That is why they are so used in the aeronautical industry and in the aerospace applications where most of the ceramic and metallic materials cannot be used [1]. Therefore, CMCs are ideal candidates for demanding low-and high temperatures environments like those in brake disks, thermal protective plates, heat exchanges, gas turbine and chemical reactors [1–4]. Efforts to minimize CFCMCs processing times and costs have led to the development of a wide range of fabrication methods, which must fill some requirements: the fabrication technique should allow complex geometries without damage to the fiber; the composite must be protected against oxidation, because the carbon fiber will react with oxygen at temperatures of as low as 500 ◦C; thermochemical and physical compatibility ∗ Corresponding author. Tel.: +55 12 3947 6441; fax: +55 12 3947 6405. E-mail address: rosarocha@iae.cta.br (R.M. Rocha). between the matrix and the fiber are necessary to control the interface characteristics [5,6]. A potential fabrication route is the precursor polymer infiltration and pyrolysis (PIP), which is being actively developed because it offers many advantages such as low processing temperature, controllable ceramic compositions, availability of complex reinforcements (three dimensional textiles) and near-net shape technologies [3,7]. This is similar in principle to the formation of the carbon matrix in the preparation of C-C composites, using a preceramic polymer instead of a phenolic resin. Polysiloxane (PSO), polycarbosilane (PCS), and polysilazane (PSZ), which give Si–O–C, Si–C, and Si–C–N ceramics upon pyrolysis, respectively, are the most often used preceramic polymers at present [7–11]. The PIP process includes basically three stages. In stage 1, fiber preform is infiltrated with precursor solution. In stage 2, the precursor is cured into a solid. In stage 3, the cured solid is pyrolyzed in inert or reactive environment. A major drawback is the large shrinkage of up to 60 vol.% and low ceramic yield of usually 70 wt.%. Several infiltration-cure-pyrolysis cycles are required to achieve dense matrices what make the process both expensive and time-consuming. A new approach to overcome this latter problem is the impregnation of the fiber with mixtures of organometallic polymers 0921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2006.08.102
R.M. Rocha et aL. Materials Science and Engineering A 437(2006)268-273 and reactive filler to limit excessive shrinkage of the ceramic Table I matrix[12-14]. According to the so called AFCOP (active filler Slurry characteristics of the three compositions prepared to form the controlled pyrolysis)process, the polymer is partially filled with filler/polymer derived ceramic matrix composite inert or active powder, in order to decrease the shrinkage. A num- Composite Slurry characteristics ber of systems have already been investigated using polysiloxane Volume ratio resins as preceramic polymers filled with different oxide, car- (MTES: PMS) (Si: polymer) bide, or metallic powders. When pyrolyzed, the filler will react with the decomposition products of the polymer and the atmo.I 0.65:1 0.25:1 1.33 sphere to form a ceramic matrix, consisting of the corresponding 0.35:1 0.33:1 metal filler carbides, nitrides, and a polymer-derived Si-O-C e The aim of the present work was to exploit this par- pressure was heated to 60, 90, 130 and 200C with a dwell ular processing method for the preparation of contin- time of 15 min in each temperature. It was applied vacuum in ous carbon fiber-reinforced ceramic composites. Si-filled the die to eliminate de condensation products during polymer poly(menthylsilsesquioxane)mixtures were used to infiltrate cross-linking ment in N2 atmosphere up to 1500C finally resulted in the 10 mm x 4 mm and pyrolyzed. The pyrolysis of the coma k two-dimensional carbon fiber via impregnation route. Heat treat- The composites were cut in small pieces of 50mm formation of a Cr/ceramic matrix composite. The effects of the ite samples was performed in two steps. First one was done slurry composition and the pyrolysis temperature on the poros- in a resistance-heated laboratory furnace at 1000"C for 30 min ity and microstructure of the composites were investigated over under flowing N2, with heating rate of 60 C/h. The second step one cycle of impregnation was performed in a tubular graphite furnace at temperatures of 1450C/2 h and 1500C/lh, both in flowing N2. The heating 2. Experimental rate was 10C/min until 1000C, 5C/min to the final temper ature and then free cooling The approach to composite fabrication in this work was Weight loss upon pyrolysis up to 1100C was character basically the preparation of a laminate composite with two- ized by thermogravimetry, heating composite specimens in a dimensional carbon fiber plain weave, arranged in 8 layers. Each thermal balance(Perkin-Elmer, TGA-7)in flowing N2 atmo- layer was put upon the previous in the same orientation and man- sphere with a constant heating rate of 10C/min. Fiber-free ually impregnated with a poly(methylsilsesquioxane)(PMS), matrix of the composite m was also examined by thermo- previously mixed with Si powder. gravimetry, using the same parameters as used for the com- The woven of carbon used as reinforcement material was posites. The bulk densities of the pyrolyzed samples were twill T-10 EKHO A, manufactured by Ural-Ukraine, derived measured according to Archimede's principle with distilled of the polyacrilonitrile precursor(PAN), carbonized, grama- water as the immersion medium; the phases were investi- ture 340 g/m with a bunch of 3000 filaments and specific gated by X-ray diffraction (XRD, Phillips Pw 18/30)using mass of 1. 48 g/cm. The slurry was prepared with the precur- Cu Ka X-radiation(1.5406 A); the microstructure investiga- sor polymer and the filler. Si powder was used as an active tion of the matrix, fibers and the interface between them was filler(Hermann C Starck, Berlin)with an average particle size conducted by scanning electron microscopy(SEM, LEO-435Pi of 10 um. The polymer phase was composed by a solid pre- and Zeiss-DSM). The room-temperature flexural strength of the cursor and a silane in which the solid precursor was diluted: composite samples was measured on bars with dimensions of poly ( methylsilsesquioxane)-PMS(MK-Wacker Chemie, Ger- 45 mm x 4.5mm x 3.0 mm in four-point bending tests(4-PBT) many), general formula of [CH3 SiO1.5]n with n=130-150 and span 40/20 mm, and a crosshead speed of 0.5 mm/min, using a a melting point of 42C; methyltrietoxysilane-MTES, general universal testing machine(Instron, model 4301). The load ver- formula CH3 Si(OC3 H5)3. Cross-linking catalysts(aluminun sus displacement curves were recorded and the maximum of the acetylacetonate)and the Si powder were added stepwise and curve gives the flexural strength. After the flexure tests, SEM homogenized through vigorous stirring for 20 min. Three was employed to observe the fracture surfaces of the compos ferent slurries composition were prepared and they are shown in Ites. Table 1. The volume ratio of mtes: PMS was decreased from composites I to Ill, whereas the volume ratio of the Si polymer 3. Results and discussion and the specific weight of the slurries were increased A small quantity of the slurry was homogeneously distributed The thermogravimetry results of the cured composites are was layered and infiltrated with the ceramic slurry, and so on. ture to 1 100C the weight change is similar for all f empera Then, the stacked composites were molded in autoclave, with It is clear that the composites show a major weight loss at the application of pressure of 0.3 MPa and left at room temperature beginning of analysis at temperatures below 200 C From 200 to for 15 h. Cross-linking of the laminates had to be carried out 600C, there is a smooth weight loss. A second distinguishable stepwise to avoid the formation of cracks caused by the products weight loss occurs at temperatures between 600 and 800C, due of the polymer condensation cure. The laminate under uniaxial to the decomposition of PMS. Above 800C, the weight of the
R.M. Rocha et al. / Materials Science and Engineering A 437 (2006) 268–273 269 and reactive filler to limit excessive shrinkage of the ceramic matrix [12–14]. According to the so called AFCOP (active filler controlled pyrolysis) process, the polymer is partially filled with inert or active powder, in order to decrease the shrinkage. A number of systems have already been investigated using polysiloxane resins as preceramic polymers filled with different oxide, carbide, or metallic powders. When pyrolyzed, the filler will react with the decomposition products of the polymer and the atmosphere to form a ceramic matrix, consisting of the corresponding metal filler carbides, nitrides, and a polymer-derived Si–O–C glass. The aim of the present work was to exploit this particular processing method for the preparation of continuous carbon fiber-reinforced ceramic composites. Si-filled poly(menthylsilsesquioxane) mixtures were used to infiltrate two-dimensional carbon fiber via impregnation route. Heat treatment in N2 atmosphere up to 1500 ◦C finally resulted in the formation of a Cf/ceramic matrix composite. The effects of the slurry composition and the pyrolysis temperature on the porosity and microstructure of the composites were investigated over one cycle of impregnation. 2. Experimental The approach to composite fabrication in this work was basically the preparation of a laminate composite with twodimensional carbon fiber plain weave, arranged in 8 layers. Each layer was put upon the previous in the same orientation and manually impregnated with a poly(methylsilsesquioxane) (PMS), previously mixed with Si powder. The woven of carbon used as reinforcement material was twill T-10 EKHO A, manufactured by Ural-Ukraine, derived of the polyacrilonitrile precursor (PAN), carbonized, gramature 340 g/m2, with a bunch of 3000 filaments and specific mass of 1.48 g/cm3. The slurry was prepared with the precursor polymer and the filler. Si powder was used as an active filler (Hermann C. Starck, Berlin) with an average particle size of 10m. The polymer phase was composed by a solid precursor and a silane in which the solid precursor was diluted: poly(methylsilsesquioxane)-PMS (MK-Wacker Chemie, Germany), general formula of [CH3SiO1.5]n with n = 130–150 and a melting point of 42 ◦C; methyltrietoxysilane-MTES, general formula CH3Si(OC3H5)3. Cross-linking catalysts (aluminun acetylacetonate) and the Si powder were added stepwise and homogenized through vigorous stirring for 20 min. Three different slurries composition were prepared and they are shown in Table 1. The volume ratio of MTES:PMS was decreased from composites I to III, whereas the volume ratio of the Si:polymer and the specific weight of the slurries were increased. A small quantity of the slurry was homogeneously distributed on the surface of carbon ply of fabric; the next ply of fabric was layered and infiltrated with the ceramic slurry, and so on. Then, the stacked composites were molded in autoclave, with application of pressure of 0.3 MPa and left at room temperature for 15 h. Cross-linking of the laminates had to be carried out stepwise to avoid the formation of cracks caused by the products of the polymer condensation cure. The laminate under uniaxial Table 1 Slurry characteristics of the three compositions prepared to form the filler/polymer derived ceramic matrix composite Composite Slurry characteristics Volume ratio (MTES:PMS) Volume ratio (Si:polymer) Specific weight (g cm−3) I 0.65:1 0.1:1 1.18 II 0.6:1 0.25:1 1.33 III 0.35:1 0.33:1 1.42 pressure was heated to 60, 90, 130 and 200 ◦C with a dwell time of 15 min in each temperature. It was applied vacuum in the die to eliminate de condensation products during polymer cross-linking. The composites were cut in small pieces of 50 mm × 10 mm × 4 mm and pyrolyzed. The pyrolysis of the composite samples was performed in two steps. First one was done in a resistance-heated laboratory furnace at 1000 ◦C for 30 min under flowing N2, with heating rate of 60 ◦C/h. The second step was performed in a tubular graphite furnace at temperatures of 1450 ◦C/2 h and 1500 ◦C/1 h, both in flowing N2. The heating rate was 10 ◦C/min until 1000 ◦C, 5 ◦C/min to the final temperature and then free cooling. Weight loss upon pyrolysis up to 1100 ◦C was characterized by thermogravimetry, heating composite specimens in a thermal balance (Perkin-Elmer, TGA-7) in flowing N2 atmosphere with a constant heating rate of 10 ◦C/min. Fiber-free matrix of the composite III was also examined by thermogravimetry, using the same parameters as used for the composites. The bulk densities of the pyrolyzed samples were measured according to Archimede’s principle with distilled water as the immersion medium; the phases were investigated by X-ray diffraction (XRD, Phillips PW 18/30) using Cu K X-radiation (1.5406 A); the microstructure investiga- ˚ tion of the matrix, fibers and the interface between them was conducted by scanning electron microscopy (SEM, LEO-435 Pi and Zeiss-DSM). The room-temperature flexural strength of the composite samples was measured on bars with dimensions of 45 mm × 4.5 mm × 3.0 mm in four-point bending tests (4-PBT), span 40/20 mm, and a crosshead speed of 0.5 mm/min, using a universal testing machine (Instron, model 4301). The load versus displacement curves were recorded and the maximum of the curve gives the flexural strength. After the flexure tests, SEM was employed to observe the fracture surfaces of the composites. 3. Results and discussion The thermogravimetry results of the cured composites are shown in Fig. 1. In the temperature range from room temperature to 1100 ◦C the weight change is similar for all composites. It is clear that the composites show a major weight loss at the beginning of analysis at temperatures below 200 ◦C. From 200 to 600 ◦C, there is a smooth weight loss. A second distinguishable weight loss occurs at temperatures between 600 and 800 ◦C, due to the decomposition of PMS. Above 800 ◦C, the weight of the
270 R.M. Rocha et al. Materials Science and Engineering A 437(2006)268-273 Mass change of the composites after pre-pyrolysis and pyrolysis at 1450C/2h andl500°/lh Composite Mass change (%) 1000°C 1500°C 8 (-)8.0±0.2 (一)1.1±0.5 10.5±0.5 (-)7.0±0.1 (+)13±0.4 (-)104±0.6 below 200C, is caused mainly by releasing of some carbon fiber organic compounds or some adsorbed humidity. The second one, between 600 and 800C, is related to the PMs decomposition The trend of the pyrolytic process of the PMS was similar to those reported by other researchers [15, 16]. It was reported that dominating degradation reaction of polysiloxane occurs with the Temperature(C) cleavage of Si-(CH3) bonds between 600 and 800C. Methane and hydrogen are the predominant decomposition products. At Thermogravimetric analysis of cured composites when heated fromroom higher temperatures dehydrogenation still occurs resulting only ature to 1100 C under flowing nitrogen(heating rate: 10C/min) in a minor weight reduction up to 1200C[15] Total weight change was determined by taking the compos omposites shows almost no change, indicating that PMS has ite weight before and after pre-pyrolysis and pyrolysis. Table 2 been converted to inorganic materials. The composite I shows presents the final weight change results. After pre-pyrolysis the highest total weight loss, about 14%o, followed by compos-(1000.C)the determined results are lower than tho ose obtained ites II and Ill with approximately 13 and 11%, respectively at the final thermogravimetry analysis, but following the same was expected to occur once the matrix of composite I contains trend. One reason for this difference can be the lower heating the highest polymer phase volume. Composites lI and ill show rate used during pre-pyrolysis. After pyrolysis at 1450oC/h almost the same weight loss up to 600C, in spite of the dif- the total weight change is very low and in the case of the com- ference in their matrix composition. It is observed that between posite Ill, with the highest amount of Si filler, the change is even 200 and 600 C the composite which has lower amount of positive This effect can be attributed to the reactions with the polymer phase shows a slightly higher weight loss than com- N2 atmosphere that take place during pyrolysis. Fig. 3 compares cure degree of the composite Il matrix due to its lower vol- 1450, 1500C and its fiber free matrix pyrolyzed at 1450C ume ratio of MTES: PMS. In order to analyze the weight loss After 1450/2 h there is the presence of some unreacted sili contribution of the matrix separately from the carbon fibers, con, crystalline phases formed from the polymer crystallization it was done thermogravimetry in the fiber free matrix of the or reactions between Si and gaseous or solids products of poly omposite Ill(Fig. 2). It can be noted that the curve of the mer degradation(SiO2 and B-SiC)and reactions with nitrogen matrix and the composite follow the same behavior. The first major weight loss, which is more pronounced in the composite, 米S2ON2 C. free matrix res 500c/h 0.6 1450°c/2h C free Matrox1450°c 506070 Temperature(C) Relative weight and rate of weight loss versus temperature up to 1000.C Fig. 3. X-ray diffraction pattens of the fiber free matrix of composite Ill after atmosphere of the composite Ill and its carbon fiber free matrix(heating pyrolysis at 1450 Cn h and the composite itself after pyrolysis at 1450C/ h
270 R.M. Rocha et al. / Materials Science and Engineering A 437 (2006) 268–273 Fig. 1. Thermogravimetric analysis of cured composites when heated from room temperature to 1100 ◦C under flowing nitrogen (heating rate: 10 ◦C/min). composites shows almost no change, indicating that PMS has been converted to inorganic materials. The composite I shows the highest total weight loss, about 14%, followed by composites II and III with approximately 13 and 11%, respectively. It was expected to occur once the matrix of composite I contains the highest polymer phase volume. Composites II and III show almost the same weight loss up to 600 ◦C, in spite of the difference in their matrix composition. It is observed that between 200 and 600 ◦C the composite III, which has lower amount of polymer phase shows a slightly higher weight loss than composite II. This effect can be associated with a possible lower cure degree of the composite III matrix due to its lower volume ratio of MTES:PMS. In order to analyze the weight loss contribution of the matrix separately from the carbon fibers, it was done thermogravimetry in the fiber free matrix of the composite III (Fig. 2). It can be noted that the curve of the matrix and the composite follow the same behavior. The first major weight loss, which is more pronounced in the composite, Fig. 2. Relative weight and rate of weight loss versus temperature up to 1000 ◦C in N2 atmosphere of the composite III and its carbon fiber free matrix (heating rate: 10 ◦C/min). Table 2 Mass change of the composites after pre-pyrolysis and pyrolysis at 1450 ◦C/2 h and 1500 ◦C/1 h Composite Mass change (%) 1000 ◦C 1450 ◦C 1500 ◦C I (−) 12.4 ± 0.3 (−) 0.6 ± 0.1 (−) 5.2 ± 0.7 II (−) 8.0 ± 0.2 (−) 1.1 ± 0.5 (−) 10.5 ± 0.5 III (−) 7.0 ± 0.1 (+) 1.3 ± 0.4 (−) 10.4 ± 0.6 below 200 ◦C, is caused mainly by releasing of some carbon fiber organic compounds or some adsorbed humidity. The second one, between 600 and 800 ◦C, is related to the PMS decomposition. The trend of the pyrolytic process of the PMS was similar to those reported by other researchers [15,16]. It was reported that dominating degradation reaction of polysiloxane occurs with the cleavage of Si–(CH3) bonds between 600 and 800 ◦C. Methane and hydrogen are the predominant decomposition products. At higher temperatures dehydrogenation still occurs resulting only in a minor weight reduction up to 1200 ◦C [15]. Total weight change was determined by taking the composite weight before and after pre-pyrolysis and pyrolysis. Table 2 presents the final weight change results. After pre-pyrolysis (1000 ◦C) the determined results are lower than those obtained at the final thermogravimetry analysis, but following the same trend. One reason for this difference can be the lower heating rate used during pre-pyrolysis. After pyrolysis at 1450 ◦C/2 h the total weight change is very low and in the case of the composite III, with the highest amount of Si filler, the change is even positive. This effect can be attributed to the reactions with the N2 atmosphere that take place during pyrolysis. Fig. 3 compares the X-ray diffraction patterns of the composite III pyrolyzed at 1450, 1500 ◦C and its fiber free matrix pyrolyzed at 1450 ◦C. After 1450 ◦C/2 h there is the presence of some unreacted silicon, crystalline phases formed from the polymer crystallization or reactions between Si and gaseous or solids products of polymer degradation (SiO2 and -SiC) and reactions with nitrogen Fig. 3. X-ray diffraction patterns of the fiber free matrix of composite III after pyrolysis at 1450 ◦C/2 h and the composite itself after pyrolysis at 1450 ◦C/2 h and 1500 ◦C/1 h
R.M. Rocha et aL. Materials Science and Engineering A 437(2006)268-273 Table 3 Bulk density and open porosity results of the composites at different temperatures Composite Bulk density (g cm 500°C 1450°C 1500°C 1.27±0.04 1.24±0.04 119±0.02 25士3 27±1 1.36±0.04 135±0.0 143±003 142±0.03 133±0.0 24士2 29士1 atmosphere(Si2ON2) The nitrogen atmosphere is supposed to react with the gaseous Sio or with the liquid Si, giving rise to i2ON2[16]. Based on these results, the low weight change after pyrolysis at 1450C is associated with the Si ON2 formation. For pyrolysis at 1500C the weight loss is more significant. Although reactions with N2 atmosphere is likely to continue taking place, the weight increase is overcame by the carbother- mal reduction of Sio, which increases at temperatures higher than 1400C and also owing to the loss of Si through gas for- mation[17]. The X-ray diffraction pattern of the composite l pyrolyzed at 1500oC shows only Sio2 and B-SiC as crystalline phases(Fig. 3). At this temperature Si filler is not identified ndicating completely reaction or evaporation by formation of gaseous species. Bulk density and open porosity of the composites are shown in Table 3. For each temperature the differences in the bulk density can be attributed to the different matrix composition Composite Ill, which has the highest amount of Si filler, shows the highest density in all temperatures. However, the open poros- y values are almost the same for each temperature regar Si: polymer ratio. After 1000C, composites exhibited about 30% of open porosity, decreasing to about 25 %o after 1450C/2h and increase to about 30% after 1500C/1 h he microstructure of the composites pyrolyzed at 1000C shows no unfilled spaces between fibers and matrix. Fig. 4 shows the microstructure of the composite Il after pre-pyrolysis The microstructure is composed by carbon fibers surrounded by PMS derived ceramic, indicating that inter-fiber regions were infiltrated during its manufacturing by just the polymer phase The lack of porous regions between fibers at this temperature proves that open porosity is mainly located in the matrix around the fiber bundles. Fig. 4b presents the micrograph of the matrix located near a fiber bundle after pre-pyrolysis. The observed grains is Si embedded in the Si-o-C ceramic yield as the XRD 10 um identified only these phases. After 1450%C, the interfiber phase shows some porosity with some whisker like structure inside Fig. 4. SEM micrographs of the composite II after pre-pyrolysis:(a)cross- the pores(Fig. 5a), resulting from gaseous phases process. This section of the C fiber involved by Si-o-C ceramic;(b)matrix around fiber porosity is likely to be related to the shrinkage of the polymer bundle after 1450Cnh phase, owing to the crystallization and with the mass loss by carbothermal reduction. However, the total open porosity at this Besides this, the formation of Si]ON2 is also responsible for temperature is lower than that found after pre-pyrolysis. The reducing porosity by the volume expansion of Si during nitri reduction of porous regions in the matrix situated in the inter- dation, compensating the polymer shrinkage. After 1500C the fiber bundles compensates the observed porosity between fibers porosity is increased. Fig. 6 shows the high level of porosity The fracture morphology of the matrix between fiber bundles is in the matrix between fiber bundles of the composite Ill. These shown in Fig. 5b. This type of fracture presents some typical results imply that at this temperature the gas phase plays a major features found in metallic material fracture, indicating that Si role for material transport as well as filler reaction during pyrol melted and filled some cavities, in this way reducing porosity. ysis
R.M. Rocha et al. / Materials Science and Engineering A 437 (2006) 268–273 271 Table 3 Bulk density and open porosity results of the composites at different temperatures Composite Bulk density (g cm−3) Open porosity (%) 1000 ◦C 1450 ◦C 1500 ◦C 1000 ◦C 1450 ◦C 1500 ◦C I 1.27 ± 0.04 1.24 ± 0.04 1.19 ± 0.02 31 ± 3 25 ± 3 27 ± 1 II 1.36 ± 0.04 1.35 ± 0.02 1.25 ± 0.02 30 ± 1 26 ± 2 30 ± 2 III 1.43 ± 0.03 1.42 ± 0.03 1.33 ± 0.02 29 ± 2 24 ± 2 29 ± 1 atmosphere (Si2ON2). The nitrogen atmosphere is supposed to react with the gaseous SiO or with the liquid Si, giving rise to Si2ON2 [16]. Based on these results, the low weight change after pyrolysis at 1450 ◦C is associated with the Si2ON2 formation. For pyrolysis at 1500 ◦C the weight loss is more significant. Although reactions with N2 atmosphere is likely to continue taking place, the weight increase is overcame by the carbothermal reduction of SiO2, which increases at temperatures higher than 1400 ◦C and also owing to the loss of Si through gas formation [17]. The X-ray diffraction pattern of the composite III pyrolyzed at 1500 ◦C shows only SiO2 and -SiC as crystalline phases (Fig. 3). At this temperature Si filler is not identified, indicating completely reaction or evaporation by formation of gaseous species. Bulk density and open porosity of the composites are shown in Table 3. For each temperature the differences in the bulk density can be attributed to the different matrix compositions. Composite III, which has the highest amount of Si filler, shows the highest density in all temperatures. However, the open porosity values are almost the same for each temperature regardless Si:polymer ratio. After 1000 ◦C, composites exhibited about 30% of open porosity, decreasing to about 25% after 1450 ◦C/2 h and increase to about 30% after 1500 ◦C/1 h. The microstructure of the composites pyrolyzed at 1000 ◦C shows no unfilled spaces between fibers and matrix. Fig. 4a shows the microstructure of the composite II after pre-pyrolysis. The microstructure is composed by carbon fibers surrounded by PMS derived ceramic, indicating that inter-fiber regions were infiltrated during its manufacturing by just the polymer phase. The lack of porous regions between fibers at this temperature proves that open porosity is mainly located in the matrix around the fiber bundles. Fig. 4b presents the micrograph of the matrix located near a fiber bundle after pre-pyrolysis. The observed grains is Si embedded in the Si–O–C ceramic yield as the XRD identified only these phases. After 1450 ◦C, the interfiber phase shows some porosity with some whisker like structure inside the pores (Fig. 5a), resulting from gaseous phases process. This porosity is likely to be related to the shrinkage of the polymer phase, owing to the crystallization and with the mass loss by carbothermal reduction. However, the total open porosity at this temperature is lower than that found after pre-pyrolysis. The reduction of porous regions in the matrix situated in the inter- fiber bundles compensates the observed porosity between fibers. The fracture morphology of the matrix between fiber bundles is shown in Fig. 5b. This type of fracture presents some typical features found in metallic material fracture, indicating that Si melted and filled some cavities, in this way reducing porosity. Fig. 4. SEM micrographs of the composite II after pre-pyrolysis: (a) crosssection of the C fiber involved by Si–O–C ceramic; (b) matrix around fiber bundle after 1450 ◦C/2 h. Besides this, the formation of Si2ON2 is also responsible for reducing porosity by the volume expansion of Si during nitridation, compensating the polymer shrinkage. After 1500 ◦C the porosity is increased. Fig. 6 shows the high level of porosity in the matrix between fiber bundles of the composite III. These results imply that at this temperature the gas phase plays a major role for material transport as well as filler reaction during pyrolysis
72 R. M. Rocha et al. Materials Science and Engineering A 437(2006)268-273 10m o um Fig. 5. SEM micrographs of the composite ll after pyrolysis at 1450Cn2 h:(a) cross-section of the C fiber involved by si-O-C ceramic; (b)matrix around fiber Fig. 6. SEM micrographs of the composites after pyrolysis at 1500C/I h:(a) composite II;(b)composite Ill. The flexural strength of the composites after one cycle of impregnation and pyrolysis at 1450C/2 h was tested by four-point bending method. The results are: 11.7+0.4 MPa for composite 1: 15.9+0.5 MPa for composite Il; 16.3+0.6 MPa for composite Il. A typical stress-displacement curve of each composite can be seen in Fig. 7. It can be noted the brittle like o failure mode. The reasons for the relatively low flexural strength 3 and the brittle failure, besides the high porosity, is concerned to o 6 a strong fiber-matrix bond. Interfacial strength can be evaluated by the morphology of the fracture surface; long fibre pullout indicates weak interfacial strength, whereas short fibre pullout implies in a strong interfacial bonding between carbon fiber and matrix. To elucidate the failure behavior of the composites, the specimens were observed by SEM after the bending tests ig. 8). It is noted that composites exhibited flat fracture surface with very short fiber pullout. The interface properties in com- Displacement (mm) posites are considered to be relative to the surface characteristic Fig. 7. Stress-displacement curves of the composites pyrolyzed at 1450C/h
272 R.M. Rocha et al. / Materials Science and Engineering A 437 (2006) 268–273 Fig. 5. SEM micrographs of the composite II after pyrolysis at 1450 ◦C/2 h: (a) cross-section of the C fiber involved by Si–O–C ceramic; (b) matrix around fiber bundle. The flexural strength of the composites after one cycle of impregnation and pyrolysis at 1450 ◦C/2 h was tested by four-point bending method. The results are: 11.7 ± 0.4 MPa for composite I; 15.9 ± 0.5 MPa for composite II; 16.3 ± 0.6 MPa for composite III. A typical stress-displacement curve of each composite can be seen in Fig. 7. It can be noted the brittle like failure mode. The reasons for the relatively low flexural strength and the brittle failure, besides the high porosity, is concerned to a strong fiber–matrix bond. Interfacial strength can be evaluated by the morphology of the fracture surface; long fibre pullout indicates weak interfacial strength, whereas short fibre pullout implies in a strong interfacial bonding between carbon fiber and matrix. To elucidate the failure behavior of the composites, the specimens were observed by SEM after the bending tests (Fig. 8). It is noted that composites exhibited flat fracture surface with very short fiber pullout. The interface properties in composites are considered to be relative to the surface characteristic Fig. 6. SEM micrographs of the composites after pyrolysis at 1500 ◦C/1 h: (a) composite II; (b) composite III. Fig. 7. Stress–displacement curves of the composites pyrolyzed at 1450 ◦C/2 h
R.M. Rocha et aL. Materials Science and Engineering A 437(2006)268-273 4. Conclusions The method of manufacturing CMCs by AFCOP proces using a system of active filler Si/poly(methtylsilsesquioxane), has proved that it is possible to produce C fiber-ceramic matrix composite in an inexpensive way. Three composites were man ufactured with one cycle of impregnation using different slurry compositions and pre-pyrolyzed at 1000C and pyrolyzed in wo distinct temperatures, 1450C/2h and 1500C/ h Regard less si: polymer ratio the inter-fiber spaces were impregnated by just the polymer phase, which became a Si-O-C based ceramic This inter-fiber phase showed a lot of porosity after pyrolysis at 1450C/2h. However. the matrix around fiber bundles were responsible for reducing porosity by means of Si melting and reactions between polymer residue, Si filler and N2 atmosphere Porosity in the matrix between fiber bundles was responsible for the change in total open porosity as a function of pyroly N sis temperature. The determined flexural strength was very low, because of the high level of porosity, and also owing to the strong fiber-matrix bond. The first could be reduced aft cles of impregnation, and the second could be changed with some modifications in the process, such as higher heating rates Acknowledgment The authors are thankful to Jose Carlos Troni for his helpful assistance in the composite preparation References [1] H. Ohnabe, S Masaki, M. Onozuka, K. Miyahara. T Sasa, Compos. Part [2] D. Belitskus, Fiber and Whisker Reinforced Ceramics for Structural App cations. Marcel Dekker Inc. New York. 1993. [3]R. Jones, A Szweda, D Petrak, Compos. Part A 30(1999)569-575 [4 W. Krenkel, F Berndt, Mater. Sci Eng. A 412(2005)177-181 Fig & SEM micrographs of the fracture surface after flexural test: (a) composite [5]K Kobayashi, H Sano, Y Uchiyama, Key Eng Mater. V 108-110(1995) [6] C.A.A. Cairo, C.R. M. Silva, M.L. Graca, J.C. Bressiani, J. Eur. Ceram. Soc.21(2001)325-329. of fibres, microstructure of matrix and processing conditions [71 Z.S.Rak,JAm Ceram Soc. 84(10)(2001)2235-2239 of the composites. Reaction between the carbon fibres and [8] w.w. Zheng, Z.H. Chen. Q.S. Ma, H.E. Hu, J Mater. Sci. 39(2004) 3521-3522 the Si-o-C based matrix on pyrolysis is a possible cause in [9] H.Q. Ly, R. Taylor, R.J. Day, I Mater. Sci. 36(2001)4027-4035 the embrittlement of the composite [8-11. It has been well [10] Q.S. Ma, Z.H. Chen, w.W. Zheng, H.E. Hu, Mater. Sci Eng. A 352(2003) with carbon fibers during pyrolysis to create strong fiber/matrix G.B. Zheng. H Sano, Y. Uchiyama, K. Kobayashi, H.M. Cheng. J. Mater bonding 19, 10J. During the fabrication of the composites, there [12)P. Greil, I. Am. Ceram Soc. 78(4)(1995)835-848 was quite a lot of time due to the low heating rate for Si atoms [13) P. Greil, J. Eur. Ceram Soc. 18(1998)1905-1914 in matrix to diffuse into carbon fibres. Moreover, the used [14] D Suttor, T Ermy, P. Greil, J Am Ceram Soc. 80(7)(1997)1831-1840. carbon fibers fabricated below 1500C have more impurity [15] G.M. Relund S Prochaska, R Doremus, J Mater Res. 12(6)(1991) elements, which are left by the precursors [ll] which favors 2716-2722 the reactions with the PMS derived Si-o-C matrix. Therefore [6] M. Scheffler, E. Pippel, J. Woltersdorf, P. Greil, Mater. Chem. Phys. 80 (2003)565-572 composites showed low flexural strength and brittle fracture [17)QWei, E. Pippel, J Woltersdorf, M Scheffler, P. Greil, Mater. Chem. Phys. behavior 73(2002)281-289
R.M. Rocha et al. / Materials Science and Engineering A 437 (2006) 268–273 273 Fig. 8. SEM micrographs of the fracture surface after flexural test: (a) composite II; (b) composite III. of fibres, microstructure of matrix and processing conditions of the composites. Reaction between the carbon fibres and the Si–O–C based matrix on pyrolysis is a possible cause in the embrittlement of the composite [8–11]. It has been well documented that the atoms in the matrix can diffuse into or react with carbon fibers during pyrolysis to create strong fiber/matrix bonding [9,10]. During the fabrication of the composites, there was quite a lot of time due to the low heating rate for Si atoms in matrix to diffuse into carbon fibres. Moreover, the used carbon fibers fabricated below 1500 ◦C have more impurity elements, which are left by the precursors [11] which favors the reactions with the PMS derived Si–O–C matrix. Therefore composites showed low flexural strength and brittle fracture behavior. 4. Conclusions The method of manufacturing CMCs by AFCOP process, using a system of active filler Si/poly(methtylsilsesquioxane), has proved that it is possible to produce C fiber–ceramic matrix composite in an inexpensive way. Three composites were manufactured with one cycle of impregnation using different slurry compositions and pre-pyrolyzed at 1000 ◦C and pyrolyzed in two distinct temperatures, 1450 ◦C/2 h and 1500 ◦C/1 h. Regardless Si:polymer ratio the inter-fiber spaces were impregnated by just the polymer phase, which became a Si–O–C based ceramic. This inter-fiber phase showed a lot of porosity after pyrolysis at 1450 ◦C/2 h. However, the matrix around fiber bundles were responsible for reducing porosity by means of Si melting and reactions between polymer residue, Si filler and N2 atmosphere. Porosity in the matrix between fiber bundles was responsible for the change in total open porosity as a function of pyrolysis temperature. The determined flexural strength was very low, because of the high level of porosity, and also owing to the strong fiber–matrix bond. The first could be reduced after more cycles of impregnation, and the second could be changed with some modifications in the process, such as higher heating rates. Acknowledgment The authors are thankful to Jose Carlos Troni for his helpful ´ assistance in the composite preparation. References [1] H. Ohnabe, S. Masaki, M. Onozuka, K. Miyahara, T. Sasa, Compos. Part A 30 (1999) 489–496. [2] D. Belitskus, Fiber and Whisker Reinforced Ceramics for Structural Applications, Marcel Dekker Inc., New York, 1993. [3] R. Jones, A. Szweda, D. Petrak, Compos. Part A 30 (1999) 569–575. [4] W. Krenkel, F. Berndt, Mater. Sci. Eng. A 412 (2005) 177–181. [5] K. Kobayashi, H. Sano, Y. Uchiyama, Key Eng. Mater. V 108–110 (1995) 145–154. [6] C.A.A. Cairo, C.R.M. Silva, M.L. Grac¸a, J.C. Bressiani, J. Eur. Ceram. Soc. 21 (2001) 325–329. [7] Z.S. Rak, J. Am. Ceram. Soc. 84 (10) (2001) 2235–2239. [8] W.W. Zheng, Z.H. Chen, Q.S. Ma, H.F. Hu, J. Mater. Sci. 39 (2004) 3521–3522. [9] H.Q. Ly, R. Taylor, R.J. Day, J. Mater. Sci. 36 (2001) 4027–4035. [10] Q.S. Ma, Z.H. Chen, W.W. Zheng, H.F. Hu, Mater. Sci. Eng. A 352 (2003) 212–216. [11] G.B. Zheng, H. Sano, Y. Uchiyama, K. Kobayashi, H.M. Cheng, J. Mater. Sci. 34 (1999) 827–834. [12] P. Greil, J. Am. Ceram. Soc. 78 (4) (1995) 835–848. [13] P. Greil, J. Eur. Ceram. Soc. 18 (1998) 1905–1914. [14] D. Suttor, T. Erny, P. Greil, J. Am. Ceram. Soc. 80 (7) (1997) 1831–1840. [15] G.M. Relund, S. Prochaska, R. Doremus, J. Mater. Res. 12 (6) (1991) 2716–2722. [16] M. Scheffler, E. Pippel, J. Woltersdorf, P. Greil, Mater. Chem. Phys. 80 (2003) 565–572. [17] Q. Wei, E. Pippel, J. Woltersdorf, M. Scheffler, P. Greil, Mater. Chem. Phys. 73 (2002) 281–289