SiC ceramic matrix composites

Materials Science and Engineering B 168 (2010) 204–207

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Materials Science and Engineering B journal homepage: www.elsevier.com/locate/mseb

Allylhydridopolycarbosilane (AHPCS) as matrix resin for C/SiC ceramic matrix composites R. Sreeja, B. Swaminathan ∗ , Anil Painuly, T.V. Sebastian, S. Packirisamy Ceramic Matrix Products Division, Propellants & Special Chemicals Group, PCM Entity, Vikram Sarabhai Space Center, Thiruvananthapuram 695022, India

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Article history: Received 1 August 2009 Received in revised form 9 November 2009 Accepted 6 December 2009 Keywords: Preceramic polymer Allylhydridopolycarbosilane Polymer impregnation and pyrolysis process (PIP) Ceramic matrix composite (CMC)

a b s t r a c t In present study, partially allyl-substituted hydridopolycarbosilane (5 mol% allyl) [AHPCS] has been characterized by spectral techniques and thermal analysis. The DSC studies show that, the polymer is self-cross-linking at lower temperatures without any incorporation of cross-linking agents. The spectral and thermal characterizations carried out at different processing stages indicate the possibility of extensive structural rearrangement accompanied by the loss of hydrogen and other reactions of C and Si containing species resulting in the conversion of the branched chain segment into a 3D SiC network structure. AHPCS gave ceramic residue of 72% and 70% at 900 and 1500 ◦ C respectively in argon atmosphere. XRD pattern of 1500 ◦ C heat-treated AHPCS, indicates the formation of silicon carbide with the particle size of 3–4 nm. AHPCS was used as matrix resin for the preparation of C/SiC composite without any interfacial coating over the T-300 carbon fabric reinforcement. Flexural strength value of 74–86 MPa for C/SiC specimen with density of 1.7 g/cm3 was obtained after four infiltration and pyrolysis cycles. © 2009 Elsevier B.V. All rights reserved.

1. Introduction The C/SiC ceramic matrix composites (CMCs) with carbon as the reinforcement and SiC as the matrix are being widely explored as thermo structural materials for various space applications. Polymer impregnation and pyrolysis (PIP) process is more advantageous compared to chemical vapor infiltration (CVI) process for its easy processing, effective microstructure control and cost [1–3]. However, the proper choice of preceramic polymer is of importance in order to get C/SiC composite using less number of PIP cycles. Polycarbosilane (PCS), [SiH(CH3 )CH2 ]n , derived from polydimethylsilane, has been widely explored as matrix resin for C/SiC composite. PCS, on heat treatment in inert atmosphere at temperatures above 1400 ◦ C give ceramic residue comprising of ␤-SiC with considerable amount of free carbon. Excess carbon leads to decreased oxidative stability, poor crystallinity and less than fully satisfactory mechanical properties [4,5]. Partially allyl-substituted hydridopolycarbosilane (5 mol% allyl) [AHPCS] has been reported to have a compositional formula [Si(CH2 CH CH2 )2 CH2 ]x [SiH2 CH2 ]n−x (where x = 0.5%) obtained by Grignard coupling of chloromethyltrichlorosilane, followed by incorporation of small amount of allyl groups and then reduction with LiAlH4 [6,7]. Of late, AHPCS has been explored as a matrix

∗ Corresponding author. Tel.: +91 471 256 4624/9446174564; fax: +91 471 256 4096. E-mail address: [email protected] (B. Swaminathan). 0921-5107/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.mseb.2009.12.033

resin for C/SiC composite for the following reasons: (i) relatively airstable, (ii) liquid at room temperature, (iii) less viscous and hence better penetration into reinforcement tows [6–9]. Even though many reports are available on the use of AHPCS as matrix resin for C/SiC composites in which boron nitride [BN] or pyrolytic carbon is used as an interfacial coating for the reinforcement [10], there appears to be no report available on the preparation of a similar composite without using any interfacial coating. Hence, in this study, C/SiC composite has been prepared from partially allyl-AHPCS and carbon (T-300) as reinforcement without using any interfacial coating; thereby avoiding the coating step and use of costly BN coating. In the present investigation, attention also has been focused on the determination of basic characterization of AHPCS, cure characteristics, thermal stability, ceramic conversion and the flexural strength of the composite, since such data is not available in literature. 2. Experimental AHPCS-SMP 10 [Starfire, USA], weight average molecular weight (Mw) 450 Da, viscosity 0.1 Pa s, color light yellow, SiC powder (Saint Gobin Ceramic Materials, Norway, Purity 99%, particle size 17 ␮m) were used as received. 1 H, 13 C and 29 Si NMR spectra were recorded in CDCl3 on Brucker Avance 300 NMR Spectrometer at 300, 75.5 and 59.6 MHz respectively. Chemical shifts were reported with respect to internal tetramethylsilane standard. DSC, DTA and TGA were performed on a TA instrument SDT 2960 at a heating rate of 10 ◦ C/min under nitrogen atmosphere. The residual free carbon

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corresponding to carbon of solvent CDCl3 used are observed at 76–78 ppm. In the 29 Si NMR spectrum, resonances corresponding to disilyl linkages are observed in the region −10.9 to −13.68 ppm. Polysilyl linkages are observed in the region −33.7 to −38.4 ppm. Silicon atoms present in the structural unit H2 SiC is observed in the

Fig. 1. DSC curve of AHPCS.

present in the pyrolyzed sample was analyzed by TGA in air at a heating rate of 10 ◦ C/min [11] as the method given in the literature. AHPCS heat-treated at 1500 ◦ C was characterized by XRD (Philips 1729 Instrument using Cu-K␣ radiation) and SiC assay by KHSO4 fusion (wet analysis). AHPCS + 10 vol% SiC slurry was used as the precursor for the ceramic matrix. The formulation was optimized so as to ensure the C/SiC volume percentage as 40/60 for the resultant composite. The slurry was coated over T-300 carbon fabric [Toray, Japan, 5-harness satin structural grade]. The coated fabric was precured at 125 ◦ C in an air oven for 30 min. It was cut into plies of size 150 mm × 150 mm and six such plies were stacked for fabricating a composite of 3 mm thickness. It was further consolidated and cured by heating from 30 ◦ C up to 300 ◦ C at a heating rate of 1 ◦ C/min in air. The green composite was then pyrolyzed up to 900 ◦ C (dwell time: 2 h) and sintered up to 1500 ◦ C under argon at a heating/cooling rate of 3 ◦ C/min and dwell time of 4 h at 1500 ◦ C. The resultant C/SiC was further densified by infiltrating with AHPCS + SiC slurry under normal pressure followed by curing at 300 ◦ C. The infiltrated composite was heat-treated up to 1500 ◦ C under argon atmosphere (PIP process). This densification cycle was repeated one more time with AHPCS + SiC slurry. Two more cycles of infiltration were carried out using AHPCS alone. Flexural samples, rectangular geometry, of size 60 × 9 × 3 mm were machined out from the final sintered composite and flexural strength of the final composite was evaluated using a Universal testing machine (ASTM C-1161).

Fig. 2.

1

3. Results and discussion

Fig. 3.

13

AHPCS was characterized by spectral and thermal analysis. DSC of AHPCS was carried out to find out the cure temperature and the heating rate of composite during pyrolysis. DSC curve of AHPCS is shown in Fig. 1. The onset of the reaction is at 125 ◦ C where, allyl double bond starts cross-linking and cure maximum is observed at 170.6 ◦ C. A sharp exotherm is observed at 230.4 ◦ C and this has been attributed to self-cross-links via the loss of hydrogen. 1 H, 13 C and 29 Si NMR spectra for the un-treated AHPCS are shown in Figs. 2–4 respectively. In the 1 H NMR spectrum of AHPCS, resonances corresponding to CH2 CH, SiH, SiCH2 and SiCH3 are observed in the region of 4.8–6.1, 3–4.4, 1.6–2.1 and 0.4–0.5 ppm respectively. In the 13 C NMR spectrum, resonances corresponding to –CH and CH2 of allyl carbons are observed at 134.7 and 114 ppm respectively. Multiple resonances observed in the region 18–30 ppm are due to –C–H2 - of allyl and –CH2 – linkages resulting from the polymerization of allyl groups on standing. Resonances corresponding to Si–CH3 carbon are observed at 0.6 to −10 ppm. Resonances

Fig. 4.

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H NMR spectrum of AHPCS.

C NMR spectrum of AHPCS.

Si NMR spectrum of AHPCS.

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Fig. 5. FTIR spectra of virgin, cured and sintered AHPCS. Fig. 7. XRD pattern of SiC from AHPCS.

region −61.9 to −65.9 ppm. Multiplicity of resonances observed in the 29 Si NMR spectrum of AHPCS is due to different stereochemical environment. In literature, data about NMR spectrum of totally allyl-AHPCS is described [6] and our results for partially 5% allylAHPCS are on similar lines confirming the structure. FTIR spectra of virgin AHPCS, cured sample (at 200 ◦ C) and heat-treated sample (at 1500 ◦ C) are shown in Fig. 5. In the virgin polymer, the characteristic vibrations of SiH and allyl (C C) are observed at 2131 and 1630 cm−1 respectively. For the cured sample, the intensity of the vibration due to C C of allyl group decreases with respect to that of Si–CH3 (1250 cm−1 ) group and this has been attributed to cross-linking involving allyl group. In the case of the sintered sample, the characteristic vibrations of SiC and silica are observed at 801 and 1086 cm−1 respectively. TG curve of AHPCS is given in Fig. 6. From the TG curve it is noticed that the weight loss takes place in three stages. The major weight loss occurs below 200 ◦ C (16 wt% loss), which can be attributed to the loss of low molecular weight oligomers. A weight loss of 4% observed in the region 250–375 ◦ C is due to the loss of smaller amounts of pyrophoric silanes and hydrogen from the SiH2 polymeric chain resulting in onset of cross-linking. The third stage of weight loss (∼3%) noticed in the region 400–500 ◦ C is attributed to the loss of the CH4 [6]. The net ceramic residue of the pure resin was found to be 70% at 1500 ◦ C in argon. AHPCS was subjected to pyrolysis at 900 ◦ C in argon atmosphere.

Fig. 6. TG curves of AHPCS in nitrogen and pyrolyzed AHPCS in air.

Free carbon content was determined as per method reported in the literature [11]. The residual free carbon present in the pyrolyzed sample was analyzed by TGA in air at a heating rate of 10 ◦ C/min up to 900 ◦ C (Fig. 6). A weight loss of 1.3% was observed at 800 ◦ C which is attributed to the loss of free carbon. The free carbon present in the pyrolyzed AHPCS is very much less compared to many other preceramic polymers. Above 800 ◦ C, a slight weight gain (0.3%) is observed which is probably due to the oxidation of amorphous SiC. XRD pattern of AHPCS heat-treated at 1500 ◦ C is given in Fig. 7. Diffraction lines corresponding to (1 1 1), (2 2 0) and (3 1 1) planes of ␤-SiC are observed at 2Â = 35.70, 60.14 and 71.93 respectively. The particle size of ␤-SiC as calculated by using Scherr equation and XRD data is 3–4 nm. SiC content of the heat-treated sample is 99% based on chemical analysis, as determined by wet analysis using KHSO4 fusion procedure, where in direct results about SiC content were obtained. However, the infrared spectra of the heat-treated sample showed the presence of silica in addition to SiC, because the extinction coefficient of silica is much greater than that of silicon carbide [12]. Formation of ceramic from preceramic polymers is a two-step process. In the first step, the preceramic polymers are pyrolyzed at 900 ◦ C in inert atmosphere during which the pendant organic groups get cleaved leaving behind carbon and silicon in the reactive stage. At this stage the reactive elements fuse together to form the ceramic, which is amorphous in nature. In the second stage, the above amorphous residue is heat-treated at 1500 ◦ C to form crystalline ceramic [13]. AHPCS + 10 vol% SiC slurry was used as a precursor matrix for the preparation of C/SiC composite. SiC powder is added as a filler to reduce the volume shrinkage on heat treatment, the final porosity and the number of PIP cycles. When heat treatment of C/SiC composite is carried out at 1500 ◦ C, SiC filler covered with ␤-SiC (formed from AHPCS) gets sintered thereby, becoming an integral part of the ceramic matrix. Flexural strength of C/SiC composite having density of 1.7 g/cc, is in the range of 74–86 MPa (5 coupons), where as the flexural strength of C/SiC composites of density 2.0–2.2 g/cc with BN coating reported is about 250 MPa [10]. The flexural strength of the present C/SiC composite (without interface coating) could be improved by coating the carbon reinforcement and by increasing the number of infiltration cycles. The pyrolytic carbon or BN interface coating is known to provide enhanced flexural strength by reducing the friction between the fiber/matrix, which inhibits crack propagation through the fiber [14]. Flexural strength is also known to increase with the number of PIP cycles [15]. The present developed C/SiC composites without any interface coating could have the poten-

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tial applications where the flexural strength in the range 50 MPa is required. The additional advantages seen in the present study are (a) on conversion to ceramic, AHPCS gives ␤-SiC and less than 2% excess carbon (b) ceramic yield (∼72%) is much higher than that of polycarbosilane (50–60%) which means that less number of PIP cycles required than PCS (c) self-cross-linking at 125–230 ◦ C temperatures and (d) higher densification efficiency, higher modulus of elasticity, less open porosity and superior oxidation resistance than PCS.

Acknowledgements

4. Conclusions

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Partially allyl-AHPCS was characterized by 1 H, 13 C and 29 Si NMR, FTIR spectroscopy, and thermal properties by TGA and DSC. It is demonstrated that air-stable, liquid and easily curable partially allyl-substituted AHPCS begins cross-linking at about 125 ◦ C and maximum cure is observed at about 170 ◦ C; gives nano-3–4 nm size ␤-SiC of 99% purity in 70% yield at 1500 ◦ C in argon atmosphere. 5% allyl-AHPCS + 10 vol% SiC serves as a precursor matrix resin for the preparation of C/SiC composite without using any interfacial coating onto the reinforcement. Flexural strength of 74–86 MPa is obtained for optimized 40/60, C/SiC composite, after four cycles of polymer infiltration, pyrolysis and sintering process without any interfacial coatings. Such composites are potential candidates for applications where the structural loads are not high. There is an additional scope to improve the flexural strength by increasing the PIP cycles and by the way of giving interface coating to the reinforcement.

The authors would like to thank the authorities of VSSC for granting the permission to present the work, Analytical and Spectroscopy Division (ASD) for thermal and spectral analyses and Materials Characterization Division (MCD) for XRD analysis. One of the authors (B. Swaminathan) acknowledges ISRO for the grant of fellowship. References