Thermal conductivity, specific heat capacity, and emissivity of ceramic matrix composites at high temperatures

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1 High Temperatures ^ High Pressures, 2003/2004, volume 35/36, pages 169 ^ 177 DOI: /htjr105 Thermal conductivity, specific heat capacity, and emissivity of ceramic matrix composites at high temperatures Ru«diger Brandt ô, Martin FrieÞ ½, Gu«nther Neuer ô ô Institut fu«r Kernenergetik und Energiesysteme (IKE), Universita«t Stuttgart, D Stuttgart, Germany, fax: ; brandt@ike.uni-stuttgart.de; ½ Deutsches Zentrum fu«r Luftund Raumfahrt ev (DLR), Institut fu«r Bauweisen- und Konstruktionsforschung, D Stuttgart, Germany, martin.friess@dlr.de Received 12 August 2002; in revised form 6 June 2003 Abstract. Ceramic matrix composites (CMCs) with low porosity are obtained in one cycle via the well established liquid silicon infiltration (LSI) process, which is characterised by short processing times and fairly low manufacturing costs. Apart from aerospace applications, such as hot structures for re-entry vehicles, more and more applications beyond this classic field of CMCs are of increasing interest, eg brake discs, zero-expansion materials, high-temperature heat exchangers, heat-sink materials, etc. By applying special process parameters the microstructure as well as the physical properties can be tailor-designed to match specific requirements. The thermophysical properties of C/C ^ SiC, especially specific heat capacity, thermal conductivity as well as total and spectral emissivity were investigated in order to find out how the selection of fibres and pretreatment of the material affect them. Thermal conductivity of carbon fibres and therefore also of the composite considerably increases with increasing pyrolysis and graphitisation temperature. 1 Introduction Ceramic matrix composites (CMCs) with low porosity are obtained via the well established liquid silicon infiltration (LSI) process, which is characterised by short processing times and fairly low manufacturing costs (Krenkel 2000). The key step is the infiltration of liquid silicon in porous carbon/carbon (C/C) composites, whose distinct microstructure is formed in the preceding pyrolysis step of carbon-fibre-reinforced plastics (CFRP). These have been produced by resin transfer moulding (RTM) or lay-up of prepregs and subsequent autoclave technique in a first step. This process leads to C/C ^ SiC composites which are characterised by high mass-specific properties in combination with extreme thermal-shock stability. Besides aerospace applications, such as hot structures for re-entry vehicles, more and more applications beyond this classic field of CMCs are of increasing interest, eg brake discs, zero-expansion materials, high-temperature heat exchangers, etc. By applying special process parameters, the microstructure as well as the physical properties can be tailor-designed to match specific requirements. Thermal conductivity of carbon fibres and thus of a C/C ^ SiC composite is considerably increased by pyrolysis and graphitisation temperature as has been reported for C/C composites (Fitzer and Terwiesch 1974; Bo«der 1982). Therefore detailed investigations were carried out in order to find out how the thermophysical properties of C/C ^ SiC composites are influenced by using different carbon fibres, precursors, and annealing process parameters of C/C. Therefore, carbon fibre fabrics were used as received or thermally pretreated and submitted to LSI processing. Another route is to graphitise the C/C intermediate with subsequent siliconising (FrieÞ and Krenkel 1999a, 1999b). Because thermal conductivity of carbon fibres after graphitisation can be higher than that of SiC, the dominating contribution to the thermal conductivity of the composite is expected to come from the fibres.

2 170 R Brandt, M FrieÞ, G Neuer 2 Sample preparation In order to investigate the influence of process parameters on the thermophysical properties of C/C ^ SiC composites, samples of various sizes were cut from C/C ^ SiC plates which were manufactured by the LSI process which is described elsewhere in more detail (Krenkel 2000). Therefore, three carbon precursors, two phenolic resins JK25 and JK27, as well as the highly aromatic resin XP60, and three different pan-based carbon fibresöhta (high tenacity), T800 (intermediate modulus), and M40 (high modulus)öas well as a pitch-based carbon fibre (NGF from Nippon Graphite Fibre) were used (see table 1). Fibre diameters ranged from 5 mm for T800 fibres to 6.5 mm for M40 fibres, and 7 mm for HTA fibres and NGF fibres. All fibres were subjected to LSI processing in the form of fabrics (plain weave, satin in the case of M40). All parameter variations together with the individual sample designations are listed in table 2. At first, CFRP plates were processed from XP60 via RTM and from the phenolic resins via autoclave technique. Then, the CFRP plates with a fibre volume content of roughly 60% (56% ^ 63%) were pyrolysed to C/C at 900 8C or optionally at C. Some specimens were annealed above C in an inductively heated graphite furnace in an argon atmosphere for 1h. In the last step, the C/C plates were infiltrated with liquid silicon at C to form C/C ^ SiC. From the C/C ^ SiC plates, small samples were cut with a diamond blade for the determination of porosity and density (Archimedes method) as well as for microstructural characterisation in SEM. The determination of silicon, carbon, and SiC in the samples was carried out by gravimetric analysis. Therefore, Table 1. Properties of carbon fibres (manufacturer's) data. HT: high tenacity, IM: intermediate modulus, HM: high modulus. Fibre HTA T800 M40 NGF YS-15 Type HT IM HM pitch Manufacturer Tenax/Akzo Toray Toray Nippon Graphite Fibre Carbon content=% 90 ± > 99 >96 Diameter=mm Density=g cm Tensile strength=mpa Tensile modulus Elongation at break=% Thermal conductivity=w m 1 K 1 in longitudinal direction after graphitisation at C ± ± ± 590 Table 2. Process parameters, physical properties, and chemical composition of C/C ^ SiC specimens. (The structure of the sample designation is: fibre/precursor pyrolysis temperature.) Specimen H/X9 H*/X9 T/X9 T*/X9 M/X9 N/X9 H/X16 H/J16 H/K16 Fibre HTA HTA a T800 T800 b M40 NGF HTA HTA HTA Precursor XP60 XP60 XP60 XP60 XP60 XP60 XP60 JK25 JK27 Pyrolysis temperature=8c Porosity=% Density=g cm Si=mass% C=mass% SiC=mass% a HTA-fibre pretreated at C; b T800-fibre pretreated at 600 8C.

3 Thermophysical properties of ceramic matrix composites at high temperatures 171 silicon was removed by dissolving in a mixture of hydrofluoric and nitric acid (20/80) at 40 8C for 48 h, whereas the carbon content was measured by burning off carbon at 700 8C for 20 h in air yielding residual SiC. A typical microstructure of C/C ^ SiC derived from CFRP with high fibre matrix bonding (H/X9) is shown in figure 1 (left). Dense C/C segments (dark), which were not accessible by liquid silicon during siliconising, are surrounded by a SiC matrix containing some residual silicon. In contrast, figure 1 (right) shows the microstructure of C/C ^ SiC obtained from a CFRP with low fibre matrix bonding (H*/X9). In this case, owing to many small cracks in the C/C, liquid silicon can enter these cracks and attack all fibres during siliconising. Consequently a large amount of carbon fibres and the carbon matrix are converted to SiC, resulting in a significantly increased SiC content (60.9% compared to 33.0% for H/X9). All other samples based on the XP60 precursor, and fibres without fibre pretreatment revealed a microstructure similar to that shown in figure 1 (left) due to high fibre matrix bonding. The microstructure based on the two types of phenolic resins is shown in figure 2. Whereas the H/J16 (JK25-type) sample has an increased SiC content (44%) and a microstructure showing C/C segments with incorporated SiC areas, the H/K16 sample is similar to the XP60-type in microstructure as well as in SiC content. Annealing of C/C specimens (graphitisation) prior to siliconising does not fundamentally change the microstructure of C/C ^ SiC samples. However, the SiC and silicon contents are increased owing to the formation of larger spacing between the C/C segments. Figure 1. SEM micrographs of C/C ^ SiC composites with HTA fibres, pyrolysed at C; left: high fibre matrix bonding (H/X9), right: low fibre matrix bonding (H*/X9). Figure 2. SEM micrographs of C/C ^ SiC composites with HTA fibres, pyrolysed at C; left: JK25-type, right: JK27-type.

4 172 R Brandt, M FrieÞ, G Neuer 3 Methods used to measure the thermophysical properties Measurements of the specific heat capacity have been carried out at IKE by means of differential scanning calorimetry on a Perkin Elmer DSC2 at temperatures up to 730 8C. Thermal conductivity has been measured directly up to 800 8C by using a commercial longitudinal comparative instrument (Holometrix, Model TCFCM). In this technique a cylindrical specimen (25 mm in diameter, 25 mm thick) is sandwiched axially between two identical reference samples of the same diameter. The whole sample stack is placed between two heating elements, controlled at different temperatures, and a heat sink. The resultant heat flow through the sample stack produces temperature gradients in all three samples which are measured with thermocouples. At thermal equilibrium, the thermal conductivity of the test sample can be calculated from the measured temperature gradients and the well-known thermal conductivity of the reference samples. For higher temperatures, thermal conductivity was calculated from thermal diffusivity, specific heat capacity, and density of the samples. Thermal diffusivity can be measured at IKE in vacuum up to C (Brandt and Neuer 1981). One face of a small disc-shaped sample, 8 mm in diameter and about 1 ^ 2 mm thick, is heated by an intensitymodulated light beam of a xenon lamp. The resulting temperature modulations propagate through the sample. From the temperature modulations on the heated and unheated sample faces (measured with IR detectors) the propagation time can be estimated, which is inversely proportional to the thermal diffusivity. The spectral and the total emissivity have been measured at IKE by the radiation comparison technique in the temperature range 800 ^ C on disc-shaped samples, 15 mm in diameter and 3 ^ 5 mm thick. The radiation from the surface is measured by means of specially developed radiation detectors combined with exchangeable interference filters to select the spectral range. The radiation detectors are calibrated against a blackbody and the sample temperature is measured within a small radially drilled hole parallel and close to the sample surface. Details of the measurement technique are described by Neuer (1995). Two apparatuses are available, one with electron beam heating for measurements in vacuum (Neuer and Jaroma-Weiland 1998) and the second with induction heating for measurements in air or in a gas atmosphere (Neuer et al 1998). 4 Results and discussion 4.1 Specific heat capacity Specific heat capacity has been measured for the samples listed in the left five columns of table 2. With respect to the measurement results these and three other (not included in table 2) C/C ^ SiC materials can be divided into two groups: one with low SiC content but high carbon and silicon contents (H/X9, T/X9, T*/X9, M/X9), and another with a high SiC content but lower carbon and silicon contents (H*/X9). The results for two representative specimens are plotted in figure 3 (dots) together with literature values for carbon, silicon, and SiC (solid lines). Calculated values according to the average composition (mass%) of each group by using the literature values for carbon, silicon, and SiC and applying the rule of mixtures (dotted lines) show good agreement with the measured data. It is evident that SiC reduces the heat capacity values whereas increasing carbon content shifts them upwards. 4.2 Thermal conductivity The influence of the fibre type on the thermal conductivity of the composite is shown in figure 4. Four different composites have been tested, all of them with about 60 vol% of fibres, the same precursor (XP60), and pyrolysed at 900 8C. Only the type of fibres (HTA, T800, NGF, and M40) has been varied. Owing to the low fabrication temperature of the HTA and NGF fibres (below C) the conductivity of these C/C ^ SiC composites is similarly low, whereas composites of the intermediate-modulus fibre T800,

5 Thermophysical properties of ceramic matrix composites at high temperatures Specific heat capacity =J g 1 K T/X9 H*/X9 graphite SiC(30%) C(64%) Si(6%) SiC(61%) C(37%) Si(2%) SiC silicon Temperature=8C Figure 3. Specific heat capacity of C/C ^ SiC composites with different fibres and annealed samples, respectively. For comparison the values for graphite, SiC, and silicon are also plotted. Thermal conductivity=w m 1 K M40 (60 vol%) T800 (65 vol%) NGF (57 vol%) HTA (65 vol%) Temperature=8C Figure 4. Influence of carbon fibre type on the thermal conductivity (in transverse direction to the fibre planes) of C/C ^ SiC composites with precursor XP60, pyrolysed at 900 8C. and especially the high-modulus fibre M40, show significantly higher values, by about 40% and 100% at 200 8C, respectively. As expected for ceramics and carbon, the thermal conductivity of all four composites decreases with increasing temperature. Because of the pronounced orientation of the fibres in one plane there is a strong anisotropy of the thermal conductivity in directions parallel and transverse to main fibre orientation. Thermal conductivity parallel to the fibre planes (not shown here) is higher by a factor 2 to 3 than in the transverse direction, as shown by the anisotropy factor in figure 5. The type of precursor also influences the thermal conductivity of the composite, as can be seen in figure 6. Here the conductivity of three composites with different precursors (XP60, JK27, and JK25) but the same type of fibres (HTA, 60 vol%) and pyrolysed at C, is plotted as a function of temperature. The composite with the XP60 precursor shows the lowest conductivity, the use of JK27 precursor leads to an increase of about 18% at 200 8C, and the JK25 precursor improves the thermal conductivity by about 64%. This is mainly due to a change in microstructure (figure 2). The JK25 precursor leads to C/C ^ SiC with a higher SiC content and thus an increase in thermal conductivity.

6 174 R Brandt, M FrieÞ, G Neuer Anisotropy factor of thermal conductivity H/X16 M/X9 H/X9 T*/X9 T/X9 H*/X Temperature=8C Figure 5. Anisotropy factor of thermal conductivity of C/C ^ SiC composites with different fibres and XP60 precursor. 20 Thermal conductivity=w m 1 K HTA/JK25 HTA/JK27 HTA/XP Temperature=8C Figure 6. Influence of precursor type on the thermal conductivity of C/C ^ SiC composites with HTA-fibres (60 vol%), pyrolysed at C. Further improvement in thermal conductivity of C/C ^ SiC can be achieved by annealing the C/C at higher temperatures, as shown in figure 7. Four different composites (different fibres and different precursors) were annealed in the C/C state between 900 8C and C. For the two composites with low thermal conductivity (HTA/XP60 and NGF/XP60), which were originally pyrolysed at 900 8C, an increase of the pyrolysis temperature to C caused an increase in thermal conductivity by about 12% and 40% at 200 8C. An increase of the annealing temperature to C improved the conductivity enormously (to about 200%). Further increase of the annealing temperature to C or C caused a drop in thermal conductivity. In view of these observations, the influence of annealing temperature on the conductivity of the two remaining composites (HTA/JK25 and HTA/JK27) was tested only in the range 1650 ^ C. For both composites the rise in annealing temperature from C to C increased the conductivity only by about 37% and 16%, respectively, but a further increase in annealing temperature to C caused a further improvement in thermal conductivity. Thus, compared to XP60, the two phenolic precursors, JK25 and JK27, reveal a different trend in the dependence of thermal conductivity on annealing temperature.

7 Thermophysical properties of ceramic matrix composites at high temperatures 175 Thermal conductivity=w m 1 K HTA/JK25 20 NGF/XP60 15 HTA/JK27 10 HTA/XP Pyrolysis temperature=8c Figure 7. Influence of pyrolysis temperature on the thermal conductivity of various C/C ^ SiC composites at 200 8C. 4.3 Spectral and total emissivity The emissivity is much less influenced by fibre type and treatment than the thermal conductivity. Such effects are here mainly affected by changes at the surface itself. Owing to the crystal structure of the fibres with high electric conductivity in the longitudinal direction, the spectral behaviour of the composite material with the fabric plane parallel to the surface is like that of metals, as shown by Neuer (1995). If the spectral emissivity at various temperatures is plotted against the wavelength (figure 8), metallic character can be observed as indicated by the so-called X-point where the spectral emissivity is independent of temperature. This effect cannot be observed in SiC ^ SiC composites. Obviously, the carbon fibres are responsible for the metallic character, because of their low electric resistance at room temperature (about 15 mo m). Spectral emissivity C/C ^ SiC, chemically etched 1000 K 1300 K X-point 1300 K 1000 K 1600 K 1300 K 1000 K SiC ± SiC Wavelength=mm Figure 8. Spectral normal emissivity of C/C ^ SiC composites (HTA/XP60), chemically etched, in comparison with a SiC ^ SiC composite. However, after annealing the samples at temperatures above C the crystal structure of the fibres is changed resulting in a remarkable increase of the spectral emissivity at long wavelengths. This is valid for the HTA fibres, as figure 9 shows, but also for the samples with T800 fibres the slope of the spectral emissivity is reduced by pretreatment of the fibres. In the case of the total emissivity (figure 10), the influence of annealing is more pronounced for samples with HTA fibres, which corresponds to the difference

8 176 R Brandt, M FrieÞ, G Neuer Spectral emissivity H/X9 H/X9 annealed T/X9 T*/X Wavelength=mm Figure 9. Spectral normal emissivity of C/C ^ SiC composites with different fibres at C. The emissivity of the sample with HTA fibre was measured before and after annealing. Total normal emissivity H/X9 H/X9 annealed T/X9 T*/X Temperature=8C Figure 10. Total normal emissivity of C/C ^ SiC with different fibres and various pretreatments of fibres (T800) or samples (HTA), respectively. between the spectral emissivity curves of both samples being larger in the untreated case than after annealing. Annealing of C/C ^ SiC samples also leads to a reduction of residual free silicon in the material, but earlier investigations (Neuer et al 1995/1996) have shown that desiliconising by chemical etching influences spectral behaviour of the emissivity much less. So it must be assumed that changes in emissivity by thermal treatment are caused by a change in fibre structure. 5 Conclusions Thermophysical properties of C/C ^ SiC, especially the thermal conductivity and the spectral emissivity, are mainly influenced by the selection of the fibres, while the specific heat capacity is due to both fibre and matrix properties. In addition, the thermal conductivity of C/C ^ SiC is significantly improved by annealing the composites in the C/C state prior to siliconising, when the use of the phenolic precursor JK25 and the pitchbased fibre NGF can further improve the thermophysical properties, especially the thermal conductivity, of C/C ^ SiC. On the basis of this knowledge it is possible to design C/C ^ SiC composites where special properties are required, eg aerospace applications, high-temperature heat exchangers, and heat-sink materials.

9 Thermophysical properties of ceramic matrix composites at high temperatures 177 Both the type and the pretreatment of fibres, as well as the surface structure of the material are important parameters leading to variations of the spectral emissivity especially at long wavelengths. Because of the spectral emissivities in the low wavelength rangeömainly contributing to the radiation at high temperaturesöthe total emissivity tends to higher values with increasing temperature. This is especially important if the material is used for protection at high temperatures. Acknowledgments. The authors wish to express their gratitude to the Deutsche Forschungsgemeinschaft (DFG) for financial support in the frame of Sonderforschungsbereich 259. References Bo«der H A, 1982 Angew. Makromol. Chem. 109/ Brandt R, Neuer G, 1981 Brennst. ^ Wa«rme ^ Kraft ^ 112 Fitzer E, Terwiesch B, 1974 Z. Werkstofftech. 5(2) 53 FrieÞ M, Krenkel W, 1999a, in Proceedings of the 9th CIMTECöWorld Ceramics Congress & Forum on New Materials Ed. P Vincenzini (ISBN: ) pp 127 ^ 134 FrieÞ M, Krenkel W, 1999b, in Verbundwerkstoffe und Werkstoffverbunde (Hamburg: John Wiley, VCH) pp 442 ^ 447 Krenkel W, 2000, PhD thesis, University of Stuttgart, Stuttgart, Germany Neuer G, 1995 Int. J. Thermophys ^ 265 Neuer G, Jaroma-Weiland G, 1998 Int. J. Thermophys ^ 929 Neuer G, Kochendo«rfer R, Gern F, 1995/1996 High Temp. ^ High Press. 27/ ^ 189 Neuer G, Pohlmann P, Schreiber E, 1998, in VDI-Berichte 1379 (Du«sseldorf: VDI) pp 173 ^ 178

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