Growth of SiC thin films on graphite for oxidation-protective coating

Growth of SiC thin films on graphite for oxidation-protective coating J.-H. Boo, a) M. C. Kim, and S.-B. Lee Department of Chemistry, Sungkyunkwan University, Suwon 440-746, Korea S.-J. Park and J.-G. Han Department of Metallurgical Engineering, Sungkyunkwan University, Suwon 440-746, Korea Received 1 December 1999; accepted 1 May 2000 We have deposited thick SiC thin films on graphite substrates in the temperature range of 700 850 C using single-molecular precursors by both thermal metal organic chemical-vapor deposition MOCVD and plasma-enhanced MOCVD PEMOCVD methods for oxidation-protection wear and tribological coating. Hexamethyldisilane HMDS, (CH 3 3 Si Si CH 3 ) 3, was used as a single-source precursor, and hydrogen and Ar were used as a bubbler and carrier gas. A highly oriented polyerystalline cubic SiC layer in the 111 direction was successfully deposited on graphite at temperatures as low as 800 C with HMDS by PEMOCVD. For thermal MOCVD, on the other hand, only amorphous SiC layers were obtained at 850 C. From this study, we confirmed that PEMOCVD was a highly effective process in improving the characteristics of the SiC layers compared to those grown by thermal MOCVD. The mechanical and oxidation-resistant properties have been assessed. The optimum SiC film was obtained at 850 C and rf power of 200 W. The maximum deposition rate and microhardness are 2 m/h and 4336 kg/mm 2 Hv, respectively. The hardness was strongly influenced by the composition ratio of the SiC protective layers. 2000 American Vacuum Society. S0734-2101 00 18204-1 I. INTRODUCTION Graphite, with its advantages of high thermal conductivity, low-thermal expansion coefficient, and low elasticity, has been widely used as a high-temperature structural material. 1,2 On the other hand, graphite can easily react with oxygen even at temperatures as low as 400 C, resulting in CO 2 formation. 2 To apply graphite as a high-temperature structural material, therefore, it is necessary to improve its resistance to oxidation. SiC is used as a semiconductor material for hightemperature, radiation-resistant, and high-power/highfrequency electronic devices because of its excellent properties. 3,4 Conventional chemical-vapor-deposited SiC has also been widely used as a coating material for structural applications because of its outstanding properties such as high thermal conductivity, high microhardness, and good chemical resistance to oxidation. 5 8 Therefore, SiC, with a similar thermal expansion coefficient as graphite, has recently been considered a good candidate protective coating to operate in high-temperature, corrosive, and high-wear environments. Due to the large lattice mismatch 50%, however, it has been very difficult to grow a thick SiC layer on graphite. Most SiC films are conventionally grown on silicon and graphite substrates by thermal chemical-vapor deposition CVD at above 1200 C using separate sources for silicon and carbon, such as silanes or chlorosilanes and various hydrocarbons. 9 11 However, it is desirable to grow the SiC films on silicon substrates at low temperature. A disadvantage of this method is the necessity to maintain a high growth a Author to whom correspondence should be addressed; electronic mail: jhboo@chem.skku.ac.kr temperature, which results in high tensile stress and crystalline lattice defects such as misfit dislocations and stacking faults in the SiC film due to the differences in lattice constants and thermal expansion coefficients between the SiC layer and the substrates. 12 In addition to this temperature problem, synthesis of SiC device structures by CVD requires precise control of the silicon-to-carbon ratio. The use of separate sources for the growth of SiC films may result in some deviations from stoichiometry in the films, leading to point defects or precipitates. A single-molecular precursor that contains Si and C atoms in the same molecule and already has a Si C bond can offer the advantage of such control and perhaps can reduce an activation barrier to SiC formation as well, resulting in low-temperature deposition. 13 17 Plasma-enhanced metal organic chemical-vapor deposition PEMOCVD can generate more reactive radicals than thermal CVD and can expect high growth rate and better film quality, especially smoothness. 18,19 In this study, therefore, we tried to make thick SiC coating layers on graphite using a new single-source precursor of hexamethyldisilane HMDS by thermal MOCVD and PEMOCVD methods to overcome the large lattice mismatch. To our knowledge, this is the first report on the growth of SiC coating layers from a precursor of HMDS by PEMOCVD. II. EXPERIMENTAL PROCEDURE The experiments were carried out in a stainless-steel vacuum chamber made in-house. Cylindrically shaped graphite model CX2123, CARBONIX Product used as a substrate was cleaned and etched using acetone, isopropyl alcohol, hydrofluoric 48% solution, and deionized water, sequentially. After cleaning the samples, the substrates were 1713 J. Vac. Sci. Technol. A 18 4, JulÕAug 2000 0734-2101Õ2000Õ18 4 Õ1713Õ5Õ$17.00 2000 American Vacuum Society 1713

1714 Boo et al.: Growth of SiC thin films on graphite 1714 pretreated in situ with an Ar plasma to make an oxygen-free surface and a buffer layer. The process of plasma pretreatment using Ar gas prior to CVD appears to increase adhesion of films on graphite, resulting in high hardness. The Ar pretreatment and deposition were carried out for 0.5 1 h at a rf power range of 100 300 W. The general deposition pressure and temperature were 2 4 10 2 Torr and 700 850 C, respectively, for thermal MOCVD and 4 6 10 1 Torr and 700 850 C, respectively, for PEMOCVD. The graphite substrate was heated using an indirect heating method, and the temperatures of the substrate were measured using an optical pyrometer. The typical conditions of the PEMOCVD process for film deposition are 100 300 W of rf power, 40 120 seem of Ar carrier gas, and 20 seem of H 2 bubbler gas. Hexamethyldisilane, (CH 3 3 Si Si CH 3 ) 3, was used as a single-molecular precursor. Due to high vapor pressure of the precursor itself, it was not necessary to heat the sources during deposition. The crystalline structures, surface morphologies, and film compositions for the as-grown SiC protective layers were characterized by x-ray diffraction XRD Rigaku, model: DMAX-33, Cu K, Ni filter, 35 kv 15 ma, X-ray photoelectron spectroscopy XPS VG Scientific Ltd., model: ESCALAB MKII, Al K, 15kV 20mA, and scanning electron microscopy Jeol, model: JSM-840A, respectively. The mechanical and oxidation-resistant properties have been checked with a Knoop microhardness tester and thermal gravimetric analyzer TGA. III. RESULTS AND DISCUSSION The resulting film was first characterized by XPS. Figure 1 a shows the x-ray photoelectron survey spectrum of the film deposited at 850 C and 200 W of rf power using HMDS. The survey spectrum clearly shows the photoelectron peaks of Si 2s, Si2p, and C 1s, and C(KVV) Auger signals, indicating the formation of silicon-carbide film. Besides these relevant peaks, there also appear the O 1s photoelectron peak and O(KVV) Auger peaks. Oxygen invariably shows up in the spectra of almost all the silicon-carbide films and can be attributed to surface contamination of the newly formed film by air and/or moisture during sample transfer. In Figs. 1 b and 1 c, the Si 2p and C 1s high-resolution XPS spectra of the same film as Fig. 1 a are shown. To compare the C 1s binding energy of this film with those of bulk silicon carbide and graphite, the peak at 283.3 ev corresponds to carbidic carbon, whereas the peak at 284.5 ev corresponds to graphitic carbon. 20 The C 1s binding energy 283 ev of the film closely matches that of the carbidic carbon of the bulk silicon carbide. The peak areas of Si 2p and C 1s strongly suggest that the films grown at 850 C and 200 W of rf power have a composition ratio Si:C 1:1.1 of silicon and carbon. The composition was varied with rf powers and deposition temperatures, and it can influence the hardness. For example, the composition ratios between Si and C are changed from 1:2 to 1:1.1 with increasing rf power and film growth temperatures. The more stoichiometric the FIG. 1. XPS spectra obtained from a SiC protective coating layer grown on graphite at 850 C and 200 W of rf power using HMDS by PEMOCVD: a survey scan spectra, b Si 2p, and c C1s high-resolution spectra. film grown under higher temperatures 700 850 C and high rf power in the range of 100 300 W, the better the hardness obtained. The structures of the films were identified by x-ray diffractometry. Figure 2 shows the x-ray diffraction patterns of the SiC films grown with HMDS by PEMOCVD Fig. 2 a and 2 b and thermal MOCVD Fig. 2 c. The XRD patterns of Figs. 2 a and 2 b showed a hint of crystalline SiC film formation, because the SiC coating layers exhibited relatively intense SiC 111 and SiC 220 peaks that appeared at 2 35 and 60, respectively. However, there is no other detectable phase in the diffraction pattern, except the peaks attributed to the graphite substrate. In Fig. 2 a, moreover, a relatively strong SiC 111 diffraction peak is observed from a SiC coating layer grown at 850 C by PEMOCVD, indicating that a highly oriented polycrystalline film in the 111 direction has been deposited on the graphite surface. The thickness of the as-grown film is about 4 m, according to the cross-sectional scanning electron microscope SEM measurement. As the deposition temperature was decreased to 800 C, a relatively poor polycrystalline SiC layer, shown J. Vac. Sci. Technol. A, Vol. 18, No. 4, JulÕAug 2000

1715 Boo et al.: Growth of SiC thin films on graphite 1715 FIG. 3. Plan-view and cross-sectional SEM images of the SiC coating layers grown on graphite using HMDS by a thermal MOCVD at 850 C, b PEMOCVD with 200 W of rf power at 700 C, c PEMOCVD with 200 W of rf power at 850 C, and d cross-sectional image of c. FIG. 2. XRD patterns of SiC protective coating layers grown on graphite using HMDS by a PEMOCVD with 200 W of rf power at 850 C, b PEMOCVD with 200 W of rf power at 800 C, and c thermal MOCVD at 850 C. in Fig. 2 b, was obtained under the same deposition condition compared to that of Fig. 2 a. Below 800 C, we only deposited more thin, carbon-rich amorphous films from the precursor of the HMDS, which suggests that the deposition temperature can be an important factor that influences film crystallinity and growth rate. The film composition ratio, determined by high-resolution XPS analysis, also changed dramatically with growth temperature and deposition process. The most stoichiometric film Si:C 1:1.1 was obtained from a SiC layer grown at 850 C and 200 W of rf power by PEMOCVD, rather than by MOCVD. For a film grown at the same temperature as Fig. 2 a, but by thermal MOCVD, we only obtained a broad peak around 2 60, which signifies an amorphous structure see Fig. 2 c. The average thickness of the as-grown SiC layers by thermal MOCVD are about 1 2 m, which is much less than those grown by PEMOCVD. This indicates that PEMOCVD can be a more highly effective process than thermal MOCVD in improving the characteristics of film quality and in making thicker films with a high growth rate. The changes in surface morphology of the as-grown films were examined by SEM. Figures 3 b and 3 c, obtained for films grown by PEMOCVD at 700 C Fig. 3 b and 850 C Fig. 3 c, show relatively smoother and largergrain, featureless surfaces of SiC coating layers than those of Fig. 3 a, grown by thermal MOCVD at 850 C. This is in good agreement with the XRD results, confirming that PEMOCVD can generate more reactive radicals than thermal CVD. The resulting high growth rate and better film quality, especially smoothness, can be explained with a nucleation effect. Because the argon plasma consists of ions, neutrals, and electrons, a number of SiC nucleation sites on the surface of the substrate will initially be increased and the electrons can make reactive radicals of the precursor atoms and/or molecules compared with that by thermal MOCVD, signifying the high growth rate. Thus, when we pretreated the substrate surface prior to SiC deposition the hardness and adhesion between the SiC layer and substrate were increased. By increasing the rf power, furthermore, a smooth film with large crystals was deposited. However, we obtained the smoothest, most oxygen-resistant, and thickest film under the rf power condition of 200 W, rather than 300 W. These results can be explained by a difference of plasma density and growth rate. In this study, therefore, the maximum deposition rate was obtained at 850 C and rf power of 200 W to be 2 m/h. This agrees well with the data of Fig. 3 d, which shows a 4- m-thick SiC coating layer with a very sharp interface between the SiC film and graphite substrate; this suggests good adhesion and a homogeneous coating layer at depth. To our knowledge, this is the first report on the growth of thicker polycrystalline SiC coating layers on graphite obtained from a HMDS precursor by PEMOCVD below 1000 C. Figures 4 a and 4 b show the variation of the film deposition rate as a function of CVD process and/or temperature Fig. 4 a and rf power Fig. 4 b. Here, the deposition rate was determined from the film thickness measured by crosssectional SEM images. Figure 4 a indicates that PEMOCVD can make a SiC layer with a higher growth rate JVST A - Vacuum, Surfaces, and Films

1716 Boo et al.: Growth of SiC thin films on graphite 1716 FIG. 4. Variation of film deposition rate as a function of a CVD process and b rf power. The changes of microhardness as a function of c rf power and d CVD process. than thermal MOCVD, and the deposition temperature also plays an important role in determining the growth rate. Figure 4 b also shows that the rf power can affect the growth rate. With increasing rf power from 200 to 300 W, however, the growth rate decreases rapidly due to a sputtering effect. In general, the plasma intensity is proportional to the rf power; thus, the highest-energy plasma density i.e., in our case, 300 W may induce sputtering processes, resulting in desorption and/or dissociation of the adsorbed species. This is why rather thinner films with a relatively low growth rate can be deposited at high rf powers, even though the same deposition condition was used during CVD. To check the mechanical properties of the as-grown SiC protective layers, we carried out a microhardness test analysis. For the hardness test under the same condition, we maintained a deposition temperature of 850 C and only varied the rf powers in the range of 100 300 W during the growth of the SiC coating layers on Si 100 by PEMOCVD. To compare the mechanical properties directly, we deposited the SiC films by thermal MOCVD under the same conditions of pressure and temperature as PEMOCVD, with the same precursor. The observed microhardness was significantly changed, from 1700 to 4336 kg/mm 2 Hv, depending on the applied rf powers see Fig. 4 c. Notice that the maximum microhardness 4336 kg/mm 2 Hv was obtained from a SiC film grown at 850 C and rf power of 200 W, rather than of 300 W. This reflects that the highly energized plasma can reduce hardness due to a sputtering effect, resulting in more carbon-rich films, or possibly, a high defect density in the film layers. Regarding the crystal quality and growth rate, judging from Figs. 2 a and 4 b, respectively, it is possible to explain that the best film with high hardness, smooth surface, and high growth rate can be deposited under a rf power of 200 W. Moreover, Fig. 4 d shows that hardness was also strongly influenced by the composition ratio of the SiC protective layers that were grown by both thermal MOCVD and PEMOCVD processes under different temperatures. In Fig. 4 d, the microhardness of the films grown by PEMOCVD FIG. 5. XRD patterns of SiC protective coating layers grown on graphite using HMDS by PEMOCVD with 200 W of rf power at 850 C under the Ar and H 2 gas ratio of a 2:1 and b 6:1. c shows the weight changes of the SiC protective coating layers from the graphite at 900 C. are significantly enhanced compared to those grown by thermal MOCVD. This indicates that the deposition temperature is the second factor to influence hardness. As mentioned in the previous section, the composition ratio of the SiC protective layers depends on the growth temperature. When the substrate temperature is increased from 700 to 850 C, a more stoichiometric SiC thin film, confirmed by XPS, was obtained by both thermal MOCVD and PEMOCVD. This indicates that the more stoichiometric the film, the greater the hardness of the film that can be deposited on the graphite surfaces. When the graphite surface was initially pretreated with Ar plasma, the hardness and adhesion between the SiC layer and graphite substrate were effectively increased due to a nucleation effect. To check the effect of an Ar-plasma treatment on hardness and crystallinity, we deposited the SiC coating layers with various Ar H 2 mixture gas conditions. Figures 5 a and 5 b show the typical XRD patterns of SiC coating layers grown at 850 C by PEMOCVD under different mixture gas conditions. The best film was obtained in the case of an Ar and H 2 gas ratio of 2:1. When we increased the Ar ratio to 6, however, a film with rather small thickness, compared to that of Fig. 5 a, was preferentially grown in the 220 J. Vac. Sci. Technol. A, Vol. 18, No. 4, JulÕAug 2000

1717 Boo et al.: Growth of SiC thin films on graphite 1717 direction. This indicates that the mixture gas ratio is also an influence on the film quality and film thickness, resulting in the different hardness and/or oxidation resistance properties. The characteristics of the oxidation protection test were also carried out using films grown with HMDS at a temperature in the range of 700 850 C and rf power of 200 300 W by PEMOCVD. The experimental condition for the thermal gravimetric analysis was 900 C of the sample temperature, with a heating rate of 50 C/min in an air atmosphere. Excellent behavior was obtained from a SiC coating layer grown at 850 C and rf power of 200 W see Fig. 5 c. This is in good agreement with the results of the microhardness test and surface morphology changes. IV. CONCLUSIONS SiC thin films for oxidation protection have been deposited on graphite in a temperature range of 700 850 C using a single-molecular precursor by thermal MOCVD and plasma-enhanced MOCVD. Hexamethyldisilane was used as a single-source precursor, and hydrogen and Ar were used as a bubbler and carrier gas. A highly oriented polycrystalline cubic SiC layer in the 111 direction was successfully deposited on graphite at a temperature as low as 800 C by PEMOCVD. For thermal MOCVD, however, only amorphous SiC layers were obtained at 850 C. From this work, we realize that PEMOCVD is a highly effective process in improving the characteristics of film compared to those of thermal MOCVD. The optimum SiC film was obtained at 850 C and rf power of 200 W. The maximum deposition rate and microhardness are 2 m/h and 4336 kg/mm 2 Hv, respectively. The hardness was strongly influenced with the composition ratio of the SiC layers. ACKNOWLEDGMENTS Two of the authors J.-H.B. and J.-G.H. acknowledge partial support from Korea Vacuum Tech., Ltd. and LG Cable & Machinery Ltd., Korea. This work was supported by the Brain Korea 21 Project and the Korea Science and Engineering Foundation Grant No. 1999-2-301-002-3. 1 W. Smith, and D. H. Leeds, Pyrolytic Graphite in Modern Materials Academic, New York, 1970. Vol. 7, p. 199. 2 E. Fitzer, Carbon 25, 163 1987. 3 D. M. Brown, IEEE Trans. Electron Devices ED-40, 325 1993 ; see, also, K. Shenai, R. S. Scott, and B. J. Baliga, ibid. ED-36, 1811 1989. 4 J. W. Palmour, J. A. Edmond, H. S. Kong, and C. H. Carter, Jr., Physica B 185, 461 1993. 5 J. Chown, R. F. Deacon, N. Singer, and A. E. S. Whote, Refractory Coatings on Graphite Academic, New York, 1963, p.81. 6 I. W. Donald and P. W. McMillan, J. Mater. Sci. 13, 1151 1978 ; see also, L. Jawed and D. C. Nagle, Mater. Res. Bull. 21, 1391 1986. 7 R. Brutson, Thin Solid Films 126, 313 1985 ; see also, X. Zhay, ibid. 215, 134 1992. 8 F. J. Buchanan and J. A. Little, Surf. Coat. Technol. 53, 137 1992. 9 S. Nishino, J. A. Powell, and H. A. Will, Appl. Phys. Lett. 42, 460 1986 ; see also, J. A. Powell and L. G. Matus, Amorphous and Crystallme Silicon Carbide Academic, New York, 1989. 10 H. Matsunami, Diamond Relat. Mater. 2, 1043 1993 ; see also, T. Ueda, H. Nishino, and H. Matsunami, J. Cryst. Growth 104, 695 1990. 11 C. E. Morosanu, Thin Films by Chemical Vapor Deposition Elsevier, Amsterdam, 1990. 12 I. Golecki, Mater. Res. Soc. Symp. Proc. 33, 1 1984. 13 I. Golecki, F. Reidinger, and J. Marti, Appl. Phys. Lett. 60, 1703 1992. 14 C. H. Wu, C. Jacob, X. J. Ning, S. Nishino, and P. Pirouz, J. Cryst. Growth 158, 480 1996. 15 J. Steckl, C. Yuan, J. P. Li, and M. J. Loboda, Appl. Phys. Lett. 63, 3347 1993 ; ibid. 64, 3000 1994. 16 J.-H. Boo, K.-S. Yu, M. Lee, and Y. Kim, Appl. Phys. Lett. 66, 3486 1995. 17 J.-H. Boo, S. A. Ustin, and W. Ho, Thin Solid Films 324, 124 1998 ; see also, K.-W. Lee, K.-S. Yu, J.-H. Boo, Y. Kim, T. Hatayama, T. Kimoto, and H. Matsunami, J. Electrochem. Soc. 144, 1474 1997. 18 H. Koinuma, Jpn. J. Appl. Phys., Part 1 25, 1811 1996. 19 C. H. Lin and H. L. Wang, Thin Solid Films 383, 171 1996. 20 J. H. Boo, K.-S. Yu, Y. Kim, S. H. Yeon, and I. N. Jung, Chem. Mater. 7, 694 1995. JVST A - Vacuum, Surfaces, and Films