High vacuum chemical vapor deposition of cubic SiC thin films on Si 001 substrates using single source precursor

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1 Surface and Coatings Technology High vacuum chemical vapor deposition of cubic SiC thin ž / films on Si 001 substrates using single source precursor J.-H. Boo a,, S.-B. Lee a, K.-S. Yu b, M.M. Sung b, Y. Kim a a Institute of Basic Science and Department of Chemistry, Sungkyunkwan Uni ersity, Suwon , South Korea b Ad anced Materials Di ision, Korea Research Institute of Chemical Technology, Taejon , South Korea Abstract Thin films of cubic SiC have been prepared on substrates in situ by high vacuum metal-organic chemical vapor deposition Ž HVMO-CVD. method using a single source precursor at various growth temperatures in the range of C. 1,3-Disilabutane, H Si-CH -SiH -CH Ž DSB that contains the same amount of silicon and carbon atoms in the same molecule was used as precursor without carrier gas. During HVMO-CVD, moreover, a series of as-deposited SiC thin films were simultaneously characterized by in situ X-ray photoelectron spectrometry Ž XPS. under the UHV condition without air exposure. XPS and Rutherford backscattering spectroscopy Ž RBS. show that the SiC films grown at above 700 C have stoichiometric composition. However, the films grown at below 700 C show Si-rich stoichiometry. Transmission electron microscopy Ž TEM. confirms the crystalline nature of the SiC films. The optimum temperatures for the formation of the epitaxial 3C SiC thin films were found to be between 900 and 1000 C on the basis of XRD and TED analysis. In this study, the best film with maximum growth rate of 0.1 m h was obtained from a SiC film grown at 900 C and Pa of DSB. The SiC Si interface is clearly 13 shown in secondary ion mass spectroscopy SIMS depth profile as judged by the sharp decrease C signals. The thickness of the as-grown films was determined using cross-sectional scanning electron microscopy Ž SEM. and RBS, and two different activation energies for 3C SiC formation were obtained from the Arrhenius plots Elsevier Science B.V. All rights reserved. Keywords: Cubic SiC thin film; High vacuum metal-organic chemical vapor deposition; Single source precursor; In situ XPS study 1. Introduction Single crystalline SiC thin film has attracted much interest for use in electronic and optoelectronic devices, and circuits designed to operate at high temperatures, high powers, high frequencies and high radiation environments due to its good electrical and mechanical Ž 2 characteristics such as electron mobility 1000 cm V. Ž 7 s 1, electron saturation velocity cm s. 2, breakdown electric field Ž , high melting point and high thermal conductivity. In particular, SiC has a wide band gap of 2.2 ev for the 3CŽ. -SiC, at Corresponding author. Fax: address: jhboo@chem.skku.ac.kr J.-H. Boo. room temperature 4, and has been exploited for SiC Si heterojunction bipolar transistors. Up to now, it is common to grow 3C SiC thin films heteroepitaxially on carbonized Si substrates by chemical vapor deposition Ž CVD. using separate Si and C sources such as SiH Ž or SiHCl. and C H Ž or CH with a carrier gas Ž H. 2, usually at temperatures higher than 1200 C 5 8. These high growth temperatures sometimes result in high tensile stress and crystalline lattice defects in the SiC films, when cooled down, due to the difference in thermal expansion coefficient between Si and SiC. Such defects and strain in heteroepitaxal films can degrade the carrier mobilities and increase the junction leakage current. Also, high growth temperatures may result in increased auto-doping and redistribution of dopants in the Si substrate. In order $ - see front matter 2000 Elsevier Science B.V. All rights reserved. PII: S

2 148 ( ) J.-H. Boo et al. Surface and Coatings Technology to achieve a lower growth temperature for 3C SiC, a clean growth ambient and a better control of growth rate are desirable. Thus, high vacuum metal-organic chemical vapor deposition Ž HVMO-CVD. at lower deposition temperatures of below 1000 C appears to be attractive in these respects. In addition, a clean sample surface can be easily prepared in ultra-high vacuum chamber. Recently, epitaxial 3C SiC films were grown on surfaces without a carbonization process at relatively low temperatures Ž C. by HVMO- CVD using the single source precursors that contain Si and C atoms in the same molecule, thus, it does not need bond formation energy between Si and C atoms, indicating the possibly of reducing SiC growth temperature In addition to the lowering of growth temperature, this method has the advantage of a much simpler deposition process and system in comparison with other CVD and molecular beam epitaxy Ž MBE. techniques, since only a single source is employed to deposit SiC thin films on without using any carrier gas. In this paper, we have studied on the epitaxial 3C SiC thin film growth on substrate without carbonization process and carrier gas using the new single molecular precursor of 1,3-disilabutane at relatively low deposition temperature in the range of C by high vacuum MO-CVD. 2. Experimental The experiments were carried out in situ in an ultrahigh vacuum system ŽESCALAB MK II, VG Scientific Ltd.. after introducing a specially designed sample preparation chamber for the X-ray photoelectron spectrometry Ž XPS. study. The CVD and XPS experiments could then be performed in the analytical chamber simultaneously and the substrate holder was placed on the analytical sample probe. The preparation chamber was modified for substrate heating using an e-beam heater. The substrate used in this study was SiŽ 001., cut into a square of mm 2. The substrate needs to be properly cleaned and prepared for the deposition to avoid residual surface impurities that can create defects in the growing films. Therefore, prior to introducing the Si substrate into the reactor, it was initially treated by a chemical cleaning process that involves degreasing, alkali treatment, acid treatment and rinsing in deionized water. This process is very similar to the method proposed by Ishizaka and Shiraki 14 to remove contamination and form a thin oxide layer on the surface. Before deposition, the e-beam and the sample holder assembly Ž made of Mo and Ta. were degassed at C for more than 1 h to minimize outgassing from the surface of the holder and filament. The substrate was then heated up to 850 C and kept at this temperature for approximately 20 min to remove the protective oxide layer. After such annealing, clean Si substrates could always be observed by in situ XPS examination. The control of substrate temperature was made by the filament current changes of the e-beam heater. The temperature of the substrate was measured with an optical pyrometer. The details of experimental set-up and deposition condition have been published elsewhere 12,13. The single precursor 1,3-disilabutane is a colorless liquid and its vapor pressure was measured to be 5333 Pa at room temperature. It was transferred into a glass bulb attached to the precursor handling system that isolated with a variable leak valve from high vacuum Ž 5 condition Pa. and was further purified by freeze-pump-thaw cycles using liquid nitrogen. SiC thin films were deposited directly on the clean surface without the carbonization process and carrier gas at varying temperatures Ž C. under high vac- Ž 4 3 uum conditions Pa.. The duration of deposition was approximately 2 6 h, and the growth rate changed depending on the experimental conditions. The grown 3C SiC films were characterized in-situ by XPS and ex situ by X-ray diffraction Ž XRD., scanning electron microscopy Ž SEM., transmission electron microscopy Ž TEM. transmission electron diffraction Ž TED., Rutherford backscattering spectroscopy Ž RBS. and secondary ion mass spectrometry Ž SIMS.. 3. Results and discussions The XPS investigation was carried out with Al K 1,2 radiation as excitation source Ž h ev.. The electron energy analyzer was operated in fixed analyzer transmission mode with constant pass energy of 50 ev. All measurements were performed at pressures lower than Pa in the analysis chamber. Corrections of the binding energy shift due to steady-state charging of the samples were made by taking the C 1s in SiC as reference at ev 15. Quantification of the XPS data was performed by normalizing the area of each peaks using the atomic sensitivity factors. Highresolution XPS spectra for Si 2p, C 1s and O 1s Ž 1 at.%. were obtained for all the SiC films at room temperature after the residual precursor evacuation. The typical XP spectra of Si 2p and C 1s for the SiC thin films grown at the various growth temperatures are presented in Fig. 1. In Fig. 1a, the Si 2p spectra of 450 C and of 300 C Ž 99.0 ev. show similar shape with the lack of any C 1s contribution due to SiC crystalline formation, indicating that no strong reaction between the precursor and substrate takes place until 450 C. However, the full width at half maximum Ž FWHM. of

3 ( ) J.-H. Boo et al. Surface and Coatings Technology Fig. 1. High resolution XPS spectra of SiC thin films grown on surfaces in situ by HVMO-CVD using 1,3-disilabutane at various deposition temperatures: Ž. a Si 2p and Ž. b C 1s peaks. the Si 2p XP peak of 450 C is broader than that of 300 C, suggesting a randomly oriented thin SiC cluster formation on the surface. This trend can distinctly be seen at the data of 600 C of which the Si 2p curve component is related to a more silicon-rich region Žits lower binding energy denoting Si atoms bonded to less carbon atoms than in SiC.. In the cases of SiC thin film samples grown above 700 C, the SiC component becomes dominant, denoting a strongly increased SiC thickness which attenuates or causes the disappearance of the Si substrate contribution. This result indicates that the Si 2p peak at ev and the C 1s peaks at ev correspond to the SiC phase. The Si:C ratio for the SiC films is for the SiC films obtained in the growth temperature range C. When the take-off angle is decreased from 90 to 20, the C Si ratio is seen to increase for the films obtained in the range C. The ratio at 90 take-off angle has the same trend, with the maximum at approximately 800 C. At higher temperatures above 1000 C, the film surfaces have less carbon than the bulk of the films. It was reported in literature that Si overlayers exist on CVD-grown SiC films at temperatures above 940 K 16. Fig. 2 shows a series of XRD patterns Ž 2 scan. obtained the as-grown SiC thin films deposited at different temperatures between 600 and 950 C for 6 h under the same CVD pressure of Pa, respectively. At 600 C, the XRD pattern shows the amorphous or nanocrystalline structure of the SiC layer. With increasing the deposition temperatures from 700 to 900 C, a significant increase of the 3C SiCŽ 001., Ž 004. diffraction peaks produces, showing a crystallinity improvement. However, in the Fig. 2, small peaks at- tributed to the 3C SiCŽ 111., Ž 220. diffraction are also appeared indicating that the Ž 001. preferred orientation of polycrystalline 3C SiC thin film was obtained at deposition temperatures in the range of C. However, the XRD pattern of 950 C shows no other SiC reflections. In addition to the Si peaks due to the substrate, however, we can see that only a single sharp peak due to the cubic structure of SiC thin film appears at for Ž 002. and for Ž 004. reflections. This signifies that at 950 C, the SiC film was grown epitaxially, resulting in the formation of the Fig. 2. X-ray diffraction patterns of SiC thin films grown on Si 001 in situ by HVMO-CVD using 1,3-disilabutane at various deposition temperatures.

4 150 ( ) J.-H. Boo et al. Surface and Coatings Technology Fig. 3. SEM images of SiC thin films grown on in situ by HVMO-CVD using 1,3-disilabutane at various deposition temperatures: Ž. a 700 C, Ž b. 800 C, Ž c. 900 C, and Ž d. 950 C. 3C SiC thin film with monocrystalline structure on substrate from the DSB. Fig. 3 shows changes in the surface morphology of the 3C SiC films grown at various temperatures Ž C. under the same deposition pressure of Pa for 6 h, respectively. The surface morphology of the films grown at 900 and 950 C show more smoother surfaces than those of the film grown at 700 and 800 C. The SEM image of 700 C shown in the Fig. 3a exhibits a circular shape of SiC crystals with submicron size on the substrates. It is believed that these areas acted as nucleation sites for the subsequent HVMO-CVD growth, as the 3C SiC films tended to be polycrystalline with very rough surface. As the substrate temperature increased to 800 C shown in the Fig. 3b, the surface roughness was more increased whereas the crystal size was increased, indicating improvement of crystal quality. With increasing the growth temperatures from 900 to 950 C, however, the surface roughness is decreased drastically and the crystal form of thin film layer is also transformed to a crystalline structure Ž see Fig. 3c,d.. The major crystal form of these deposited films were rectangular in shape on the substrates at 950 C and a voids are also shown with the bases of the inverted rectangular pyramidal voids located in the SiC film subsurface region. Growth temperature and pressure are very important factors in the attainment of single crystallinity in the films, as shown in the micrographs of plane view TEM image and TED patterns of Fig. 4, the 3C SiC films grown on substrates at various temperature Ž 700, 850, 900, 1000 C. and deposition pressure of Pa. Polycrystalline crystals can be seen in Fig. 4a, when the growth temperature was at 700 C. The insets of Fig. 4 show selected area electron diffraction Ž SAD. patterns of SiC film grown on SiŽ As the substrate temperature increased to 900 C from 700 C, the grain size was increased and a distinct spot-patterns containing rings were also observed. Increasing the temperature further to 1000 C allowed monocrystalline films to grow on the substrate Ž see Fig. 4d.. The spots seen in the SAD pattern of Fig. 4d are due to the SiC film; the inner strong spots to Si substrate, and outer weak spots to 3C SiC. Besides the major diffraction spots in other samples, the weak extra spots due to twinning in the diffraction pattern should be noted. The contrast seen in the cross-sectional view TEM micrographs is believed to be associated with mass contrast due to surface roughness and or misori- Fig. 4. Cross-sectional TEM images of SiC thin films grown on in situ by HVMO-CVD using 1,3-disilabutane at various deposition temperatures: Ž a. 700 C, Ž b. 850 C, Ž c. 900 C, and Ž d C. The insets of each figure show the selected TED pattern of the same films.

5 ( ) J.-H. Boo et al. Surface and Coatings Technology Ž. Ž. Fig. 5. Variation of film growth rate as a function of deposition temperature a and the Arrhenius plot b. ented or twinned regions surrounding the epitaxial growth which is similar to the films grown by CVD 9. The dependence of film growth rate on the film deposition temperatures was also studied. The SiC growth rates on have been obtained by the changes of film thickness measured by SEM, RBS, and ellipsometry as a function of temperature as shown in Fig. 5a. Two distinct growth temperature regions are apparent, i.e. the growth rate increases exponentially with deposition temperature in the range of C. This behavior is characteristic of a deposition process well known as kinetically controlled deposition in which the surface decomposition of the precursor of DSB is the rate determining step. The activation energy for the SiC film deposition calculated experimentally from the slope of Fig. 5b is approximately 67.4 kj mol. In the higher temperature region over 800 C, however, the Arrhenius plot shows a positive slope with negative activation energy Ž kj mol., indicating that the growth rate is controlled by the mass transport of reagents through the boundary layer to the growth surface. This is termed not only the region of diffusion controlled growth, but also possibly some portion of desorption and or pre-reaction reacted with CO or residual gases during thin film growth. The typical deposition rates for as-grown SiC films are 0.03 m h at 700 C, increasing to 0.1 m h at 850 C, and then decreasing to 0.02 m h at 1000 C. The RBS measurements were performed using 2.24 MeV 4 He 2 particles at normal incidence. A good Fig. 6. RBS and SIMS data for the SiC thin films grown on surfaces using 1,3-disilabutane at 1000 C: Ž a. RBS and Ž b. SIMS. The spectra Ž. i and Ž ii. in Ž a. are random RBS Ž. i and channeling spectra Ž ii., respectively.

6 152 ( ) J.-H. Boo et al. Surface and Coatings Technology depth resolution was achieved by using a scattering angle of 165. Crystalline quality and interface misfit were also checked by channeling of the film layers along the 001 direction. The last direct evidence of the crystallographic quality of the deposits is given by RBS channeling. Fig. 6a presents a comparison between the random RBS and channeling spectra for a typical HVMO-CVD sample, 1200 A thick with a growth time of 6 h, growth temperature of 1000 C and growth pressure of Pa. A very low minimum yield value, min, which is the ratio of the minimum counts between the aligned and random RBS spectra, is obtained, suggesting a high crystalline quality. The relationship between Ž min min channeling scatter- ing yield random scattering yield. 100; % value and crystallinity can directly be explained with the minimum value of min. Because a difference between random RBS Ž. i and channeling spectra Ž ii. are proportioned to the degree of the film quality. As growth temperatures increase and the min values decrease, thin film with good crystallinity can thus be produced. Positive SIMS depth profiles of all the samples using a 14.5 kv, 20 na, Cs beam were also performed to investigate the compositional variations of the SiC films with sputter time, and hence depth. A typical profile of the 3C SiC films Ž 1000 C. are shown in Fig. 6b. The SiC Si substrate interface is clearly shown in this depth profile as judged by the sharp decrease in the 16 O, 1 H, 13 C signals. However, it is observed that there is an increase in 16 O as one goes deeper into the interface of sample. A possible explanation is that there may be incorporation of oxygen during the growth of the SiC film or not removal native oxide layer. The species of 1 H, 13 C are also present and their amount increases with growth temperature, indicating that some of H or C species in the 3C SiC film are a result of unreacted DSB or DSB reaction intermediates. 4. Conclusions Thin films of cubic SiC have been prepared on by HVMO-CVD using 1,3-disilabutane at various temperatures. During CVD, moreover, a series of as-depositing SiC thin films were simultaneously characterized by in situ XPS under the UHV condition without air exposure. The optimum temperatures for the formation of the epitaxial 3C SiC thin films were found to be between 900 and 1000 C on the basis of XRD and TED analysis. XPS and RBS show that the SiC films grown at above 700 C have stoichiometric composition. However, the films grown at below 700 C show Si-rich stoichiometry. In this study, the best film with maximum growth rate of 0.1 m h was obtained from a SiC film grown at 900 C and Pa of DSB. The SiC Si interface is clearly shown in SIMS depth profile as judged by the sharp decrease 13 C signals. A dependence of film growth rate on the film deposition temperatures was also studied. Two distinct growth temperature regions are apparent, that is, the growth rate increases exponentially with deposition temperature in the range of C. This behavior is characteristic of a deposition process well known as kinetically controlled deposition in which the surface decomposition of the precursor of DSB is the rate determining step. The activation energy for the SiC film deposition calculated experimentally from the slope of Arrhenius plot is approximately 67.4 kj mol. In the higher temperature region over 800 C, however, a negative activation energy Ž kj mol. was obtained, indicating that the growth rate is controlled by the mass transport of reagents through the diffusion controlled reaction. Acknowledgements Support of this research by the Korea Science and Engineering Foundation Ž project No and by the Ministry of Education through BK21 program is gratefully acknowledged. References 1 W.E. Nelson, F.A. Halden, A. Rosengreen, J. Appl. Phys. 37 Ž D.K. Ferry, Phys. Rev. B 12 Ž W.V. Muench, I. Pfaffender, J. Appl. Phys. 48 Ž N.W. Jepps, T.F. Page, in: P. Krishna Ž Ed.., Progress in Crystal Growth and Characterization, 7, Pergamon, Oxford, 1983, pp S. Nishino, Y. Hazuki, H. Matsunami, T. Tanaka, J. Electrochem. Soc. 127 Ž P. Liaw, R.F. Davis, J. Electrochem. Soc. 132 Ž S. Nishino, H. Suhara, H. Ono, H. Matsunami, J. Appl. Phys. 61 Ž Y. Furumura, M. Doki, F. Mieno, T. Eshita, T. Suzuki, M. Maeda, J. Electrochem. Soc. 135 Ž I. Golecki, F. Reidinger, J. Marti, Appl. Phys. Lett. 60 Ž C. Yuan, A.J. Steckl, M.J. Loboda, Appl. Phys. Lett. 64 Ž J.-H. Boo, K.-S. Yu, M. Lee, Y. Kim, Appl. Phys. Lett. 66 Ž K.-W. Lee, K.-S. Yu, J.-H. Boo et al., J. Electrochem. Soc. 144 Ž K.-W. Lee, K.-S. Yu, Y. Kim, J. Crystal Growth 179 Ž A. Ishizaka, Y. Shiraki, J. Electrochem. Soc. 133 Ž A.T.S. Wee, Z.C. Feng, H.H. Hng et al., Appl. Surf. Sci. 81 Ž F. Bozso, J.T. Yates Jr., W.J. Choyke, L. Muchlhoff, J. Appl. Phys. 57 Ž