AD-A259 247 lI[I I """lll II Annual Report I I I Low Temperature Deposition and Characterization of N- and P-Type Silicon Carbide Thin Films 3 ]and Associated Ohmic and Schottky Contacts I Supported under Grant #N00014-92-J-1500 Office of the Chief of Naval Research Report for the period January 1, 1992-December 31, 1992 I R. F. Davis and R. J. Nemanich* R. S. Kern, R. Patterson, L. B. Rowland, L. Spellman-Porter, and S. Tanaka Materials Science and Engineering Department II *Department of Physics North Carolina State University 5] Raleigh, NC 27695 I Ao5 1993 SD N December, 1992 \IUIIU!ll 1 93-00255 kDi. -(cid:127)i REPORT DOCUMENTATION PAGE oI.S No 0704 0188 ,ne0.0--t0o og" " we"g..f~ge Iwg~.#M . % 0'9d. .....q I..e .@$ . .'.'.jg W..,.I.: ,J. oU~.I.s.q ~~, c.ep r 4 .0 it.. .)IUi I. AGENCY USE ONLY gLeev# OafnA) I. RJaEPnOuRaTr yD,A T1E9 92 I 3A. nREnPuORaTl- JTaYPnEu aArNyD D1A, T1E9S 9C2O-VDEeREcDember 31, 1992 4. TITLE AND SUBTITLE S. FUNDING NUMBERS Low Temperature Deposition and Characterization of N- and P-Type Silicon Carbide Thin Films and Associated Ohmic and R&T: sic0002--01 Schottky Contacts 1261 6. AUTHORtS) N00179 Robert F. Davis N66005 4B855 7. PERFORMING ORkmANIZATION NAME(S) AND AOORESS(ES) B. PERFORMING ORGANIZATION REPORT NUMBER North Carolina State University Hillsborough Street #N00014-92-J-1500 Raleigh, NC 27695 1. SPONSORINGi MONITORING AGENCY NAME(S) AND ADDRESS(ES) 10. SPONSORING. MONITORING Sponsoring: ONR, 800 N. Quincy, Arlington, VA 22217 AGENCY REPORT NUMBER Monitoring: Office of Naval Research Resider, N66005 The Ohio State Univ. Research Center 1314 Kinnear Road Columbus, OH 43212-1194 11. SUPPLEMENTARY NOTES tIa. OISTRIBUTION , AVAILABILITY STATEMENT t2b. OISTRIBUTION CODE Approved for Public Release-Distribution Unlimited IJ. ABSTRACT (Maemum1O0words) Single-crystal epitaxial films of cubic P(3C)-SiC(1 11) and AIN(0001) have been deposited on a(6H)-SiC(0001) substrates oriented 3-4* towards [1120] at 1050"C via gas-source molecular beam epitaxy using disilane (Si2H6), ethylene (C2H4), thermal evaporation of Al and activated N species derived from an ECR plasma. High resolution transmission electron microscopy revealed that the nucleation and growth of the 13(3C)-SiC regions occurred primarily on terraces between closely spaced steps. Pseudomorphic bilayer structures containing J0(3C)-SiC and 2H-AIN have been grown under the same conditions for the first time. HREED and cross-sectional HRTEM showed all layers to be monocrystalline. Initial high temperature chemical interdiffusion studies between SiC and AIN show that all components diffuse very slowly across the interface. AHRTEM and SAS are being used to determine the concentration profiles. Thin film solid solutions of AIN and SiC have been deposited using similar techniques and conditions as the individual compounds. Metal contacts of Ti, Pt and Hf deposited at RT on n-type a(6H)-SiC(0001) exhibit rectifying behavior with ideality factors between 1.01 and 1.09. The Pt and Hf contacts had leakage currents of 5x10"8 A/cm2 at -10V. Values of barrier heights for all contacts were within a few tenths of 1.0eV which is indicative that the Fermi level is pinned at the SiC surface. 14. SUBJECT TiRMS 15. NUMBER OF PAGES films, SiC, AIN, gas source molecular beam epitaxy, transmission electron 62 microscopy, chemical interdiffusion, metal contacts, Ti, Pt, Hf, ideality factors, 16. PRICE CODE Fermi level pinning It. SECURITY CLASSIFICATION If SECURITY CLASSIFICATION 19. S.CURITY CLASSIFICATION 20 LIMITATION OF ABSTRACT OF RE1PORT OF THIS PAGE Of ABSTRACT UNCLAS UNCLAS UNCLAS SAR j NN ,SA4j. /1113-'Ui) *." ..'0,(cid:127) "- , l(cid:127)1 '.. * : I Table of Contents I I. Introduction I II. Gas-Source Molecular Beam Epitaxy of Monocrystalline 6 Beta-SiC on Vicinal Alpha(6H)-SiC III. Chemical and Electrical Mechanisms in Titanium, 16 Platinum, and Hafnium Contacts to Alpha(6H) Silicon Carbide IV. Epitaxial Growth of AIN by Plasma-Assisted 25 Gas-Source Molecular Beam Epitaxy V. Aluminum Nitride/Silicon Carbide Multilayer 36 Heterostructure Produced by Plasma-Assisted Gas-Source Molecular Beam Epitaxy VI. Determination of the Diffusivity of the Si, C, Al, 46 and N at the Interface of the SiC-AIN Diffusion Couple VII. Solid Solutions of AIN and SiC Grown by 56 Plasma-Assisted, Gas-Source Molecular Beam Epitaxy VIII. Distribution List 62 I I I BaOsston Par (cid:127) NTIS Gft&I DTIC TAB Unannounoed [ Just i ati onC, By I (cid:127)etribut ton/ Availability Codes _S IL blot 1speoi.. I I. Introduction Silicon carbide (SiC) is a wide bandgap material that exhibits polytypism, a one- dimensional polymorphism arising from the various possible stacking sequences of the silicon and carbon layers. The lone cubic polytype, P3-SiC, crystallizes in the zincblende structure and is commonly referred to as 3C-SiC. In addition, there are also approximately 250 other rhombohedral and hexagonal polytypes [1] that are all classed under the heading of a-SiC. The most common of the a-SiC polytypes is 6H-SiC, where the 6 refers to the number of Si/C bilayers along the closest packed direction in the unit cell and the H indicates that the crystal structure is hexagonal. Beta (3C)-SiC is of considerable interest for electronic applications that utilize its attractive physical and electronic properties such as wide bandgap (2.2 eV at 300K) [2], high breakdown electric field (2.5x106 V/cm) [3], high thermal conductivity (3.9 W/cm 0C) [41, high melting point (3103K at 30 atm) [5], high saturated drift velocity (2x107 m/s) [61, and small dielectric constant (9.7) [7]. Primarily due to its higher electron mobility than that of the hexagonal polytypes, such as 6H-SiC [8], 13-SiC remains preferable to hexagonal SiC for most device applications. Most 3C-SiC thin film growth to date has been performed on Si substrates. Large-area, crack-free, and relatively thick (up to 30 gm) epitaxial 3C-SiC thin films have been grown on Si (100) by exposing the Si substrate to a C-bearing gaseous species prior to further SiC growth [7, 9, 10]. However, these films exhibited large numbers of line and planar defects due to large lattice and thermal mismatches between SiC and Si. One particular type of planar defect, the inversion domain boundary (1DB), was eliminated with the use of Si (100) substrates cut 2*-4* toward [011] [11-13]. Growth on Si substrates has allowed much understanding of SiC growth processes and device development to occur, but the large thermal and lattice mismatches between SiC and Si hamper further development using Si substrates. As a result, great effort has been made to develop methods for growth SiC single crystal substrates for homoepitaxial growth of SiC thin films. Since the 1950's, monocrystalline single crystals of 6H-SiC have been grown at using the Lely sublimation process [14]. However, nucleation was uncontrolled using this process and control of resultant polytypes was difficult. SiC single crystals inadvertently formed during the industrial Acheson process have also been used as substrates for SiC growth. However, neither these crystals or those formed using the Lely process are large enough for practical device applications. Recently, using a seeded sublimation-growth process, boules of single polytype 6H-SiC of > 1 inch diameter of much higher quality of that obtained using the Lely process have been grown. The use of single crystals of the 6H polytype cut from these boules has given a significant boost to SiC device development. SiC epitaxial thin film growth on hexagonal SiC substrates has been reported since the 1960's. The use of nominally on-axis SiC substrates has usually resulted in growth of 3C-SiC films. Films of 3C-SiC (111) grown by CVD have been formed on 6H-SiC substrates less than 1' off (0001) [15]. Films of 3C-SiC on 6H-SiC substrates have typically had much lower defect densities than those grown on Si substrates. The major defects present in 3C-SiC/6H- SiC films have been double positioning boundaries (DPB) [16]. Despite the presence of DPBs, the resultant material was of sufficient quality to further device development of SiC. The use of off-axis 6H-SiC (0001) substrates has resulted in growth of high-quality monocrystalline 6H-SiC layers with very low defect densities [17]. In addition, the use of more advanced deposition techniques, such as molecular beam epitaxy (MBE), has been reported for SiC in order to reduce the growth temperature and from about 1400-1500'C on 6H-SiC substrates. Si and C electron-beam sources have been used to I epitaxially deposit SiC on 6H-SiC (0001) at temperatures of 1150°C [18]. Ion-beam deposition of epitaxial 3C-SiC on 6H-SiC has also been obtained at the temperature of 750'C using mass- separated ion beams of 30Si+ and 13C+ [19]. Aluminum nitride (AIN) is also of particular interest at this time because of its very large bandgap. It is the only intermediate phase in the Al-N system and normally forms in the wurtzite (2H-A1N) structure. Most current uses of AIN center on its mechanical properties, such as high hardness (9 on Mohs scale), chemical stability, and decomposition temperature of Sabout 2000*C [201. Properties such as high electrical resistivity (typically > 1013 (2-cm), high thermal conductivity (3.2 W/cm K) [21], and low dielectric constant (e - 9.0) make it useful as a potential substrate material for semiconductor devices as well as for heat sinks. The wurtzite form has a bandgap of 6.28 eV [22] and is a direct transition, thus it is of great interest for optoelectronic applications in the ultraviolet region. Because of the difference in bandgaps (2.28 eV for 3C-SiC and 6.28 eV for 2H-A1N) between the materials, a considerable range of wide bandgap materials, made with these I materials, should be possible. Two procedures for bandgap engineering are solid solutions and multilayers. A particularly important factor is that the two materials have a lattice mismatch of less than one percent. Research in ceramic systems suggests that complete solid solubility of AIN in SiC may I exist [23]. Solid solutions of the wurtzite crystal structure should have Eg from 3.33 eV to 6.28 eV at 0 K. Although it has not been measured, the bandgap of cubic AIN has been estimated to be around 5.11 eV at absolute zero and is believed to be indirect [24]. Cubic solid solutions should thus have Eg from 2.28 eV to roughly 5.11 eV at 0 K and would be indirect at all compositions if theory holds true. I Because of their similarity in structure and close lattice and thermal match, AiN-SiC heterostructures are feasible for electronic and optoelectronic devices in the blue and infrared region. Monocrystalline AIN layers have been formed by CVD on SiC substrates [25] and SiC I 2 I I layers have been formed on AIN substrates formed by AIN sputtering on single crystal W [26]. In addition, theory on electronic structure and bonding at SiC/AIN interfaces [24] exists and critical layer thicknesses for misfit dislocation formation have been calculated for cubic AlN/SiC [27). Note that AIN (at least in the wurtzite structure) is a direct-gap material and SiC is an indirect gap material. Superlattices of these materials would have a different band structure than either constituent element. The Brillouin zone of a superlattice in the direction normal to the interfaces is reduced in size. This reduction in zone size relative to bulk semiconductors causes the superlattice bands to be "folded into" this new, smaller zone. This folding can cause certain superlattice states to occur at different points in k space than the corresponding bulk material states [28]. This can lead to direct transitions between materials which in the bulk form have indirect transitions. This has been demonstrated in the case of GaAsO.4P0.6/GaP and GaAs0.2P0.8/GaP superlattices, where both constituents are indirect in the bulk form [29]. Whether this is possible in the case of A1N/SiC is unknown, but very intriguing. It may be possible to obtain direct transitions throughout nearly the entire bandgap range with use of superlattices of AIN and SiC. Use of solid solutions in superlattices introduces additional degrees of freedom. For example, the bandgap can be varied independently of the lattice constant with proper choice of layer thickness and composition if superlattices of solid solutions of AIN and SiC were formed. Due to the potential applications of solid solutions and superlattice structures of these two materials, an MBE/ALE system was commissioned, designed, and constructed for growth of the individual compounds of SiC and AIN, as well as solid solutions and heterostructures of these two materials. Dithisimal studies concerned with the kinetics and mechanisms of mass transport of Si, C, Al and N at the SiC/AIN interface are also being conducted in tandem with the deposition investigations. A very important additional goal of this research is to understand what controls the contact electrical characteristics of specific metals to n-type 6H-SiC and to use this information to form good ohmic and Schottky contacts. A list of five metals to be studied, which consists of Ti, Pt, Hf, Co, and Sr, was created at the beginning of this research project. The selection process began by taking the simplest case, an ideal contact which behaves according to Schottky-Mott theory. This theory proposes that when an intimate metal-semiconductor contact is made the Fermi levels align, creating an energy barrier equal to the difference between the workfunction of the metal and the electron affinity of the semiconductor. It is the height of this barrier which determines how the contact will behave; for ohmic contacts it is desirable to have either no barrier or a negative barrier to electron flow, while for a good Schottky contact a large barrier is desired. Although metals were chosen optimistically, i.e. on the basis that they will form ideal contacts, some evidence exists that the contact properties will be more complicated. J. Pelletier 3 I I et al. [30] have reported Fermi level pinning in 6H-SiC due to intrinsic surface states, suggesting little dependence of barrier height on the workfunction of the metal. In addition, L. J. Brilnson [31, 32] predicts the pinning rate to be higher for more covalently bonded materials. Other complications may arise if the surface is not chemically pristine. A major part of this project will be devoted to determining whether the contacts behave at all ideally, and if not, whether the Fermi level is pinned by intrinsic or extrinsic effects. Along with examining the barriers of the pure metal contacts, the chemistry upon annealing will be studied and correlated with the resulting electrical behavior. The electrical behavior will be quantified both macroscopically in terms of current-voltage characteristics and microscopically in terms of barrier height. Identification of the phases formed will present the opportunity to attribute the electrical characteristics to the new phase in contact with silicon carbide. Within this reporting period, 3C-SiC and AIN thin films were grown on single crystal wafers of 6IH-SiC cut 3o-40 off (0001) towards [ 1120]. The crystalline quality of these films as determined using reflection high-energy electron diffraction (RHEED). Surface morphology was examined using scanning electron microscopy (SEM). Defects and interfaces were characterized using high-resolution transmission electron microscopy (HRTEM). From this analysis, reasons are discussed for the formation of 3C-SiC despite the presence of large numbers of surface steps on off-axis substrates. Research concerned with the growth of (AIN)x(SiC)I.x solid solutions and AIN/SiC heterostructures has also been conducted as well as chemical interdiffusion between these phases. Lastly, the results of the investigations concerned with the use of Ti, Pt and Hf metals as rectifying contacts on n-type 6H-SiC are described. Several techniques have been and are being implemented to measure barrier heights. Electrical measurements will be presented, along with some observed trends. The experimental procedures, results, discussion of these results, conclusions and plans for future efforts for each of the topics noted above are presented in the following sections. Each of these sections is self-contained with its own figures, tables and references. REFERENCES 1 . G. R. Fisher and P. Barnes, Philos. Mag. B 61, 217 (1990). 2. H. P. Philipp and E. A. Taft, in Silicon Carbide,A High Temperature Semiconductor, edited by J. R. O'Connor and J. Smiltens (Pergamon, New York, 1960), p. 371. 3. W. von Muench and I. Pfaffender, J. Appl. Phys. 48, 4831 (1977). 4. E. A. Bergemeister, W. von Muench, and E. Pettenpaul, J. Appl. Phys. 50, 5790 (1974). 5. R. I. Skace and G. A. Slack, in Silicon Carbide,A High Temperature Semiconductor, edited by J. R. O'Connor and J. Smiltens (Pergamon, New York, 1960), p. 24. 6. W. von Muench and E. Pettenpaul, J. Appl. Phys. 48, 4823 (1977). 7. S. Nishino, Y. Hazuki, H. Matsunami, and T. Tanaka, J. Electrochem Soc. 127, 2674 (1980). 4 8. P. Das and K. Ferry, Solid State Electronics, 19, 851 (1976). 9. K. Sasaki, E. Sakuma, S. Misawa, S. Yoshida, and S. Gonda, Appl. Phys. Lett. 45, 72 (1984). 10. P. Liaw and R. F. Davis, J. Electrochem. Soc. 132, 642 (1985). 11. K. Shibahara, S. Nishino, and H. Matsunami, J. Cryst. Growth 78, 538 (1986). 12. J. A. Powell, L. G. Matus, M. A. Kuczmarski, C. M. Chorey, T. T. Cheng, and P. Pirouz, Appl. Phys. Lett. 51, 823 (1987). 13. H. S. Kong, Y. C. Wang, J. T. Glass, and R. F. Davis, J. Mater. Res 3, 521 (1988). 14. J. A. Lely, Ber. Deut. Keram. Ges. 32, 229 (1955).- 15. H. S. Kong, J. T. Glass, and R. F. Davis, Appl. Phys. Lett. 49, 1074 (1986). 16. H. S. Kong, B. L. Jiang, J. T. Glass, G. A. Rozgonyi, and K. L. More, J. Appl. Phys. 63, 2645 (1988). 17. H. S. Kong, J. T. Glass, and R. F. Davis, J. Appl. Phys. 64, 2672 (1988). 18. S. Kaneda, Y. Sakamoto, T. Mihara, and T. Tanaka, J. Cryst. Growth 81, 536 (1987). 19. S. P. Withrow, K. L. More, R. A. Zuhr, and T. E. Haynes, Vacuum 39, 1065 (1990). 20. C. F. Cline and J. S. Kahn, J. Electrochem. Soc, 110, 773 (1963). 21. G. A. Slack, J. Phys. Chem. Solids 34, 321 (1973). 22. W. M. Yim, E. J. Stofko, P. J. Zanzucci, J. I. Pankove, M. Ettenberg, and S. L. Gilbert, J. Appl. Phys. 44, 292 (1973). 23. See, for example, R. Ruh and A. Zangvil, J. Am. Ceram. Soc. 65, 260 (1982). 24. W. R. L. Lambrecht and B. Segall, Phys. Rev. B 43, 7070 (1991). 25. T. L. Chu, D. W. Ing, and A. J. Norieka, Solid-State Electron. 10, 1023 (1967). 26. R. F. Rutz and J. J. Cuomo, in Silicon Carbide-1973, ed. by R. C. Marshall, J. W. Faust, Jr., and C. E. Ryan, Univ. of South Carolina Press, Columbia, p. 72 (1974). 27. M. E. Sherwin and T. J. Drummond, J. Appl. Phys. 69, 8423 (1991). 28. G. C. Osbourn, J. Vac. Sci. Technol. B 1, 379 (1983). 29. P. L. Gourley, R. M. Biefeld, G. C. Osbourn, and I. J. Fritz, Proceedings of 1982 Int'l Symposium on GaAs and Related Compounds (Institute of Physics, Berkshire, 1983), p. 248. 30. J. Pelletier, D. Gervais, and C. Pomot, J. Appl. 55, 994 (1984). 31. L. J. Brillson, Phys. Rev. B, 18, 2431 (1978). 32. L. J. Brillson, Surf. Sci. Rep., 2, 123 (1982). I I I I I I II. Gas-Source Molecular Beam Epitaxy of Monocrystalline Beta-SiC on Vicinal Alpha(6H)-SiC A Communication Submitted for Consideration for Publication to The Journal of Materials Research By L. B. Rowland*, S. Tanaka, R. S. Kern and Robert F. Davis Department of Materials Science and Engineering North Carolina State University Box 7907 Raleigh, North Carolina 27695-7907 November, 1992 ABSTRACT Single-crystal epitaxial films of cubic 03(3C)-SiC(1 11) have been deposited on hexagonal a(6H)-SiC(0001) substrates oriented 3-4* towards [1120] at 1050-1250*C via gas-source molecular beam epitaxy using disilane (Si2H6) and ethylene (C2H4). High resolution transmission electron microscopy revealed that the nucleation and growth of the i0(3C)-SiC regions occurred primarily on terraces between closely spaced steps because of reduced rates of surface migration at the low growth temperatures. Double positioning boundaries were observed at the intersections of these regions. *Present Address Naval Research Laboratory, Code 6861 4555 Overlook Av., SW Washington, DC 20375-5320 6 Polytypes are special one-dimensional polymorphs which differ only in the stacking sequence along the closest-packed direction. Silicon carbide occurs in one cubic (zincblende) polytype referred to as 3C- or 3-SiC, where the 3 refers to the number of planes in the periodic sequence. The hexagonal (wurtzite) polytype also exists in this material. Both polytypes occur in more complex, intermixed forms yielding a wider range of ordered, larger period, hexagonal or rhombohedral structures of which 6H is the most common. All of these noncubic polytypes are known collectively as a-SiC. The growth of 3C- and 6H-SiC thin films has been achieved primarily via chemical vapor deposition (CVD) (see Ref. #1 for a review of this research). Monocrystalline Si(100) wafers have been the principal substrate of choice for the deposition of 3C. It is now a common first step to transform the surface region of these wafers to 13-SiC by reaction with a C-containing gas to reduce the effects of the large mismatches in lattice parameters (-20%) and coefficients of thermal expansion (-10%). The epitaxial growth of SiC films on a(6H)-SiC(0001) substrates via CVD has been reported for three decadesI. Single phase P-SiC films result2,3 when the [0001] direction of the 6H wafer is oriented off-axis <10. The primary defects in these films are double positioning boundaries (DPB)4. The use of vicinal 6H-SiC(0001) substrates cut 3-40 towards [1120] have resulted in high-quality monocrystalline 6H-SiC layers with low defect densities5,6, including the absence of DPBs. Solid- and gas-source (GS) molecular beam epitaxy (MBE) techniques have also been employed for deposition of SiC films7,8. Kaneda et al.7 used on-axis a(6H)-SiC(0001) substrates and electron-beam evaporated Si and C sources. Epitaxial 3C-SiC(1 11) films were obtained at particular Si-to-C flux ratios in the temperature range of 1150-1400*C, as determined by reflection high-energy electron diffraction (RHEED). No information was given by these authors regarding either the microstructure or the type of defects present in these films. By contrast, Yoshinobu et al.8 employed the periodic introduction of Si2H6 and C H to 2 2 achieve 3C-SiC growth on vicinal 6H-SiC(000I) and 6H-SiC(OI 14(cid:127) substrates at 850-11600C. Films grown on vicinal 6H-SiC (0001) contained DPBs while those grown on 6H-SiC (0114) were free of these defect. Smooth films were obtained at the lowest growth rates used in the study (< 0.01 jrn/hr). In the present research, the 3C-SiC films were grown via GSMBE between 1050 and 1250*C on a(6H)-SiC(0001) wafers oriented 3--4 off [0001] towards [1120] and produced by Cree Research, Inc.using a seeded Lely sublimation method. The MBE growth system has been described previously9.Each wafer was sequentially cleaned prior to growth using a 10% HF etch at room temperature for 5 min, rinsed in DI water for 2 min and heated in the MBE chamber for 5 min at the growth temperature to achieve the desorption and decomposition of any remaining hydrocarbon species and native oxide, respectively. The source gases of Si I- 2 6 7