High-Temperature Nucleation of Cubic Silicon Carbide on (0001) Hexagonal-SiC Nominal Surfaces Laurence
Latu-Romain,†,‡
Didier
Chaussende,*,†
and Michel
Pons‡
Laboratoire des Mate´ riaux et du Ge´ nie Physique (UMR CNRS 5628), INPGrenoble-Minatec, 3 parVis Louis Ne´ el, BP257, 38016 Grenoble cedex 1, France, and Laboratoire de Thermodynamique et Physico-Chimie Me´ tallurgiques (UMR CNRS 5616), INPGrenoble-CNRS, BP 75, 38402 Saint Martin d’He` res, France
CRYSTAL GROWTH & DESIGN 2006 VOL. 6, NO. 12 2788-2794
ReceiVed July 4, 2006; ReVised Manuscript ReceiVed October 9, 2006
ABSTRACT: The development of 3C-SiC crystals from oriented hexagonal seed has always suffered from systematic twinning that appears during the nucleation step of the layer. To investigate the possibility to reduce or eliminate the incoherent twin boundaries at high temperature (for conditions close to bulk growth ones), we conducted an experimental study on 3C-SiC nucleation. A mechanism for the selection of one 3C-SiC orientation among the two possible is proposed. It is based on a strong interaction between the R-SiC substrate steps and the anisotropic lateral expansion of the β-SiC domains. This model is confirmed by cross-sectional high resolution transmission microscopy observations of the R-β interface. The mechanism is discussed with respect to the surface polarity (Si or C faces), the miscut angle, and the substrate polytype. 1. Introduction Silicon carbide (SiC) is considered the most promising wide band gap semiconductor for high power, high temperature, and high-frequency electronic devices. Among the huge number of SiC polytypes, the β-SiC (or 3C-SiC) form with the zinc blende structure is the only cubic one. All the others collectively denoted as R-SiC occur in the hexagonal or rhombohedral structure. Because of the cubic lattice isotropy and the high value of the electron mobility, the 3C-SiC presents some additional interest compared to the R-SiC polytypes, which should make it more and more considered in the forthcoming years. For example, it could be an excellent candidate for n-type inversion channel MOS applications.1 Another point of interest for the 3C-SiC polytype is the (100) orientation, which is the most realistic substrate for the stabilization of cubic nitrides as AlN, GaN, and their alloys.2 Despite its high potential, the availability of 3C wafers is still a challenging issue as no one has succeeded in growing 3C-SiC bulk single crystals. A few groups have attempted to grow by chemical vapor deposition (CVD) either self-nucleated 3C-SiC crystals on a graphite susceptor3 or large area freestanding 3C-SiC wafers by heteroepitaxy on “undulant-Si (001)” substrates.4 In the first case, crystals of very high structural quality were obtained,5 but the size, about several millimeters in diameter, was not compatible with the applications. In the second case, up to 6-in.-diameter free-standing wafers have been demonstrated, but the crystals still contain too many structural defects that are intrinsically related to the heteroepitaxial SiC/ Si system.6 Consequently, expected electronic performances could not be obtained.7 The reasons for the absence of bulk 3C-SiC crystals are twofold. The first one is the availability of 3C-SiC seeds. Although both the Acheson and the Lely processes provided a lot of 6H-SiC seeds for the development of hexagonal materials, β-SiC seeds were never obtained. The second one is the availability of an adapted growth process. * Corresponding author: Tel: +33 4 56 52 93 29. Fax: +33 4 56 52 93 01. E-mail:
[email protected]. † Laboratoire des Mate ´ riaux et du Ge´nie Physique (UMR CNRS 5628), INPGrenoble-Minatec. ‡ Laboratoire de Thermodynamique et Physico-Chimie Me ´ tallurgiques (UMR CNRS 5616), INPGrenoble-CNRS.
The seeded sublimation method (or physical vapor transport, PVT) is the most common technique for the growth of 6H- and 4H-SiC ingots.8-10 The high temperatures involved (>2000 °C) make the stabilization of the β-SiC polytype very difficult.11 To bypass the lack of 3C-SiC seeds, the most realistic way is to heteropitaxially grow oriented β-SiC crystals on commercially available (0001) R-SiC nominal substrates. Interfacial problems are then drastically reduced as the lattice mismatch is only 0.09% and the differential thermal expansion is close to zero. However, in this case twins are usually observed, coming from the two possible orientations of the cubic 3C-SiC axis on the hexagonal 6H-SiC basis.12 As a consequence this approach requires one to monitor first the heteropolytypic R-β transition and second the selection of one orientation among the two possible, i.e., elimination of twin boundaries (TB). Many authors have reported some attempts to the eliminate TB in the heteroepitaxial growth of 3C-SiC. They pointed out some parameters that strongly affect the TB density, such as seed polarity, misorientation, C/Si ratio in the gas phase, and temperature.13-15 One of the key points is the seed surface preparation. Substrate defects such as structural defects or subsurface damage are 3C nucleation centers.16,17 With an adapted pregrowth in-situ etching, it was possible to reduce the 3C nucleus density, the resulting TB density being reduced as well.18 The most mature works are those of Neudeck et al.19 who demonstrated the formation of TB free epilayers by a perfect control of first the formation of a step free (0001) 6HSiC surface and second the 2D nucleation of the 3C polytype on such a surface and then its lateral expansion by a pure layer by layer growth mode. However, those results are restricted to growth on mesa (1950 °C), the heteropolytypic R-β transition is not favored anymore (Figure 4), because of the first thermodynamics discussed above and second activation of the spiral grow mode (kinetics). The latter continuously supplies the surface with new steps. The replication of the substrate polytype is thus preferred through the step flow growth mode, whereas 2D nucleation of 3C-SiC is avoided. The β Orientation Selection. On a nominal (0001) surface, the probability to nucleate one 3C-SiC orientation or the other (twinned) is perfectly similar. It has been demonstrated by Neudeck et al.19 who obtained one orientation or the other on a set of step-free mesas. As 3C-SiC nucleation systematically proceeds via 2D nucleation, it should be similar even for a vicinal (0001) surface. The effect of stacking order just below the topmost bilayer, which has been extensively investigated and discussed by Chien et al.28 has not proven to be significant on the 3C orientation selection in our study. See for example Figure 11 where both 3C orientations are obtained on the same R-SiC terrace; this picture will be detailed in the following. It means that the selection mechanism for the 3C-SiC orientation does not take place during the nucleation step (formation of the 2D nuclei) but after, i.e., during the expansion of those nuclei. From all our results, selection occurs only for a short window of experimental conditions: between 1850 and 1950 °C for a miscut angle of 0.37° (Figure 4). Under such conditions, 3C domains develop with particular shapes, elongated along the substrate step edges (see Figures 3 and 6). Steps are thus playing an important role. The scenario for the 3C domain expansion is shown in Figure 9. If the terrace length is larger
Latu-Romain et al.
Figure 9. Effect of the substrate step on the 3C-SiC nucleus lateral expansion along the step edge. (A) Formation of a 2D nucleus on the substrate terrace. (B) Expansion of the nucleus toward all directions. (C) Directional expansion of the 3C nucleus along the substrate step edge.
Figure 10. According to the Wulff construction, faceting of the two twinned 3C-SiC domains leads to different lateral expansion. They could be (A) on the same terrace and (B) on a different terrace. The normal direction is the .
Figure 11. Cross-sectional HRTEM observation of the R-β interface. A very thin twinned 3C domain suggests that it is overgrown at the early stage of growth.
than the adatoms mean free path, a 2D nucleus can form on the step (Figure 9A). This nucleus can then develop toward all directions and reach the step edge (Figure 9B). At this stage, a strong potential sink appears as it consists of a particular surface kink. This kink could be considered as a triple interface between the β-SiC nucleus-R-SiC substrate step edge and vapor phase. This kink strongly enhanced the adatoms incorporation, and a fast lateral growth of the 3C domains is then observed (Figure 9C). If the step density is too low (θ ) 0.09°, for example), the 2D nucleus is too far from the step edge and its expansion is independent of the substrate step edge. The resulting 3C domains are unshaped (Figure 6a). This first model addresses the observed shape of the 3C-SiC domains. However, it cannot explain the selection of the orientation. For this, some additional points have to be considered. For instance, what happens when two nuclei of the two different orientations form on the same terrace (Figure 10A)? On the basis of the model presented by Wulfhekel et al.29 for the faceting of the 6H-SiC step edges, we could propose the following. Along the c-axis of the crystal, both 3C-SiC domains are of three-fold symmetry. Within each domain, the stacking is hcp, but from one domain to the other, the lattice is
High Temperature Nucleation of β-SiC on R-SiC
rotated by 60°. As a consequence, the close packed step edges change their nature going from one domain to the other if we consider the same direction. The character of steps running along the substrate low index directions alternates and shows either single or double dangling bonds per atoms.30 Then, depending on the substrate step orientation with respect to the 3C steps orientation, one 3C domain could expand much faster than the other. It is what we observed in Figure 6b (EBSD picture), where the 3C(2) domain is very small compared to 3C(1) on the same terrace. With the same idea, two twinned 3C-SiC domains on two terraces (one over the other, see Figure 10B) could expand laterally at different rates. For instance, if the upper 3C domain has two dangling bonds per atom at the step edge and the lower 3C domain only one in the step flow direction, then fast overlapping of the lower 3C domain is expected. This leads to a twin in the basal plane between those two domains (see Figure 11). Such observation has already been mentioned for the heteroepitaxial growth of 3C-SiC from the vapor-liquid-solid mechanism at low temperature.31 To sum up, the selection mechanism is based on a strong interaction between the R-SiC substrate steps and the anisotropic lateral expansion of the β-SiC domains. TEM observations of the interface for different samples for which selection of a 3C orientation occurred revealed that the microtwin (overlapped 3C domain) are always very thin, less than 1 µm. It means that if the selection mechanism is able to take place (adapted growth conditions and substrate orientation), it is very effective in the early stage of growth. After that, the elimination of incoherent twin becomes not very probable even for very large thicknesses. Polarity Influence. The effect of polarity on mechanisms involved in nucleation, crystal growth, doping incorporation, defects occurrences, and so on is extensively documented in the literature. The main parameter explaining the strong behavior difference between silicon and carbon polarities is their surface free energy, which is more than twice as high for the Si-face than for the C-face. As a consequence, growth front morphologies observed for similar process conditions are rather different. To decrease its high surface energy, the Si-face prefers to change its shape by enhanced bunching and faceting.32 This was clearly evidenced in Figure 1. Such differences in surface energy will affect both the R-β transition and the selection of the β-SiC orientations. Because of its lower surface energy, the free energy of formation of a critical 2D nucleus is lower on the carbon face.33 This could explain why we observed at 1800 °C some 3C-SiC formed on the C-face, whereas no 3C-SiC was nucleated on the Si-face (see Figures 4 and 8). However, when temperature increases, the increased adatoms mobility promotes the faceting of the Si-face. The formation of perfectly nominal facets thus favors the 2D nucleation on the Si-face. For the whole range of experimental conditions used, faceting of the carbon face has never been observed. The orientation selection mechanism, which is based on an anisotropic lateral growth rate (which sounds similar to faceting), could not take place on the C-face. As a consequence, no particular shape of the 3C domains has been noted (Figure 7), and both 3C-SiC orientations have a 1:1 proportion. Polytype Effect. Many experiments have been conducted on 4H-SiC substrates. Even if similar conclusions could be extracted, results were much more dispersed and hard to reproduce. The precise origin of this behavior still is unclear. 5. Conclusion With bulk growth in view, we have investigated at high temperature with the continuous feed-physical vapor transport
Crystal Growth & Design, Vol. 6, No. 12, 2006 2793
process, the nucleation of the β-SiC polytype on R-SiC {0001} nominal surfaces. First, the R-β heteropolytypic transition has been studied. Optimal conditions have been discussed with respect to thermodynamics and kinetics (surface growth mode). Second, the selection of one 3C-SiC orientation among the two possible, i.e. elimination of incoherent twin boundaries, has been presented. A model for this selection has been proposed. It is based on a strong interaction between the R-SiC substrate steps and the lateral expansion of the β-SiC domains. If several conditions are fulfilled at the same time (adapted substrate step density and orientation, growth conditions that allow faceting of the 3C domains...), the selection mechanism is likely to occur and is very effective within the first micrometer of growth. We believe that this model is general and should be true whatever the growth process (CVD, LPE, PVT). Finally, as the surface energy of the C-face is lower than the Si-face, such selection mechanism has not been observed. Acknowledgment. This work is in part supported by the DGA (French Ministry of Defence). The authors would like to thank particularly L. Rapenne, P. Chaudouet, and Dr. E. Pernot for their support and Drs. G. Ferro and E. K. Polychroniadis for fruitful discussions. References (1) Scho¨ner, A.; Bakowski, M.; Ericsson, P.; Stro¨mberg, H.; Nagasawa, H.; Abe, M. Mater. Sci. Forum 2005, 801, 483-485. (2) Founta, S.; Gogneau, N.; Martinez-Guerrero, E.; Ferro, G.; Monteil, Y.; Daudin, B.; Mariette, H. Mater. Sci. Forum 2004, 457-460, 1561-1564. (3) Gorin, S. N.; Ivanova, L. M. Phys. Status Solidi B 1997, 202, 221245. (4) Nagasawa, H.; Yagi, K.; Kawahara, T. J. Cryst. Growth 2002, 237239, 1244-1249. (5) Vetter, W. M.; Dudley, M. J. Cryst. Growth 2004, 260, 201-208. (6) Polychroniadis, E.; Syvajarvi, M.; Yakimova, R.; Stoemenos, J. J. Cryst. Growth 2004, 263, 68-75. (7) Krieger, M.; Pensl, G.; Bakowski, M.; Schoener, A.; Nagasawa, H.; Abe, M. Mater. Sci. Forum 2005, 483-485, 441-444. (8) Nishizawa, S. I.; Kato, T.; Kitou, Y.; Oyanagi, N.; Hirose, F.; Yamaguchi, H.; Bahng, W.; Arai, K. Mater. Sci. Forum 2004, 457460, 29-34. (9) Muller, S. G.; Glass, R. C.; Hobgood, H. M.; Tsvetkov, V. F.; Brady, M.; Henshall, D.; Malta, D.; Singh, R.; Palmour, J.; Carter, C. H. Mater. Sci. Eng. B 2001, B80, 327-331. (10) Ohtani, N.; Fujimoto, T.; Katsuno, M.; Aigo, T.; Yashiro, H. J. Cryst. Growth 2002, 237-239, 1180-1186. (11) Furusho, T.; Sasaki, M.; Ohshima, S.; Nishino, S. J. Cryst. Growth 2003, 249, 216-221. (12) Kong, H. S.; Jiang, B. L.; Glass, J. T.; Rozgonyi, G. A.; More, K. L. J. Appl. Phys. 1988, 63, 2645-2650. (13) Kong, H. S.; Glass, J. T.; Davis, R. F. J. Mater. Res. 1989, 4, 204214. (14) Nishino, K.; Kimoto, T.; Matsunami, H. Jpn. J. Appl. Phys. Part 1 1997, 36, 5202-5207. (15) Powell, J. A.; Larkin, D. J.; Matus, L. G.; Choyke, W. J.; Bradshaw, J. L.; Henderson, L.; Yoganathan, M.; Yang, J.; Pirouz, P. Appl. Phys. Lett. 1990, 56, 1353-1355. (16) Hallin, C.; Konstantinov, A. O.; Pecz, B.; Kordina, O.; Janzen, E. Diamond Relat. Mater. 1997, 6, 1297-1300. (17) Powell, J. A.; Petit, J. B.; Edgar, J. H.; Jenkins, I. G.; Matus, L. G.; Yang, J. W.; Pirouz, P.; Choyke, W. J.; Clemen, L.; Yoganathan, M. Appl. Phys. Lett. 1991, 59, 333-335. (18) Xie, Z. Y.; Edgar, J. H.; Burkland, B. K.; George, J. T.; Chaudhuri, J. J. Cryst. Growth 2001, 224, 235-243. (19) Neudeck, P. G.; Powell, J. A. In Silicon Carbide; Springer: New York, 2004; pp 179-205. (20) Chaussende, D.; Ucar, M.; Auvray, L.; Baillet, F.; Pons, M.; Madar, R. Cryst. Growth Des. 2005, 5, 1539-1544. (21) NOVASiC, http://www.novasic.com. (22) Vicente, P.; Pernot, E.; Chaussende, D.; Camassel, J. Mater. Sci. Forum 2002, 389-393, 729-732. (23) Chaussende, D.; Chaudouet, P.; Auvray, L.; Pons, M.; Madar, R. Mater. Sci. Forum 2004, 457-460, 387-390.
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