and heteroepitaxial growth of SiC on - Springer Link

Journal 206 of Electronic Materials, Vol. 28, No. 3, 1999

Paper Fissel, Pfennighaus, Kaiser, Schröter,Special and Issue Richter

Mechanisms of Homo- and Heteroepitaxial Growth of SiC on α-SiC(0001) by Solid-Source Molecular Beam Epitaxy A. FISSEL, K. PFENNIGHAUS, U. KAISER, B. SCHRÖTER, and W. RICHTER Institut für Festkörperphysik, Friedrich-Schiller-Universität Jena, Max-Wien-Platz 1, D-07743 Jena, Germany

Epitaxial growth of SiC on hexagonal (or α)-SiC(0001) has been performed by means of solid-source molecular beam epitaxy (MBE). The solid-source MBE growth conditions have been analyzed concerning the supersaturation and the excess phase formation of silicon and carbon. In general, our results demonstrate that control of the Si/C ratio and the supersaturation (S) is essential for the growth mode and the kind of polytype grown. Low temperature (T1450K under near surface equilibrium conditions, different growth modes, and conditions stabilizing the growth of certain polytypes have been found. With a step decrease of S, a step-flow growth mode of both 4Hand 6H-SiC was obtained and, for the first time in case of epitaxial SiC growth, a homogeneous nucleation of α-SiC at more C-rich conditions has been realized. Conditions stabilizing the growth of certain polytypes have been estimated by thermodynamic calculations considering the influence of polytype structure on the supersaturation and the surface energy. Based on these results, we have demonstrated the growth of a double-heterostructure by firstly growing a 3C-SiC layer on 4H-SiC(0001) at low temperature and a subsequent growth of 4H-SiC under near surface equilibrium conditions on a C-stabilized surface on top of this layer. Key words: Growth conditions, growth mechanisms, heterostructures, molecular beam epitaxy (MBE), polytypism, silicon carbide

INTRODUCTION SiC is a wide-band-gap semiconductor material of great technological interest for devices operating at high temperatures, high power, high frequency and in harsh environments.1,2 One of the most interesting properties of SiC is the occurrence of different structures (polytypes) in this material with different physical properties. Therefore, it may be possible to build devices from heterostructures consisting of one semiconducting material with confined electrons in a twodimensional (2D) gas3 only by a different stacking of SiC monolayers in [0001] direction. In this context, thin epitaxial films of SiC on SiC(0001) substrates with definite layer structure are of interest. The molecular beam epitaxy (MBE) is a suitable method to prepare such structures because of the controlled deposition process within an atomic layer (Received August 26, 1998; accepted November 23, 1998) 206

range and the clean growth conditions. Another benefit, which also originates from the ultra high vacuum (UHV), is the possibility of the in situ evaluation of the growth process by means of reflection high-energy electron diffraction (RHEED). Up to now, MBE technique has become more and more attractive for the epitaxial growth of SiC.4–9 However, most of the work concerning the epitaxial growth of SiC has been focused on the chemical vapor deposition (CVD) technique.10–14 Recently, epitaxial growth of SiC by CVD has experienced a large progress and the SiC epitaxial layers have reached a quality suitable for device application [see for example the special issue of phys. stat. sol. (b) 202 (1997)]. The breakthrough has been realized by performing homoepitaxial SiC growth on vicinal (off-oriented) SiC(0001) surfaces by a so called “step-controlled” epitaxy process.15,16 In step-controlled epitaxy, the grown layer inherits the stacking order of the substrate steps on the vicinal surface through a step-flow

Mechanisms of Homo- and Heteroepitaxial Growth of SiC on α-SiC(0001) by Solid-Source MBE

Fig. 1. Arrhenius plot of equilibrium Si vapor pressure to stabilize the α-SiC(0001) surface against graphitization. The dotted line corresponds to an activation energy of 4.5 eV.

Fig. 2. Supersaturation S as a function of Si vapor pressure for a Csupply corresponding to pc = 2 × 10–5 Pa at T = 1600K and T = 1500K (inset).

growth mode. The critical growth conditions for the transition from the step-flow growth mode to the growth by 2D nucleation have recently been determined in dependence of the growth rate and the offangle for the CVD-growth on 6H-SiC(0001) at temperatures between 1200 and 1600°C by Kimoto and Matsunami.17 The growth of heteropolytypic structures, however, demands definite nucleation conditions. At first, therefore, conditions have to be established not only to realize a perfect layer-by-layer growth in the epitaxial process, but also conditions stabilizing the growth of a certain polytype. No nucleation of α-SiC has been reported so far for the epitaxial growth by MBE or CVD. Many papers have dealt with both thermodynamic and kinetic origins of polytype growth of SiC,18,19 but a conclusive theory is still missing. Most of these theories are based on the influence of impurities, the Si/C-ratio and the temperature. It is believed that 3CSiC is the stable structure in the nucleation stage. Consequently, nucleation far from the equilibrium,

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where a rearrangement of the growing film is not possible, results only in 3C-SiC growth.20 At high temperatures (or low supersaturations), different polytypes have been grown at certain conditions by means of sublimation. An important factor affecting the crystal polytype is the seed polarity. Because of the polarity in the [0001] direction, the SiC surface is covered by a outermost atomic layer of Si (Si-face) or C (C-face). On the (0001) Si-face always the substrate polytype has been reproduced, whereas on the (000 1) C-face under certain conditions 4H-SiC occurs.21,22 This was found regardless of the modification of the substrate. Stein at al.,21 therefore, suggested that the different surface energies of the C-face and the Si-face have a marked influence on the kind of polytype growing on this surface. Maltsev et al.23 suggest that the occurrence of the 4H-SiC on the Cface is due to the planar character of the bond between the two carbon atoms (sp2-hybridization), which breaks the original symmetry of the seed. Furthermore, from the thermodynamic aspect, 4H should always occur at more C-rich conditions.18,24 Moreover, 4H is preferentially grown in the presence of selected impurities like Sn, Ge, Pb in the vapor phase.24 Only in one case, 6HSiC was grown on the Si-face of 4H-SiC by means of sublimation growth (modified Lely method)25 at temperatures between 2000 and 2300°C and high growth rates in the range of 2.5 mm/h. Under the presence of supersaturated Si vapor always 3C-SiC has been grown despite of the high temperature of about 2000°C. In general, on the SiC(0001) Si-face, there is a greater tendency for the occurrence of 3C-SiC than on the (0001)C-face.25 We report about results of different growth studies of SiC on α-SiC(0001), necessary to realize multiheterostructures of different polytypes, such as 4H/ 3C/4H. Experiments have been performed between 1200 and 1600K by means of solid-source MBE. In this context, we demonstrate the specific role of surface superstructures occurring on SiC(0001) and, moreover, of an alternating supply of Si and C for the stabilization of the low-temperature epitaxial 3C-SiC growth via 2D nucleation. For the first time, nucleation of α-SiC has been realized in an epitaxial growth process. The mechanisms of formation and growth of different polytypes and, moreover, of observed growth features have been investigated and analyzed in the framework of growth parameters, such as temperature, supersaturation and Si/C ratio. EXPERIMENTAL CONDITIONS The SiC films were grown in a RIBER-EVA 32 MBE system. The source materials of high-purity Si and high-purity C were evaporated separately by means of electron beam guns and controlled by a massspectrometer based flux meter. More detailed experimental conditions are decribed elsewhere.26,27 As substrates were used: on-axis and 3–8.5° off-axis SiC(0001) wafers from Cree Research Inc. The MBE growth mode, surface morphology and structure were investigated by in situ reflection high-energy electron dif-

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fraction (RHEED), ex situ by transmission electron microscopy (TEM), atomic force microscopy (AFM) and scanning electron microscopy (SEM) investigations. The grown polytype has been identified by means of electron channelling patterns.28 Ex situ preparation conditions are to be found in Ref. 29. In situ, prior the epitaxial growth, surfaces has been prepared by in situ sublimation etching in a Si flux to remove surface imperfections. Whereas after ex situ preparation no ordered step structure was found, a well developed step structure has been obtained after a two hour in situ etching, with steps typically 2–6 monolayer (ML) in height. EXPERIMENTAL SOLID-SOURCE MBE GROWTH CONDITIONS AT HIGH TEMPERATURES In our growth experiments, different growth features and growth modes have been obtained in dependence on the growth conditions. Therefore, we have established first these growth conditions for the solidsource MBE concerning the supersaturation (S) and the prevention of any excess phase formation of carbon, which were only roughly known up to now.4 Recently, we have already established the low temperature growth conditions (T < 1300K) and have demonstrated that a stable layer-by-layer growth has been realized only by an atomic level controlled surface-stabilized growth using Si-rich surface superstructures.30–34 At higher temperatures (T > 1400K), the graphitization is already a serious problem in the solid-source MBE growth which can only be prevented if a certain excess of Si over SiC is present in the gas phase.4 In our investigations, the formation of C-rich ( 3 × 3 )R30°-superstructure has been detected before graphitization by the occurrence of superstructure spots in RHEED diffraction pattern. Therefore, the Si-flux stabilizing the surface against graphitization was determined as a function of temperature by the first occurence of the superstructure. The corresponding vapor pressure of Si has been calculated by the Hertz-Knudsen relationship:

where v and s refer to the vapor and solid phase, respectively. Moreover, the vapor phase composition above the SiC surface can be shifted between the SiCC equilibrium and the SiC-Si one36,37 depending on the growth environment. Comparing our results for the Si equilibrium vapor pressure above SiC to prevent the graphitization with data from other groups concerning the Si equilibrium vapor pressure, we can find agreement to values given by Karpov37 for the so called SiC-C equilibrium. This is consistent with the fact that in our investigation the Si vapor phase is in equilibrium with the C-determined superstructure formed on the SiC surface. The supersaturation S at the growing surface is mainly determined by the equilibrium partial pressures of SiC2 and Si 2C at the substrate temperatures used. Both partial pressures are a strong function of temperature, but also of the given Si vapor pressure over the surface in accordance with the reactions given above. That means that the supersaturation S, which is determined by the relation for the minority component carbon: S = p C/ (pSi2C +2pSiC2), becomes also a strong function of the Si overpressure (corresponding to the Si supply) and, consequently, S cannot be controlled independently concerning Si and C. That is demonstrated by the data presented in Fig. 2. These data have been calculated by the functions given by Glass et al.38 for the equilibrium partial pressures of SiC2 and Si 2C, and a carbon flux corresponding to a vapor pressure of 2 × 10–5 Pa, which compares to a growth rate of about 30 nm/h. A strong change in S is seen from these curves by varying the Si flux, which corresponds to a shift of growth conditions from the SiC-C to the SiC-Si equilibrium. At a certain Si flux, slightly above the value to stabilize the SiC surface against graphitization, a maximum in S occurs. This effect may be of importance if the equilibrium conditions change during the growth, for example, due to effects like a changing surface morphol-

F = pV/(2πmkT)1/2, where F is the flux, pV-vapor pressure of the component, m-molar weight of the specie, k-Boltzmann constant, and T is the substrate temperature. The results of these investigations are given in Fig 1. The slope in the Arrhenius plot corresponds to an activation energy of 4.5±0.3 eV. At the temperatures used and a Si flux stabilizing the surface against graphitization, the vapor over SiC mainly consists of the species Si, Si2C and SiC2,35 which we assume result from the following reactions at the SiC surface: (1) SiC(s) → Si(v) + C(s) (2) 2SiC(s) → SiC2 (v) + Si(v) (3) C(s) + 2Si(v) → Si2C(v) (4) SiC2(v) + 3Si(v) → 2Si2C(v)

Fig. 3. Supersaturation S as a function of temperature for different Si overpressures and a C-supply corresponding to pc = 3 × 10–5 Pa. The inset shows the range of low supersaturation. The dotted line indicates the upper limit to prevent Si clustering, whereas arrows in the inset indicate the lower limit to prevent graphitization.


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same conditions (temperature, growth rate) growth can take place either by a step flow growth mode or via nucleation. RESULTS OF GROWTH EXPERIMENTS AND DISCUSSION Low Temperature Growth

Fig. 4. Cross-sectional HRTEM image of a 3C layer grown on on-axis 6H-SiC(0001) at 1330K and a growth rate of 50 nm/h. The bright line on the right hand side is due to a double position boundary (DPB).

Fig. 5. AFM images of 3C-SiC layers grown on on-axis 6H-SiC(0001) at 1430K with continuous supply of Si and C (left hand side) and a discontinuous supply of C (right hand side).

ogy. A further increase in Si vapor pressure results in a drop of supersaturation down to an almost constant level. Under these conditions, a stable growth via step-flow should be possible on terraces with lengths short enough to prevent a SiC nucleation. However, the excess Si flux should be restricted to a value low enough to prevent the formation of Si clusters or droplets at defects or at step edges giving rise to 3CSiC growth by a subsequent carbonization process. On the other hand, the excess Si supply has to be high enough to prevent the graphitization which restricts a further lowering of supersaturation, too. These restrictions (or limits) are considered in Fig. 3 by the dotted line (upper limit) and by the arrows in the inset (lower limit), respectively, where the supersaturation is presented as a function of temperature for different Si vapor pressures and a fixed C-supply. The supersaturations given in Fig. 2 and Fig. 3 are calculated concerning a flat SiC surface. If there exist steps and the surface adatom diffusion is sufficient, the adatom density (or supersaturation) will be drastically reduced by fast incorporation of adatoms into step edges in accordance to the known relationship based on the BCF model:39 S = pC / (pSi2C + 2pSiC2) + [1 – pC/(pSi2C + 2pSiC2)]/cosh(λ0/2λS), where λ0 and λS are the terrace length and the average diffusion length of adatoms, respectively. That becomes important for SiC growth at lower temperatures, where, depending on the terrace length, at the

In our growth experiments, step controlled epitaxial growth has been realized on vicinal (3.5–8° offaxis) substrates cut toward the [1120] direction at temperatures down to 1300K using a growth rate of 50 nm/h. That means that the diffusion length of adatoms becomes comparable to the terrace length at these conditions and the supersaturation cannot reach a critical value giving rise to a nucleation. This is supported by experiments performed on on-axis substrates at the same conditions. In this case, always 3C-SiC has been grown via nucleation. The nucleation can be interpreted by a multilayer nucleation process as indicated by a TEM micrograph (Fig. 4) of a 3C-SiC layer grown on on-axis 6H-SiC at 1330K. In this picture, there is seen an incoherent twin boundary (or double position boundary (DPB)), which is the main defect occuring in epitaxial 3C-SiC layers. DPBs result from the nucleation and growth of 3C in different stacking sequences, such as ABC and ACB, on adjacent terraces.40 On the left hand side of this boundary, there is a terrace structure with terrace length in the range of 15 nm and steps 3- and 6bilayer in height. Because there were no steps on the substrate before, this terrace structure results from the nucleation process itself and implies that 3 bilayer high nuclei are preferred. This may result from the slightly different surface energy of the A, B, and C terminated terraces.41 Moreover, the observed terrace length corresponds well to a length which we would expect for a 3.5° off-axis 6H-SiC(0001) substrate (assuming 3 SiC-monolayer steps height), where a step flow growth mode occurs at these conditions. The mobility of adatoms was found to be drastically increased by a discontinuous source supply. This is demonstrated in experiments performed on on-axis 6H-SiC at T = 1430K and certain Si flux, but, with a discontinuous supply of carbon. The carbon flux has been switched by a shutter corresponding to deposit one monolayer in each cycle. Between each deposition step a short 5 s growth interruption was performed. In this way, which is known as a migration enhanced epitaxy (MEE), the epitaxial layer quality was significantly improved. As seen in Fig. 5, the size of the DPB domains increases up to some hundreds of microns, a value which is comparable to thick SiC layers grown by CVD at much higher temperatures.42 Furthermore, many steps can be seen within the domains indicating a step flow growth mode. XRD-measurements of the film demonstrate the good crystalline quality of the grown layer. No differences in the FWHM value in the Θ-2Θ-scan ( 1450K) growth on on-axis substrates (terrace length above 1 µm) in dependence of supersaturation (S) and the supply ratio. At high supersaturation (above 50), corresponding to temperatures below 1500K and a growth rate of 50 nm/h, 2D nucleation of 3C-SiC on terraces frequently occurs. Decreasing S by a factor of 2, which corresponds to T = 1500K, and pSi > 5 × 10–4 Pa, the SiC layers grow by a step flow growth mode, however, frequently triangular shaped inclusions of 3C were observed. Mostly, these triangles are correlated with defects. A formation mechanism for such triangular defects in 4H epitaxial layers has been proposed by Zhou et. al.46 based on the interaction of partial dislocations with surface steps. However, this model cannot explain the occurrence of these defects also in 6H epilayers, because it is restricted to 4H. Another more likely explanation is based on the 3C-SiC nucleation on triangular stacking faults, which were induced by substrate imperfections.47,48 The stacking faults suppress step-flow growth and promote the development of large terraces. The nucleation of 3C at

Fig. 6. SEM images of different growth figures obtained for the epitaxial growth of 4H on nominal on-axis 4H-SiC(0001) at temperatures between 1550 and 1600K and the following conditions: a) S < 10, b) S < 10 and more C-rich, c) S > 10 and C-rich; d) simple illustration of the growth feature of Fig. 6b indicating a nucleation process at step edges.


those terraces occurs as a result of higher supersaturation. Increasing the temperature to 1600K and keeping the other conditons fixed, no 3C inclusions were found within the layers (Fig. 6a). At more C-rich conditions, besides the usual step flow growth mode, however, sawtooth like structures were also found at step edges on larger terraces extending across the terraces. Such conditions correspond to a Si vapor pressure below 5 × 10–4 Pa of Fig. 2, which corresponds simultaneously to a slightly higher supersaturation. In many cases, nearly equal distances between the single teeth have been observed (Fig. 6b). For that reason, we propose a new growth mechanism based on an onedimensional (1D) homogeneous nucleation49,50 at step edges. This nucleation process can be explained as follows: At high temperatures, the adatoms have a high mobility and can reach existing steps rapidly. If there exists a high dangling bond density at the step there is a high probability to incorporate the adatoms directly into the step edge giving rise to a step-flow growth mode. Along close-packed step-edges, however, there is a lower probability to incorporate the adatoms and diffusion along the edges is enhanced. Because of the longer lifetime of adatoms at these edges, the adatom concentration increases rapidly and can reach the value necessary to form a nucleus at the step edge, as illustrated in Fig. 6d. This value is in fact lower than the value necessary to form 2D nuclei on terraces. The formed nuclei grow then rapidly in size across the terrace especially in [1120] direction, where the growth rate is 20% faster than in [1100] direction.51 The fastest part of the nucleus then can reach the next step edge. This may be also an additional mechanism for step bunching and can give an explanation for the enhanced step bunching under C-rich conditions in CVD.52 In growth experiments on 4H-SiC at C-rich conditions, as indicated by a permanent presence of the Crich superstructure, the layer surface was found to be relatively rough and, moreover, no step structures were visible (Fig. 6c). This demonstrates that under this condition, the layers are grown via nucleation. The layers grown under these conditions, however, were of 4H polytype. The vanishing of the step structure can be attributed partly to a sublimation etch process, because of the easier removement of atoms from kink positions and, moreover, to the nucleation process itself. The occurrence of 2D nucleation at this low supersaturation is a surprising fact and can only originate from a reduced mobility of adatoms on top of the C-rich surface and, moreover, by the presence of many nucleation sites due to the superstructure. Large vacancies have been reported to form C-rich superstructures,53 which can act as such nucleation centers. Furthermore, the energy of nucleus formation can be expected to be reduced if such nucleation centers are present at the surface. Growing SiC under conditions of a higher Si supply or lower temperatures on such C-rich surface always results in 3C-SiC films.


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is large at low supersaturation, the free energy ∆G* of formation of a critical nucleus in case of homogeneous nucleation is expressed by:54 ∆G* = 16πΩ 2σ3/3∆µ2 where σ is the surface free energy, Ω is the molecular volume, ∆µ = –kTln(p/p0) and p is the vapor pressure of carbon (corresponding to the supply), p0 is the equilibrium vapor pressure (p0 = pSi2C + 2 pSiC2). At higher supersaturation, corresponding to a smaller critical size, the free energy in case of the (111) or (0001) face, respectively, can be expressed by:54 ∆G* = b4 σ2/∆µ

Fig. 7. Cross-sectional HRTEM image of 4H/3C/ on-axis 4H-SiC double-heterostructure grown at 1300K (3C-SiC) and 1600K (Crich)(4H-SiC).

Fig. 8. Supersaturation S as a function of Si overpressure at 1600K and a C-supply corresponding to pc = 2 × 10–5 Pa for 4H-SiC and 3C-SiC.

In first experiments based on the investigations presented here, a thin five monolayer thick 3C-SiC layer has been grown on an on-axis 4H-SiC(0001) substrate at low temperature by an alternating supply of Si and C. On top of this layer again 4H-SiC has been grown at high temperature under C-rich conditions. Despite the fact that the layer structure was not homogeneous across the substrate, the layer shows many regions of 1 µm large areas consisting of a 4H/ 3C/4H-SiC double-heterostructure (Fig. 7). THERMODYNAMICAL CONSIDERATIONS OF POLYTYPE GROWTH The occurrence of 4H in the nucleation stage on 4HSiC and 3C-SiC at low supersaturation and more Crich conditions may be explained within the framework of classical nucleation theory by means of a reversed Ostwald’s step rule. After this rule, the formation of the metastable phase can take place at first if its nucleation probability is initially higher (or its nucleation energy is lower). Considering that the critical size of the nuclei

where b is the bonding length. Considering the surface free energy only, then in any case 3C-SiC should occur in the nucleation stage because of the lower surface energy σ3C/σ4H = (1742 × 10–7 J/cm2) /(1800 × 10–7 J/cm2),55,56 respectively. However, we have also to consider that the polytype structure exerts an influence on the partial pressure of vapor components and, consequently, on the supersaturation S and the Si/C ratio at the surface. It has been reported, that the vapor above 3C-SiC is enriched with silicon more than the vapor above 4HSiC.18 Furthermore, at the same temperature, the equilibrium vapor pressure over 4H should be lower than over 3C. This was already attributed to the different heat of formation (∆Hf298) of 3C-SiC (–62.8 kJ/mol) and 4H-SiC (–66.6 kJ/mol).18 In our consideration, this different heat of formation has been taken into account to explain the polytype growth more quantitatively. Considering the dissociation reactions given above, the equilibrium constants of reaction (1) and (2) (and, therefore, the heat of reactions) are affected by the different heat of formation of SiC. In case of 3C-SiC, the heat of reaction (1) and (2) should be lower by a factor of 3.8 kJ/mol (reaction 1) and 7.6 kJ/mol (reaction 2), respectively, in comparison to 4H-SiC. The equilibrium vapor pressures of Si and SiC2, therefore, should be higher above the 3C-SiC surface. To quantify the higher equilibrium vapor pressure of SiC2 in case of 3C-SiC, the SiC2 partial pressures calculated after the equation given by Glass et al.38 have been corrected by a factor of exp[(2*3.8 kJ/mol)/RT] (R- molar gas constant), in accordance to the lower heat of reaction (2). As demonstrated by the calculations presented in Fig. 8, this higher SiC2 equilibrium vapor pressure leads consequently to a lower supersaturation above the 3C-SiC surface during the growth in comparison to 4H-SiC, especially at more C-rich conditions. Furthermore, from the data presented in Fig. 8 and the surface energies given above, we have estimated the energy necessary to form a critical nucleus. In Fig. 9, the results are plotted as a function of Si vapor pressure. From this plot, it is clearly seen that at more C-rich conditions, the formation of 4H-SiC is prefered, whereas with a higher supply of silicon 3C-SiC becomes more stable in the nucleation stage. Varying

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the supersaturation by changing the C-supply at a fixed Si vapor pressure, the occurence of 4H-SiC is prefered at low supersaturation (Fig. 10). Both results agree very well to our experimental finding and results presented in literature.for example, 18, 23–25 Moreover, we have also checked the influence of a mechanochemical activation due to misfit stresses caused by different lattice constants and expansion coefficents of the substrate and the epitaxial layer. Such a misfit leads to a further increase of saturation vapor pressures p0 above, for example, the cubic SiCphase on hexagonal SiC due to the strain energy ε = (∆a)2γ/2, (p = p0exp(2ε/κT)),54 where γ is the bonding force constant amounting to 3.1 N/cm.35 However, because of the low misfit of (a3C – a4H) = 6 × 1013 m,57 the strain energy in our case is only in the range of 0.5 meV and, therefore, does not influence the equilibrium vapor significantly. In summary, the effect of polytype structure on the vapor pressure will overcompensate the influence of the surface energy up to a certain Si vapor pressure and supersaturation, respectively. Further increasing of both parameters results in a crossover to a higher nucleation probability of 3C-SiC. In general, 4H-SiC nucleation should occur always at C-rich conditions, independent on the supersaturation, and, moreover, at low supersaturation using a medium supply of Si. We assume that for 6H-SiC the situation is similar, however, the curves (Figs. 7–9) in this case should be closer to the 3C-SiC. This likely makes it more difficult to grow 6H-SiC via nucleation at low temperatures.

Fig. 9. Free energy of formation of critical nucleus as a function of Si overpressure. The data of supersaturation were derived from the data presented in Fig. 8. The inset shows the range of low Si overpressure.

CONCLUDING REMARKS Epitaxial growth of SiC on α-SiC(0001) in the temperature range between 1200 and 1600K has been performed and investigated by means of solid-source MBE. The solid-source MBE growth conditions have been analyzed concerning the supersaturation, the Si/C supply ratio and the excess phase formation of silicon and carbon. In summary, the heteroepitaxial growth of 3C-SiC on vicinal α-SiC has to be performed via nucleation at low temperatures (T