CRYSTAL GROWTH & DESIGN
Vapor–Liquid–Solid Growth of 3C-SiC on r-SiC Substrates. 1. Growth Mechanism
2008 VOL. 8, NO. 3 1044–1050
Maher Soueidan,*,†,‡ Gabriel Ferro,† Olivier Kim-Hak,† François Cauwet,† and Bilal Nsouli‡ Laboratoire des Multimateriaux et Interfaces, UMR-CNRS 5615, UniVersité Claude Bernard Lyon 1, 43 Bd du 11 noV. 1918, 69622 Villeurbanne, France, and Lebanese Atomic Energy Commission-CNRS, P.O. Box 11-8281, Riad El Solh 1107 2260 Beirut, Lebanon ReceiVed June 1, 2007; ReVised Manuscript ReceiVed NoVember 21, 2007
ABSTRACT: We report on the heteroepitaxial growth of 3C-SiC layers by a vapor–liquid–solid (VLS) mechanism on various R-SiC substrates, namely, on- and off-axis for both 4H- and 6H-SiC(0001), Si and C faces. The Si-Ge melts, in which the Si content was varied from 10 to 50 atom%, were fed by 3 sccm of propane. The growth temperature was varied from 1200 to 1600 °C. It was found that single domain 3C-SiC layers can be obtained on 6H-SiC off- and on-axis and 4H-SiC on-axis, while the other types of substrates gave twinned 3C-SiC materials. As a general rule, one has to increase temperature when decreasing the Si content of the melt to avoid twin formation. It was also found that twinned 3C-SiC layers form at low temperatures, while homoepitaxy is achieved at high temperatures. Some growth mechanisms are proposed to explain the possibility of achieving either homoepitaxial or 3C-SiC layers (twinned or twin free) by changing the growth conditions. Concerning the selection of one orientation of the 3C layer for twin elimination, correlation with the carbon solubility in the melt and surface characteristics of the R-SiC seed are discussed. Introduction In recent years, the research activity in silicon carbide (SiC) has considerably increased due to the need for electronic devices capable of operation at high power levels and high temperatures. In addition to its excellent physical and electronic properties, this semiconducting compound exhibits polymorphism along the (0001) axis called polytypism. Each polytype differs from one to another only in the stacking sequence of the C-Si bilayers (see ref 1 for more details on SiC material). The three most common SiC polytypes are the 3C (or β), 4H, and 6H (both R). The 3C polytype is the only one with a cubic (zinc-blende) structure. As a consequence, it has specific advantages over the hexagonal forms that make it desirable for device applications. For instance, the reduced density of near-interface traps may help increase the drift mobility,2,3 or its lower bandgap would be more appropriate for inversion channel MOS applications.4 Furthermore, it has been shown very recently that the stacking faults do not move during device operation in 3C material, whereas this is the case in 4H.5 The main problem for 3C based electronics is the absence of commercially available substrate of this polytype. Silicon was often used as a host substrate for 3C-SiC layers, but the grown material suffers from high amounts of crystalline defects and important curvature due to lattice and thermal expansion mismatch, respectively. On the other hand, when onaxis R-SiC(0001) substrates are used, such a mismatch is almost negligible so that higher quality layers are expected to be obtained. However, the material commonly grown by chemical vapor deposition (CVD) or sublimation contains a high density of twins (also called double positioning boundary or DPB) due to a 60° rotation of the initial 3C stacking on the (0001) hexagonal plane, each orientation having equivalent probability of nucleation. Elimination or partial reduction of the twin density is thus a difficult issue. The presence of screw dislocations in * Corresponding author. E-mail:
[email protected]. Tel: +33 4 72 43 82 31. Fax: +33 4 72 44 06 18. † Université Claude Bernard Lyon 1. ‡ Lebanese Atomic Energy Commission-CNRS.
Table 1. Influence of the Substrate Nature on the SiC Deposita
3C-SiC formation twin-free layer
4H, Si face
6H
on-axis 4° off 8° off
on-axis on-axis 3.5° off Si face C face Si face
yes yes
yes no
yes no
yes yes
yes no
yes yes
a The term “yes” does not mean that it is always positive since it often depends on other growth conditions such as temperature or melt composition.
the substrate generates also some polytype inclusion due to a higher step density around these dislocations, which favors substrate polytype replication by step-controlled epitaxy.6 Recently, twin-free (also called single domain) 3C-SiC layers were obtained using nonstandard growth techniques such as continuous-feed physical vapor transport (CF-PVT)7 or vapor– liquid–solid (VLS) mechanism in Si-Ge melt.8,9 These promising results need to be deepen to evaluate their potential for proposing an alternative to R-SiC based electronics. The nucleation step of 3C-SiC on R-SiC by a VLS mechanism using a Si-Ge melt has been studied in another paper.10 In this paper, we focus on the growth aspects by evaluating the layer characteristics (twin density, roughness, polytype) as a function of temperature, melt composition, and R-SiC substrate nature (polytype, polarity, and misorientation). We will show that twin-free material can be grown on a wide range of conditions. A model will be proposed to explain the transition between homoepitaxial and heteroepitaxial growth. Experimental Procedures The experimental setup is composed of a water-cooled, quartz vertical cold-wall reactor. The temperature of the RF-heated graphite susceptor is controlled by an optical pyrometer. To improve the thickness uniformity, a susceptor rotation of ∼60 rpm was used. The carrier was high purity Ar (5 slm), and the carbon source was a mixture of H2 (95%) and propane (5%). For all the experiments, the propane flux was 3 sccm. Various types of 4H- and 6H-SiC(0001) oriented substrates were used in this study: Si face on-axis and off-axis, 8° and 4° for 4H and 3.5° for 6H. C face substrates were also used but only for 6H polytype and on-axis orientation (see Table 1 for the listing of the substrates).
10.1021/cg070499+ CCC: $40.75 2008 American Chemical Society Published on Web 01/23/2008
Vapor–Liquid–Solid Growth of 3C-SiC
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Figure 1. AFM image showing the surface morphology of 3C-SiC island on 6H-SiC on-axis after a contact with the liquid (before VLS growth) (a) scan 20 × 20 µm2 and (b) scan 2 × 2 µm2. A pregrowth surface preparation of each seed was performed in the same way as described in a previous report.11 After that, graphite glue was used to stick the seed (8 × 8 to 10 × 10 mm2) at the center of a graphite crucible, which was also stuck on top of the susceptor. The crucible dimensions were 20 and 2 mm, respectively, for inner diameter and depth. Si (6N) and Ge (6N) pieces were stacked inside the crucible above the SiC seed. Two melt compositions were mainly studied, that is, Si25Ge75 and Si50Ge50, with some experiments down to 10 at% Si. After introduction of the assembly in the reactor and 30 min evacuation, 5 slm of Ar was flowed. The temperature was varied from 1200 to 1600 °C at which point propane was introduced to start the growth. Two growth procedures were used in this work. Procedure 1 is the simplest with only one temperature plateau when the temperature is lower than 1450 °C. Because of the high evaporation rate of the melts at high temperature, which prevents long growth time deposition, we also used a two-step process called procedure 2. In this procedure, a 5 min growth at high temperature (g1450 °C), for the nucleation step, was followed by a second stage of 55 min at a lower temperature (1250 °C) where evaporation is low. Propane was switched off during the temperature transition, which took 1 min, and added again in the reactor at low temperature. This two-step process does not affect the final result since the layer quality (3C or homoepitaxy, twin formation) depends essentially on the initial nucleation.10 After 1 h of growth, the melt was removed, as described in ref 11, before cooling down. The remaining Si-Ge alloy inside the crucible was eliminated by wet chemical etching in HF-HNO3 solution. The layers were mainly characterized by Normarski optical microscopy since the observation of the surface morphology is very often sufficient to determine the layer polytype and the presence of twins.8 Confirmation was also routinely obtained from micro-Raman spectroscopy using an Ar laser (488 nm) that was focused to form a spot of a few µm2. Electron backscattering diffraction (EBSD) was used on selected samples for orientation and phase mapping of the deposit. The denomination “twin-free” will refer to the entire sample surface.
Results One of the major results displayed in ref 10 was the nucleation of elongated 3C-SiC islands on the R-SiC surface (see Figure 1) simply during the initial heating ramp, when the seed is in contact with the liquid just after the melting. These islands form without any addition of propane in the reactor. The proposed mechanism for their formation involves a slight dissolution of the seed during the temperature ramp up followed by the precipitation under the form of SiC of the excess of C inside the melt. This precipitation of SiC from the liquid phase is supposed to be so fast that the nucleation occurs on the terraces and not at the step edges. The cubic polytype is thus obtained. The formation of an entire 3C layer after a complete growth run suggests that the main effect of propane addition, when the temperature plateau is reached, is to enlarge the initial islands at the expense of substrate polytype replication by stepcontrolled epitaxy.
Figure 2. Typical surface morphology of 3C-SiC layers grown by VLS in a Si25Ge75 melt at 1400 °C on (a) C face 6H-SiC on-axis, (b) Si face 4H-SiC 4° off, and (c) Si face 4H-SiC 8° off.8
As mentioned above, several kinds of R-SiC substrates were used. The main results obtained on these substrates are summarized in Table 1. One can see that 3C-SiC heteroepitaxial layers could be obtained on all these kinds of substrates, even on highly misoriented (8° off). For the substrate that always gave twinned 3C layers (namely, Si face 4H misoriented and C face 6H on-axis), the surface morphology is rather disturbed as can be seen in Figure 2. The twin density is very high, much higher than in the worst layers grown on the three other types of substrates giving single domain layers (namely 6H and 4H on-axis, and 6H 3.5° off, all of them being Si face). The worst case is the 4H 8° off substrate showing a surface morphology composed of 60° rotated triangles of a few micrometers. EBSD characterization performed on such a layer confirmed not only the two possible orientations of the 3C domains (Figure 3a) but also the presence of 4H inclusions (Figure 3b). These inclusions were difficult to localize and characterize by Raman spectroscopy. Because the most interesting aspect is the different substrates on which twin-free layers were grown, we will present here a detailed study of the growth conditions for each of them. The results are summarized in Figure 4. One can see that singledomain 3C-SiC layers can be obtained only with specific conditions of temperature and melt composition. These condition windows depend on the seed nature. As a general rule, one has to increase the temperature when decreasing the Si content of the melt to avoid twin formation. This trend is more pronounced
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Figure 3. EBSD phase mapping of the same layer as shown in Figure 2c; panels a and b represent the same area but with different color contrasts to enhance the presence of 4H inclusions.
on 6H substrates than on 4H ones. The experience shows that the surface morphology of these twin-free layers does not vary significantly with the growth conditions (temperature or melt composition) but rather depends more on the type of substrate used. An illustration of such differences is given in Figure 5. For example, using 6H-SiC 3.5° off, the surface is very rough with more jagged steps than on the layers grown on on-axis seeds. The surface morphology obtained on 4H-SiC on-axis is also clearly different since it is the smoothest one with no clear step bunching distinguishable by optical microscopy. The three graphs in Figure 4 show also that twinned 3C-SiC layers are formed at low temperature, whereas homoepitaxy is achieved at high temperature. The main difference between these graphs concerns the boundary conditions between single-domain 3C-SiC and homoepitaxy. While direct transition from singledomain 3C to homoepitaxy is observed with 6H-SiC 3.5°off substrate, in the other two cases the transition occurs with mixing of polytype or twinned 3C-SiC complete layers. The typical surface morphologies of homoepitaxial layers for these three different substrates are given in Figure 6. In contrast to the case of twin-free 3C layers, the morphology does not differ significantly with the seed nature. It is composed of large terraces separated by steps of several hundreds of nanometers in height. Note the 120° angle of the step tips which differ from the 60° tips obtained in the case of 3C-SiC layers. This difference allows a fast identification of the layer polytype, which is usually confirmed by Raman spectrometry. The typical surface morphology of the layers grown on 6HSiC and 4H-SiC on-axis substrates but in the transition domain between 3C-SiC and homoepitaxial layers are shown in Figure 7. One can see that in the case of the 6H-SiC substrate (Figure 7a), the layer presents two different morphologies, one with
Figure 4. Experimental conditions to obtain 3C-SiC single-domain layers as a function of temperature and melt composition, (a) on 6HSiC on-axis, (b) on 4H-SiC on-axis, and (c) on 6H-SiC 3.5° off. All these substrates are Si face.
large and parallel terraces and the other with jagged steps and 60° tips uniformly oriented. Micro-Raman analyses (Figure 7c) unambiguously show that the area with large and parallel terraces is homoepitaxial and the other area is 3C-SiC, probably single-domain. In the case of the 4H-SiC substrate, the surface morphology is highly disturbed due to the apparent presence of a dense network of twins. However, one cannot exclude the presence of homoepitaxial inclusions since it was detected on another highly disturbed 3C layer (Figures 2c and 3b). These inclusions may be too small to be evidenced by micro-Raman spectroscopy. Discussion For all the R-SiC substrates used in this study, homoepitaxial layers were always achieved at high temperatures. Since the growth of a 3C-SiC layer by VLS involves the formation of 3C-SiC islands during the initial contact with the liquid at low temperature, one can suggest that these islands are dissolved when reaching high temperature due to an increase of carbon solubility in the melt. In all cases also, the increase in Si content of the melt results in a decrease in the minimum temperature of appearance of the homoepitaxial layers. This can be easily explained by the increase of carbon solubility inside the melt due to Si content increase (pure Si dissolves more C than Ge does12).
Vapor–Liquid–Solid Growth of 3C-SiC
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Figure 5. Surface morphology of single-domain 3C-SiC layers grown by VLS on various substrates (a) 6H-SiC on-axis, (b) 6H-SiC 3.5° off, and (c) 4H-SiC on-axis.
Figure 6. Surface morphology of homoepitaxial layers grown by VLS on different Si face substrates: (a) 6H-SiC on-axis, (b) 6H-SiC 3.5° off, and (c) 4H-SiC on-axis.
The main difference between the results obtained on the various substrates lies on the transition between homoepitaxy and 3C single-domain conditions in Figure 4. We will first consider the cases of 6H substrates (on- and off-axis) and then try to see if the explanation can be extended to the case of 4H on-axis. When comparing the results obtained on these two different 6H substrates, two important parameters have to be taken into account: first, the geometrical aspects related to the substrate misorientation, that is, the variation of step density at the surface and the inclination of the basal plane, and second, the thermodynamic aspect related to temperature since homoepitaxy occurs at significantly lower temperatures in the case of an off-axis substrate. Concerning the thermodynamic aspects, as 3C-SiC layers were obtained even at temperatures as low as 1200 °C, the 3CSiC islands are most probably formed below this temperature during the initial heating ramp. If we consider now that the island nucleation occurs only once during the temperature increase, that is, below 1200 °C, these islands are thus at
the origin of both twinned and twin-free 3C layers (at low and moderate temperature), but they do not prevent the formation of homoepitaxial layers at higher temperatures. After formation of these islands, a further increase in temperature should have the only effect of dissolving these islands due to increased C solubility in the melt. The size of the islands should thus decrease when the growth temperature increases, until their disappearance at too high a temperature. If we now consider the geometrical aspects, the increase of the misorientation angle involves an increase not only of the step density but also of the vertical component of the lateral growth. Considering these various parameters, we can propose a growth model illustrated in Figure 8. It is mainly based on the competition between homoepitaxy at the seed surface and enlargement of the 3C-SiC islands upon propane addition into the reactor to start the VLS growth. At low temperatures, these islands are tall enough compared to the vertical component of the lateral growth to allow their expansion in all directions. The strong lateral growth, inherent to growth from the melt, will allow covering the entire surface with 3C-SiC, starting from
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Figure 7. Surface morphology of SiC epitaxial layers grown at 1450 °C (a) on 6H-SiC on-axis substrate in Si50Ge50 melt and (b) on 4H-SiC on-axis substrate in Si25Ge75 melt; (c) Raman spectra collected on areas with different morphology in a.
Figure 8. Mechanism proposal concerning the competition between homoepitaxy and 3C-SiC. (a) 6H-SiC 3.5° off seed and low growth temperature, (b) 6H-SiC 3.5° off seed and high growth temperature, and (c) 6H-SiC on-axis seed and high growth temperature.
these islands and at the expense of homoepitaxial growth, even in the case of 3.5° off substrate (Figure 8a). If one keeps now this misorientation and increases the growth temperature, the islands height will decrease down to a point where it becomes insufficient compared to the vertical component of the lateral growth (Figure 8b). In this case, the homoepitaxial layer will overgrow the 3C-SiC islands to give a complete homoepitaxial film at the end. Finally, if one considers an on-axis substrate with identically high temperature (Figure 8c), then the smaller islands are still sufficiently tall, compared to the vertical component of the growth, to give a complete 3C-SiC layer by lateral enlargement, as in the case of Figure 8a. When comparing to CVD, one can notice that the same general trend is found: increasing the misorientation decreases the temperature of homoepitaxial growth. In CVD, this is due to the reduction in the terraces width associated with misorientation increase,
whereas in our case it is rather related to an increase of the vertical component of the lateral growth. A mixing of 3C and 6H within the layers was obtained on the on-axis 6H substrate in the boundary zone between twinfree 3C and homoepitaxy (Figure 4a). This could come from local inhomogeneities in island height or density so that the dissolution induced by the temperature increase can be more effective to promote homoepitaxial growth in some parts where these islands are smaller or less dense. This polytype mixing was not detected on 4H-SiC substrate, but, as mentioned before, this could be due to the dense network of twins that significantly disturb the layer (Figure 7b). One cannot exclude, without performing any polytype mapping, that small homoepitaxial inclusions are present. This would be in better accordance with the present model of competition between island enlargement and homoepitaxy.
Vapor–Liquid–Solid Growth of 3C-SiC
Finally, concerning the single-domain 3C-SiC layers formation, it occurs in defined conditions of temperature and melts composition. One has to notice that these temperatures are well above the ones of 3C-SiC island formation (3C-SiC layers are obtained even at 1200 °C). It means that, for a fixed melt composition, the nucleation step is the same for all these layers. So the difference in layer characteristics (twin density or polytype) should come from the difference in thermal treatment after this nucleation. As a general rule, one has to increase temperature when decreasing the Si content of the melt to avoid twin formation, even if these conditions change from one substrate polytype to another. If one compares now this trend with the C solubility evolution inside the melt, some similarities can be found. Indeed, the C solubility increases when increasing the Si content of the melt and/or the temperature. So, a trend seems to emerge between single-domain 3C materials and C solubility in the liquid. All seem to happen as if the singledomain zones of Figure 4 were following C isosolubility conditions. Note that this hypothetical optimal C solubility range is rather restricted. Assuming this hypothesis of C solubility driven mechanism, the selection of one orientation of the 3C layer can occur at two stages: (1) during the inevitable partial dissolution of the islands upon heating up to the temperature plateau (before propane addition) or (2) during the lateral enlargement of the islands upon propane addition in the reactor. If the dissolution stage is assumed, it should mean that the dissolution kinetics is orientation dependent so that one family of islands (monooriented) is dissolved faster than the other. In this case, there is a specific C solubility in the melt for higher dissolution selectivity between the two orientations of the 3C islands. If we consider now the island enlargement stage upon propane addition, it should mean that the C solubility in the melt affects the selectivity of the lateral growth rate orientation so that one family of islands may be enlarged faster than the other. If the 3C orientation selection depends only on the C solubility in the liquid, using different SiC polytypes as substrates should not affect the position of the single-domain zone in the graphs of Figure 4. But when comparing the use of on-axis 4H and 6H-SiC substrates, one can notice a significant shift toward low temperatures in the case of 4H. This should mean that another parameter than C solubility in the melt has to be taken into account, and it is most probably related to the seed surface. The two main differences between on-axis 4H and 6H-SiC(0001) surfaces are (1) the free surface energy of these basal planes and (2) the morphology or surface reconstruction. Concerning the free surface energy, they were found to be different between the two polytypes,13 which means that the interfacial energy between the melt and these crystals should also be different.14 This parameter may have an effect on the 3C-SiC islands formation by accelerating or slowing down the dissolution-precipitation mechanism. Concerning the surface morphology of the seed, we have shown that the islands forming on a 6H-SiC(0001) on-axis seed are elongated and parallel to the step edges. Such surfaces are known to give very regular and parallel step reconstruction, much more regular than 4H ones.15,16 In the present study, the seeds were etched under H2 at 1400 °C for 10 min, which results in a more undulated and disturbed surface.16 This may have an effect on the elongated form of the 3C islands and possibly on the selection of their orientation. But the exact interaction between the step structure and the 3C nucleation is still unclear. Latu-Romain et al.17 raised the possibility of orientation selection by a preferential 3C-SiC islands enlargement along the step edges. This selection could
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happen by considering a blocking effect of the steps and the potential sink for adatoms at the R-SiC/3C/gas phase triple point along the step edges. However, this mechanism does not really fit with our results for two reasons. First, the height of the 3CSiC islands, which are formed by the contact with the Si-Ge melt, is often much higher (few tens of nanometers) than that of the substrate surface steps (∼0.7 nm) (see Figure 1b). Second, the majority of the 3C islands in the present study covers several terraces each. As consequence, in our case the step edges do not seem to constitute a real blockage for the lateral growth of the islands. Conclusion We have shown that single domain 3C-SiC layers can be grown on various kinds of R-SiC substrates using a VLS mechanism in Si-Ge melts. The optimal conditions (temperature and melt composition) to avoid twin formation depend on the chosen kind of substrate (polytype, misorientation). It was found possible to grow twin-free 3C-SiC layers either on 6H or 4HSiC substrates, which is very interesting for device application if one aims at fabricating 3C/6H or 3C/4H heterostructures. It was not possible to avoid twin formation on Si face 4H offoriented and C face 6H on-axis substrates. For each substrate giving a twin-free 3C layer, the range of these optimal conditions is large enough for reproducibility purposes. The resulting surface morphology of these twin-free layers was found to change significantly from one substrate to another. Starting from the observation of 3C-SiC island formation on the surface during the initial heating ramp, we proposed a growth mechanism mainly based on the competition between homoepitaxy and enlargement of these islands upon propane addition. This allows explaining the results obtained on the different kinds of substrates. Concerning the mechanism for selection of one 3CSiC orientation, the understanding is less advanced though some links were proposed with the carbon solubility inside the melt and/or with the seed surface nature (surface energy, step regularity). Acknowledgment. M.S. gratefully acknowledges the financial support by the Lebanese National Council for Scientific Research (CNRSL). The authors thank the CECOMO of Lyon for the Micro-Raman measurements and D. Chaussende, P. Chaudouet, and F. Robaut for their help and discussion concerning the EBSD characterizations.
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1050 Crystal Growth & Design, Vol. 8, No. 3, 2008 (10) Soueidan, M.; Ferro, G.; Kim-Hak, O.; Robaut, F.; Dezellus, O.; Dazord, J.; Cauwet, F.; Viala, J. C.; Nsouli, B. Acta Mater. 2007, 55, 6873. (11) Ferro, G.; Jacquier, C. New J. Chem. 2004, 28, 889. (12) Scace, R. I.; Slack, G. A. J. Chem. Phys., 1959, 30 (6), 1551. (13) Syväjärvi, M.; Yakimova, R.; Janzèn, E. J. Electrochem. Soc. 1999, 146 (4), 1565. (14) Yakimova, R.; Syväjärvi, M.; Janzèn, E. Mater. Sci. Forum 1998, 264– 268, 159.
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CG070499+
