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Early stages of growth of β-SiC on Si by MBE
,. . . . . . . .
ELSEVIER
CRYSTAL OIROWTH
Journal of Crystal Growth 157 (1995) 392-399
Early stages of growth of fl-SiC on Si by MBE K . Z e k e n t e s b , . , V. P a p a i o a n n o u a, B. P e c z c, j. S t o e m e n o s a a Physics Department, Aristotle University of Thessaloniki, GR-54006 Thessaloniki, Greece b Foundation for Research and Technology, Heraklio, Crete, Greece c Research Institute for Technical Physics, Hungarian Academy of Sciences, P.O. Box 76, H-1325 Budapest, Hungary
Abstract
The early stages of growth of /3-SIC formed by carbonization in a MBE system at different substrate temperatures and flow rates were systematically studied by cross-section transmission electron microscopy (TEM), high-resolution TEM (HRTEM) and atomic force microscopy (AFM). Fluctuations of the moir6 pattern reveal local strain variations at the /3-SiC/Si interface, suggesting that the various /3-SIC island nuclei, formed during the carbonization process, are not restricted to having a parallel epitaxy relative to the substrate. The roughness at the back side of the fl-SiC is attributed to the interdiffusion of the Si and C atoms during the carbonization process in order to form the fl-SiC. The formation of /3-SIC islands is related to the Si migration from the Si surface.
1. Introduction
fl-SiC is a very promising semiconductor for high-temperature, high-frequency and high-power electronic devices because of its wide band gap, high-saturated electron velocity and high-breakdown electric field [ I]. Recently, a technique for the epitaxial growth of fl-SiC on Si substrates by chemical vapor deposition (CVD) was developed. An essential step in this method is the formation of a thin carbonized layer before the deposition of SiC by CVD [2]. The perfection of this layer is crucial for the quality of the overgrown /3-SIC. Ultra-high vacuum molecular beam epitaxy (MBE) is the most appropriate technique for the growth of the first carbonized layers due to the better process control and the very clean growth ambient
* Corresponding author. Fax: +30 31 214276.
which can be provided. In order to avoid implications due to the presence of hydrocarbons and hydrogen, a solid carbon source was used for the growth of the carbonized layers. Up to now very few attempts have been reported for MBE growth by solid carbon sources [3]. In the present work, the carbonized layers were grown by MBE at different substrate temperatures and different carbon flow rates. The structural characteristics of these layers were systematically studied by cross-section TEM, HRTEM and AFM.
2. Experimental procedure
The /3-SIC layers were formed by the carbonization of a Si wafer in a molecular beam epitaxy (MBE) system. Pure pyrolytic graphite was used as the carbon source. Pyrolytic graphite is the purest dense form of carbon. However, high purity leads to extremely low vapor pressure therefore, temperatures
0022-0248/95/$09.50 © 1995 Elsevier Science B.V. All rights reserved SSDI 0022-0248(95)00330-4
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K. Zekentes et al./Journal of Crystal Growth 157 (1995) 392-399
of about 2500°C are necessary in order to obtain practical deposition rates in the MBE system. The conversion of Si(001) surface to B-SiC has been performed by exposing the Si surface to the flux of elemental carbon. The substrates were 2 inch Si(001) wafers misoriented by 2 ° to 4 ° towards [110]. In order to ensure a perfectly clean Si surface, a silicon layer about 200 n m thick was epitaxially deposited at 725°C on the Si wafer in the MBE system prior to the carbon evaporation. Different carbon sublimation rates and substrate temperatures, in the range of 650 to 950°C have been used to grow about twenty B-SiC layers. The growth process was controlled in-situ by reflection high-energy electron diffraction (RHEED). The grown layers were characterized by plane view and cross-section TEM. The surface morphology of the B-SiC films was also studied by AFM.
3. Results and discussion In order to study the early stage of growth of B-SiC, seventeen specimens were grown. The conditions of growth and the structural characteristics of
the B-SiC layers are shown in Table 1, the following qualitative results can be drawn from them. (i) Due to 21% misfit in the B - S i C / S i system, the 3D-epitaxial growth is the most favorable; in this way, the interfacial energy is significantly reduced. Therefore, in the early stage of the growth, the epitaxy proceeds solely by islands, all the B-SiC islands are in good epitaxial relation with the Si substrate. (ii) At a constant deposition temperature, the density of the B-SiC islands increases as the carbon flow rate increases. For example, the specimens 7a and 7f were deposited at 725°C. From them, in specimen 7a, which was deposited at a low flow rate of 2 X 10 -4 n m / m i n , the island density was 3 X 10 8 c m - 2 . In specimen 7f, which was deposited at a rate of 10 -1 n m / m i n , the island density was 2 X 10 ~1 cm -2. Thus by increasing the flow rate, the island density increases. In spite of the 2.5 longer deposition time in specimen 7f compared to that of 7a, the grain size in the latter was twice of those in the former. This behavior can be explained by considering that the density of the stable islands is affected by the deposition rate and the deposition temperature, while the
Table 1 Conditions of carbonizationand structural characteristics of the /3-SICfilms Specimen Deposition Flow rate Deposition Meanthickness Meansize of Islanddensity Structural Surfacecovering number temperature(C°) (nm/min) time(min) (nm) grain (nm) (cm-2 ) characteristics(%) 9a 9b 9C 8a 8b 7a 7b 7c 7d 7e 7f 7g 7h 7i 7j 6a 6b
970 960 960 860 860 725 725 725 725 725 725 725 725 725 725 660 600
0.0003 0.0020 0.1700 0.1900 0.2250 0.0002 0.0010 0.0060 0.0300 0.0800 0.1000 2.0000 5.4000 6.0000 16.0000 0.4500 0.4500
90 120 204 120 120 3 40 30 204 16 8 2 1.5 2 1 8 25
0.025 0.26 35 23 27 0.005 0.025 0.02 6 1.2 0.85 17 8.1 18 16 0.1 0
200 250 150 40 80 14 4.5 1.8 18 7.2 7 15 8 10 16 5 0
S.C. = Single crystalline;con. = continuous;a. con. = almost continuous.
3.0 X 10 6 2.0 X 107 a. con. con. con. 3.0 × 108 5.0 × 10z° 2.0 X 10~1 1.0x 10tl 2.5 X 10tl 2.0 X 10~a con. con. con. con. 1.0 X 1012 0
S.C. 95% >> 95% >> 45% >> 80% >> 60% >> 95% >> 95% >> 95% >> 100% >> 95% >> 95% >> 85% >> 90% >> 85% >> 85% >> 90% -
0.1 1 70 95 95 0.05 0.8 0.5 25 10 8 100 95 100 100 20 0
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K. Zekentes et al./ Journal of Crystal Growth 157 (1995) 392-399
size of the islands is mainly determined by the deposition time. Although the deposition time in specimen 7a was shorter than in 7f, the mean size of the islands was larger due to the lower grain density. In the case of specimen 7d, which was deposited at a high flow rate for a deposition time of 204 s, the island density was high; however, due to the very long deposition time, the mean grain size is also large, 18 nm. (iii) The density of the stable nuclei is being reduced as the substrate temperature increases at a constant flow rate. Compare, for example, specimen 9a and 7a, grown at 970 and 725°C, respectively, having comparable flow rates, 3 X l 0 - 4 and 2 X 10 - 4 rim. In spite of the very long deposition time, the island density of specimen 9a was two orders of magnitude lower than of specimen 7a. Very large islands were formed in the case of specimen 9a, as shown in Table 1, in agreement to paragraph (ii). (iv) The lowest temperature for the formation of /3-SIC is 660°C, below this temperature no reaction occurs. The highest island density 1 X 1012 c m - 2 was observed when a high flow rate was combined with the lowest deposition temperature for SiC formation, as is the case for specimen 6a. The results (i) to (iii) are in agreement with the classical nucleation theory [4]. The islands of r-SiC are developed along the (110) crystallographic directions, very often exhibiting a dendritic growth, as shown in the plane view micrograph in Fig. la. The islands are heavily twinned as shown in the dark field (DF) micrograph in Fig. lb. The related diffraction pattern in the inset of Fig. lb reveals that the islands of /3-SIC are in epitaxial relation with the Si substrate although slightly misoriented. The misorientation of the nuclei during the early stage of growth was also confirmed, in commercially available wafers [5], from the rotation of the moir6 patterns, which are produced by the diffraction of the e-beam in the two overlapping lattices at the /3-SiC/Si interface, as shown in the plane view micrograph in Fig. 2. Moir6 fringes were formed because the electron beam was originated from the superposition of the two lattices. Various defects and local strain variations result in the observed fringe irregularities. The residual strain at the interface was estimated from the periodicity of the moir6 patterns. From Fig. 2, the periodicity of the
Fig. 1. Plane view micrographs from specimen 9b. (a) General view of the /3-SIC islands, which were formed on the Si substrate. Notice the strong black/white contrast around the islands. The preferential growth along the (110) directions is evident. The dendritic growth of the islands is shown in the high magnification micrograph in the inset. (b) DF micrograph taken from the 220 /3-SIC diffraction spot denoted by an arrow in the inset of the figure. Satellite spots due to double diffraction are evident. The high density of the twins results in the laminar structure of the island.
displacement-type moir6 pattern for the (220) reflection was measured and the mean D220 spacing was found to be 0.81 nm. This value was compared with the theoretical one, which is given by Eq. (1), o2n~o
dsi d s i c dsi _ dsic
(1)
where dsi a n d dsi C are the d lattice spacings for Si and r-SiC, respectively, when (220) is the operating reflection. From Eq. (1), the theoretical value of the O2n~o spacing of the moir6 pattern is 0.778 nm. Comparing the O2r~o spacing of the experimental and
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395
parallel epitaxy relative to the substrate. The actual misorientation, (9, of the lattice can be estimated from the misorientation, to, of the moir6 pattern according to Eq. (2), ,o = 0 / a ,
Fig. 2. Plane view micrographfrom the fl-SiC/Si interface, moir6 pattem along the [110] and [110] directions are evident. Distortion of the moir6 patterns at the low angle boundaries are denoted by arrows. The related diffraction patterns are shown in the inset.
theoretical values of the moir6 pattern, we can deduce the remaining misfit due to the residual strain of the lattices t5r which, in this case, was estimated to be 8 r = 0.009. Since the misfit due to the residual strain of the lattices is only 0.9%, it is evident that the 21% misfit is mainly relieved at the /3-SiC/Si interface by misfit dislocation and low angle boundaries, as will be discussed in the next paragraph. Rotation of the moir6 pattern up to 11 ° has also been found, as shown in Fig. 2. Therefore, the various island nuclei are not restricted to having a
(2)
where (5 is the misfit. Actual misorientations of the lattice up to 2.5 ° were found. The mean size of the nuclei is about 10 nm and due to their misorientation low angle boundaries are formed during the coalescence, which is another source of defects. Low angle boundaries are denoted by arrows in Fig. 2 resulting in the distortion of the moir6 pattern as they cross the boundaries. The misorientation of the nuclei during the early stage of growth results in the relaxation of the two lattices as molecular dynamic calculations have shown [6]. HRTEM observations also reveal the misorientation of the/3-SIC islands. The introduction of the defects at the grain boundary (GB) during the coalescence of the nuclei A and B in the cross-section H R T E M is shown in Fig. 3, a parallel shift of the lattices at the boundary is observed denoted by an arrow at the GB. It is noticed that the nucleus denoted by the letter B is slightly misoriented compared with the Si substrate; in this case, stacking faults are generated denoted by the letter SF in Fig. 3. In the case of the nucleus A, where all the misfit is absorbed by the misfit dislocations at the interface,
Fig. 3. Cross-sectionHRTEM micrographfrom the /3-SiC/Si interface. A low angle grain boundary (GB) due to the coalescenceof the two islands, denoted by the letters A and B, is evident. Planar defects related with island B are shown in the inset.
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no other defects are generated inside the nucleus. In this case, an extra fl-SiC plane is introduced in every four silicon lattice planes in order to absorb the 21% misfit, as is shown at the right-hand side of Fig. 3. Most of the defects which propagate inside the/3-SIC films are generated due to the small misorientation of the/3-SIC nuclei during the early stage of growth. The formation of /3-SIC islands during the carbonization process is related with a significant Si migration from the Si substrate. The /3-SIC islands are surrounded by a strong black and white contrast, as shown in Fig. la. The contrast is partially attributed to the lattice distortion due to strain development and also to the thinner Si substrate around the islands due to Si surface migration, for the formation of the /3-SIC. The Si surface migration is evident from the depression of the silicon surface in the vicinity of the islands, as shown by the cross-section micrograph in Fig. 4a. The loss of silicon around the fl-SiC island was also confirmed by the AFM micrograph in Fig. 4b. It was possible by AFM image processing to estimate the total volume of the missing silicon around each island and to compare it with the volume of the corresponding SiC island. It was found that the two volumes are equal. This is expected if we consider that the conversion of one unit volume of Si to SiC results in an increase in volume only by 3.5%. Surface diffusion of Si at relatively low temperatures of 650 to 970°C is expected because according to the recent theoretical calculations, the activation energy for surface diffusion in Si is very low, 0.6 eV [7]. In contrast the activation energy of bulk Si self-diffusion is very high, about 5 eV, and therefore bulk diffusion becomes significant above 1200°C [8]. As the carbonization process advances the surface is covered by SiC which inhibits Si surface diffusion. In this case, the Si, which is needed for the growth of SiC, is provided by open areas on the surface which are the sources of Si supply. Therefore, pits are formed having the shape of reversed tetragonal pyramids bounded by {111} crystallographic facets [9], as shown in the cross-section micrograph in Fig. 5a and the AFM micrograph in Fig. 5b. Even when the surface is completely covered by SiC, the cavities continue emitting Si atoms. The possible paths for migration are through the existing misfit dislocations and the other extended defects, as expected by the Kirkendall effect,
c flux MSF:s
and
icrotwi n~
-"2_
Fig. 4. (a) Cross-section TEM micrograph from specimen 9b. A fl-SiC island with well-developed {111} facets is shown. Depression of Si at the vicinity of the island reveals surface Si migration. (b) AFM micrograph from the /3-SIC island is evident (specimen 9b). In the inset, the profile of the surface along one run is shown. (c) Diagram of fl-SiC island formation by Si surface diffusion during the carbonization process.
K. Zekentes et al./ Journal of Crystal Growth 157 (1995) 392-399
397
diagrammatically shown in Fig. 5c. Reverse preferential growth of SiC occurs in these cavities, as shown in Fig. 6, revealing significant carbon diffusion through the overgrown SiC. The /3-SiC/Si interface is not atomically flat due to the strong reaction of the carbonized zone with the
Fig. 6. (a) Cross-sectionTEM micrograph from a thick r-SiC, the cavities are now covered by a continuous /3-SIClayer. Additional /3-SIC islands were grown on the back side of the /3-SIC layer inside the cavity.
i P[an~
Carbon o
5i
"
C
Fig. 5. (a) Pits in the Si substrate having the form of reverse tetragonal pyramids bounded by {111} crystallographic facets are shown in the cross-sectionTEM micrograph. (b) AFM micrograph from a continuous r-SiC layer (specimen 9c) revealing pits at the surface. (c) The supply of Si from the cavities through structural defects for the formation of /3-SIC at the surface of the /3-SIC layer is diagrammatically shown.
Si substrate to form /3-SIC. The reaction occurs in depth by interdiffusion of carbon and Si atoms. Spatial differences in the degree of interdiffusion result in the formation of uneven nuclei having shifted lattice planes. The introduction of defects at the boundaries during the coalescence of the neighboring nuclei due to the shifted lattice planes is shown in Fig. 3. The uneven steps are also responsible for the high density of antiphase boundaries (APBs). The epitaxial growth of r - s i c having the zincblende structure, on a Si substrate having an odd number of steps of atomic layers, results in the formation of APBs. This problem can be avoided if the steps are an even number of atomic layers high. The formation of APBs during the epitaxial growth of GaAs on Si is suppressed if vicinal 4°-off (100) substrate is used because the steps are even [10]. The use of vicinal Si substrate in the case of epitaxially grown r - s i c does not prevent the formation of APBs [11], because the carbonization process is not a simple epitaxy but rather a strong chemical reaction, as Kim et al. [12] pointed out, which occurs at a depth of several atomic layers by interdiffusion of Si and carbon atoms resulting in a rough interface, as shown in Fig. 3. A limited carbon diffusion into Si was also predicted by molecular dynamic calculations [6]. The roughness of the /3-SiC/Si interface was also confirmed by X T E M observations in commercially available r - s i c films grown epitaxially on Si wafers by the CVD method [11]. The back side of
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K. Zekentes et al. / Journal of Crystal Growth 157 (1995) 392-399
Fig. 7. AFM micrograph from the back side of the t-SiC after dissolution of the Si substrate. The roughness of the t-SiC at the interface in atomic scale is evident.
the /3-SIC films, which actually characterizes the quality of the /3-SCI/Si interface as seen from the /3-SIC side, was studied by AFM in atomic scale. The back side was studied after the dissolution of the Si substrate. As solvent, a solution of 95% HNO 3 and 5% HF was used, which dissolves the Si leaving the /3-SIC intact. The distortion of the atomic plane as well as the roughness at the back side becomes evident by AFM, as shown in the 3D image in Fig. 7. The roughness of the surface must not be attributed to the roughness of the Si substrate that, even in the case of a vicinal wafer, is atomically smooth, but to the strong reaction of the carbonized zone with the silicon substrate to form /3-SIC.
misoriented, resulting in the formation of planar defects. Deposition at lower temperatures, around 725°C, gives the best SiC films. This result is in agreement with similar observations in SiC films deposited by CVD [13]. Thick SiC films formed by carbonization are partially polycrystalline due to the formation of multiple-first and higher-order twins.
4. Conclusions
References
SiC is formed at temperatures as low as 660°C. In the early stage of growth, the epitaxy proceeds solely by islands which develop {111} facets resulting in the formation of microtwins. The growth of the islands occurs by Si surface diffusion from the surroundings of the SiC islands. The islands are slightly
Acknowledgements This work was supported by EU BRITE-EURAM under contract No. BRE2-CT92-0211.
[1] P.A. Ivanov and V.E. Cheinokov Semicond. Sci. Technol. 7 (1992) 863. [2] S. Nishimo, J.A. Powell and H.A. Will, Appl. Phys. Lett. 45 (1983) 460. [3] Durupt et al., Thin Solid Films 88 (1992) 291. [4] J.P. Hirth and G.M. Pound, Condensation and Evaporation (Pergamon, Oxford, 1963).
K. Zekentes et al./ Journal of Crystal Growth 157 (1995) 392-399
[5] CREE Research Inc., 2810 Meridian Parkway, Durham, NC27713. [6] M. Kitabataka, M. Deguch and T. Hirao, J. Appl. Phys. 74 (1993) 4438. [7] G. Brocks, P.J. Kelly and R. Car, 20th Int. Conf. on the Physics of Semiconductors, Thessaloniki, Greece, Aug. 6-10, 1990, Eds. E.M. Anastassakis and J.D. Ioannopoulos (World Scientific, Singapore, 1990) p. 127. [8] S.M. Hu, Mater. Res. Soc. Proc. 59 (1985) 223. [9] J.P. Li, A.J. Steckl, I. Golecki, F. Reidinger, L. Wang, X.J. Ning and P. Pirows, Appl. Phys. Lett. 62 (1993) 3135.
399
[10] M. Kawabe and T. Ueda, Jap. J. Appl. Phys. 25 (1986) L285. [11] J. Stoemenos, C. Dezauzier, G. Arnaud, S. Contreras, J. Camassel, J. Pascual and J.L. Robert, Mater. Sci. Eng. B 29 (1995) L285. [12] H.J. Kim, R.F. Davis, X.B. Cox and R.W. Linton, J. Electrochem. Soc. 134 (1987) 2269. [13] H. Kobashi, M. Hirai, M. Kusaka, M. lwami, Y. Yokota, H. Kado and T. Tohda, 21th Int. Conf. on the Physics of Semiconductors, Beijing, China, Aug. 10-14, 1992, Eds. P. Jiang and H.Z. Zheng (World Scientific, Singapore, 1992) p. 373.
ELSEVIER
CRYSTAL OIROWTH
Journal of Crystal Growth 157 (1995) 392-399
Early stages of growth of fl-SiC on Si by MBE K . Z e k e n t e s b , . , V. P a p a i o a n n o u a, B. P e c z c, j. S t o e m e n o s a a Physics Department, Aristotle University of Thessaloniki, GR-54006 Thessaloniki, Greece b Foundation for Research and Technology, Heraklio, Crete, Greece c Research Institute for Technical Physics, Hungarian Academy of Sciences, P.O. Box 76, H-1325 Budapest, Hungary
Abstract
The early stages of growth of /3-SIC formed by carbonization in a MBE system at different substrate temperatures and flow rates were systematically studied by cross-section transmission electron microscopy (TEM), high-resolution TEM (HRTEM) and atomic force microscopy (AFM). Fluctuations of the moir6 pattern reveal local strain variations at the /3-SiC/Si interface, suggesting that the various /3-SIC island nuclei, formed during the carbonization process, are not restricted to having a parallel epitaxy relative to the substrate. The roughness at the back side of the fl-SiC is attributed to the interdiffusion of the Si and C atoms during the carbonization process in order to form the fl-SiC. The formation of /3-SIC islands is related to the Si migration from the Si surface.
1. Introduction
fl-SiC is a very promising semiconductor for high-temperature, high-frequency and high-power electronic devices because of its wide band gap, high-saturated electron velocity and high-breakdown electric field [ I]. Recently, a technique for the epitaxial growth of fl-SiC on Si substrates by chemical vapor deposition (CVD) was developed. An essential step in this method is the formation of a thin carbonized layer before the deposition of SiC by CVD [2]. The perfection of this layer is crucial for the quality of the overgrown /3-SIC. Ultra-high vacuum molecular beam epitaxy (MBE) is the most appropriate technique for the growth of the first carbonized layers due to the better process control and the very clean growth ambient
* Corresponding author. Fax: +30 31 214276.
which can be provided. In order to avoid implications due to the presence of hydrocarbons and hydrogen, a solid carbon source was used for the growth of the carbonized layers. Up to now very few attempts have been reported for MBE growth by solid carbon sources [3]. In the present work, the carbonized layers were grown by MBE at different substrate temperatures and different carbon flow rates. The structural characteristics of these layers were systematically studied by cross-section TEM, HRTEM and AFM.
2. Experimental procedure
The /3-SIC layers were formed by the carbonization of a Si wafer in a molecular beam epitaxy (MBE) system. Pure pyrolytic graphite was used as the carbon source. Pyrolytic graphite is the purest dense form of carbon. However, high purity leads to extremely low vapor pressure therefore, temperatures
0022-0248/95/$09.50 © 1995 Elsevier Science B.V. All rights reserved SSDI 0022-0248(95)00330-4
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of about 2500°C are necessary in order to obtain practical deposition rates in the MBE system. The conversion of Si(001) surface to B-SiC has been performed by exposing the Si surface to the flux of elemental carbon. The substrates were 2 inch Si(001) wafers misoriented by 2 ° to 4 ° towards [110]. In order to ensure a perfectly clean Si surface, a silicon layer about 200 n m thick was epitaxially deposited at 725°C on the Si wafer in the MBE system prior to the carbon evaporation. Different carbon sublimation rates and substrate temperatures, in the range of 650 to 950°C have been used to grow about twenty B-SiC layers. The growth process was controlled in-situ by reflection high-energy electron diffraction (RHEED). The grown layers were characterized by plane view and cross-section TEM. The surface morphology of the B-SiC films was also studied by AFM.
3. Results and discussion In order to study the early stage of growth of B-SiC, seventeen specimens were grown. The conditions of growth and the structural characteristics of
the B-SiC layers are shown in Table 1, the following qualitative results can be drawn from them. (i) Due to 21% misfit in the B - S i C / S i system, the 3D-epitaxial growth is the most favorable; in this way, the interfacial energy is significantly reduced. Therefore, in the early stage of the growth, the epitaxy proceeds solely by islands, all the B-SiC islands are in good epitaxial relation with the Si substrate. (ii) At a constant deposition temperature, the density of the B-SiC islands increases as the carbon flow rate increases. For example, the specimens 7a and 7f were deposited at 725°C. From them, in specimen 7a, which was deposited at a low flow rate of 2 X 10 -4 n m / m i n , the island density was 3 X 10 8 c m - 2 . In specimen 7f, which was deposited at a rate of 10 -1 n m / m i n , the island density was 2 X 10 ~1 cm -2. Thus by increasing the flow rate, the island density increases. In spite of the 2.5 longer deposition time in specimen 7f compared to that of 7a, the grain size in the latter was twice of those in the former. This behavior can be explained by considering that the density of the stable islands is affected by the deposition rate and the deposition temperature, while the
Table 1 Conditions of carbonizationand structural characteristics of the /3-SICfilms Specimen Deposition Flow rate Deposition Meanthickness Meansize of Islanddensity Structural Surfacecovering number temperature(C°) (nm/min) time(min) (nm) grain (nm) (cm-2 ) characteristics(%) 9a 9b 9C 8a 8b 7a 7b 7c 7d 7e 7f 7g 7h 7i 7j 6a 6b
970 960 960 860 860 725 725 725 725 725 725 725 725 725 725 660 600
0.0003 0.0020 0.1700 0.1900 0.2250 0.0002 0.0010 0.0060 0.0300 0.0800 0.1000 2.0000 5.4000 6.0000 16.0000 0.4500 0.4500
90 120 204 120 120 3 40 30 204 16 8 2 1.5 2 1 8 25
0.025 0.26 35 23 27 0.005 0.025 0.02 6 1.2 0.85 17 8.1 18 16 0.1 0
200 250 150 40 80 14 4.5 1.8 18 7.2 7 15 8 10 16 5 0
S.C. = Single crystalline;con. = continuous;a. con. = almost continuous.
3.0 X 10 6 2.0 X 107 a. con. con. con. 3.0 × 108 5.0 × 10z° 2.0 X 10~1 1.0x 10tl 2.5 X 10tl 2.0 X 10~a con. con. con. con. 1.0 X 1012 0
S.C. 95% >> 95% >> 45% >> 80% >> 60% >> 95% >> 95% >> 95% >> 100% >> 95% >> 95% >> 85% >> 90% >> 85% >> 85% >> 90% -
0.1 1 70 95 95 0.05 0.8 0.5 25 10 8 100 95 100 100 20 0
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size of the islands is mainly determined by the deposition time. Although the deposition time in specimen 7a was shorter than in 7f, the mean size of the islands was larger due to the lower grain density. In the case of specimen 7d, which was deposited at a high flow rate for a deposition time of 204 s, the island density was high; however, due to the very long deposition time, the mean grain size is also large, 18 nm. (iii) The density of the stable nuclei is being reduced as the substrate temperature increases at a constant flow rate. Compare, for example, specimen 9a and 7a, grown at 970 and 725°C, respectively, having comparable flow rates, 3 X l 0 - 4 and 2 X 10 - 4 rim. In spite of the very long deposition time, the island density of specimen 9a was two orders of magnitude lower than of specimen 7a. Very large islands were formed in the case of specimen 9a, as shown in Table 1, in agreement to paragraph (ii). (iv) The lowest temperature for the formation of /3-SIC is 660°C, below this temperature no reaction occurs. The highest island density 1 X 1012 c m - 2 was observed when a high flow rate was combined with the lowest deposition temperature for SiC formation, as is the case for specimen 6a. The results (i) to (iii) are in agreement with the classical nucleation theory [4]. The islands of r-SiC are developed along the (110) crystallographic directions, very often exhibiting a dendritic growth, as shown in the plane view micrograph in Fig. la. The islands are heavily twinned as shown in the dark field (DF) micrograph in Fig. lb. The related diffraction pattern in the inset of Fig. lb reveals that the islands of /3-SIC are in epitaxial relation with the Si substrate although slightly misoriented. The misorientation of the nuclei during the early stage of growth was also confirmed, in commercially available wafers [5], from the rotation of the moir6 patterns, which are produced by the diffraction of the e-beam in the two overlapping lattices at the /3-SiC/Si interface, as shown in the plane view micrograph in Fig. 2. Moir6 fringes were formed because the electron beam was originated from the superposition of the two lattices. Various defects and local strain variations result in the observed fringe irregularities. The residual strain at the interface was estimated from the periodicity of the moir6 patterns. From Fig. 2, the periodicity of the
Fig. 1. Plane view micrographs from specimen 9b. (a) General view of the /3-SIC islands, which were formed on the Si substrate. Notice the strong black/white contrast around the islands. The preferential growth along the (110) directions is evident. The dendritic growth of the islands is shown in the high magnification micrograph in the inset. (b) DF micrograph taken from the 220 /3-SIC diffraction spot denoted by an arrow in the inset of the figure. Satellite spots due to double diffraction are evident. The high density of the twins results in the laminar structure of the island.
displacement-type moir6 pattern for the (220) reflection was measured and the mean D220 spacing was found to be 0.81 nm. This value was compared with the theoretical one, which is given by Eq. (1), o2n~o
dsi d s i c dsi _ dsic
(1)
where dsi a n d dsi C are the d lattice spacings for Si and r-SiC, respectively, when (220) is the operating reflection. From Eq. (1), the theoretical value of the O2n~o spacing of the moir6 pattern is 0.778 nm. Comparing the O2r~o spacing of the experimental and
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395
parallel epitaxy relative to the substrate. The actual misorientation, (9, of the lattice can be estimated from the misorientation, to, of the moir6 pattern according to Eq. (2), ,o = 0 / a ,
Fig. 2. Plane view micrographfrom the fl-SiC/Si interface, moir6 pattem along the [110] and [110] directions are evident. Distortion of the moir6 patterns at the low angle boundaries are denoted by arrows. The related diffraction patterns are shown in the inset.
theoretical values of the moir6 pattern, we can deduce the remaining misfit due to the residual strain of the lattices t5r which, in this case, was estimated to be 8 r = 0.009. Since the misfit due to the residual strain of the lattices is only 0.9%, it is evident that the 21% misfit is mainly relieved at the /3-SiC/Si interface by misfit dislocation and low angle boundaries, as will be discussed in the next paragraph. Rotation of the moir6 pattern up to 11 ° has also been found, as shown in Fig. 2. Therefore, the various island nuclei are not restricted to having a
(2)
where (5 is the misfit. Actual misorientations of the lattice up to 2.5 ° were found. The mean size of the nuclei is about 10 nm and due to their misorientation low angle boundaries are formed during the coalescence, which is another source of defects. Low angle boundaries are denoted by arrows in Fig. 2 resulting in the distortion of the moir6 pattern as they cross the boundaries. The misorientation of the nuclei during the early stage of growth results in the relaxation of the two lattices as molecular dynamic calculations have shown [6]. HRTEM observations also reveal the misorientation of the/3-SIC islands. The introduction of the defects at the grain boundary (GB) during the coalescence of the nuclei A and B in the cross-section H R T E M is shown in Fig. 3, a parallel shift of the lattices at the boundary is observed denoted by an arrow at the GB. It is noticed that the nucleus denoted by the letter B is slightly misoriented compared with the Si substrate; in this case, stacking faults are generated denoted by the letter SF in Fig. 3. In the case of the nucleus A, where all the misfit is absorbed by the misfit dislocations at the interface,
Fig. 3. Cross-sectionHRTEM micrographfrom the /3-SiC/Si interface. A low angle grain boundary (GB) due to the coalescenceof the two islands, denoted by the letters A and B, is evident. Planar defects related with island B are shown in the inset.
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no other defects are generated inside the nucleus. In this case, an extra fl-SiC plane is introduced in every four silicon lattice planes in order to absorb the 21% misfit, as is shown at the right-hand side of Fig. 3. Most of the defects which propagate inside the/3-SIC films are generated due to the small misorientation of the/3-SIC nuclei during the early stage of growth. The formation of /3-SIC islands during the carbonization process is related with a significant Si migration from the Si substrate. The /3-SIC islands are surrounded by a strong black and white contrast, as shown in Fig. la. The contrast is partially attributed to the lattice distortion due to strain development and also to the thinner Si substrate around the islands due to Si surface migration, for the formation of the /3-SIC. The Si surface migration is evident from the depression of the silicon surface in the vicinity of the islands, as shown by the cross-section micrograph in Fig. 4a. The loss of silicon around the fl-SiC island was also confirmed by the AFM micrograph in Fig. 4b. It was possible by AFM image processing to estimate the total volume of the missing silicon around each island and to compare it with the volume of the corresponding SiC island. It was found that the two volumes are equal. This is expected if we consider that the conversion of one unit volume of Si to SiC results in an increase in volume only by 3.5%. Surface diffusion of Si at relatively low temperatures of 650 to 970°C is expected because according to the recent theoretical calculations, the activation energy for surface diffusion in Si is very low, 0.6 eV [7]. In contrast the activation energy of bulk Si self-diffusion is very high, about 5 eV, and therefore bulk diffusion becomes significant above 1200°C [8]. As the carbonization process advances the surface is covered by SiC which inhibits Si surface diffusion. In this case, the Si, which is needed for the growth of SiC, is provided by open areas on the surface which are the sources of Si supply. Therefore, pits are formed having the shape of reversed tetragonal pyramids bounded by {111} crystallographic facets [9], as shown in the cross-section micrograph in Fig. 5a and the AFM micrograph in Fig. 5b. Even when the surface is completely covered by SiC, the cavities continue emitting Si atoms. The possible paths for migration are through the existing misfit dislocations and the other extended defects, as expected by the Kirkendall effect,
c flux MSF:s
and
icrotwi n~
-"2_
Fig. 4. (a) Cross-section TEM micrograph from specimen 9b. A fl-SiC island with well-developed {111} facets is shown. Depression of Si at the vicinity of the island reveals surface Si migration. (b) AFM micrograph from the /3-SIC island is evident (specimen 9b). In the inset, the profile of the surface along one run is shown. (c) Diagram of fl-SiC island formation by Si surface diffusion during the carbonization process.
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diagrammatically shown in Fig. 5c. Reverse preferential growth of SiC occurs in these cavities, as shown in Fig. 6, revealing significant carbon diffusion through the overgrown SiC. The /3-SiC/Si interface is not atomically flat due to the strong reaction of the carbonized zone with the
Fig. 6. (a) Cross-sectionTEM micrograph from a thick r-SiC, the cavities are now covered by a continuous /3-SIClayer. Additional /3-SIC islands were grown on the back side of the /3-SIC layer inside the cavity.
i P[an~
Carbon o
5i
"
C
Fig. 5. (a) Pits in the Si substrate having the form of reverse tetragonal pyramids bounded by {111} crystallographic facets are shown in the cross-sectionTEM micrograph. (b) AFM micrograph from a continuous r-SiC layer (specimen 9c) revealing pits at the surface. (c) The supply of Si from the cavities through structural defects for the formation of /3-SIC at the surface of the /3-SIC layer is diagrammatically shown.
Si substrate to form /3-SIC. The reaction occurs in depth by interdiffusion of carbon and Si atoms. Spatial differences in the degree of interdiffusion result in the formation of uneven nuclei having shifted lattice planes. The introduction of defects at the boundaries during the coalescence of the neighboring nuclei due to the shifted lattice planes is shown in Fig. 3. The uneven steps are also responsible for the high density of antiphase boundaries (APBs). The epitaxial growth of r - s i c having the zincblende structure, on a Si substrate having an odd number of steps of atomic layers, results in the formation of APBs. This problem can be avoided if the steps are an even number of atomic layers high. The formation of APBs during the epitaxial growth of GaAs on Si is suppressed if vicinal 4°-off (100) substrate is used because the steps are even [10]. The use of vicinal Si substrate in the case of epitaxially grown r - s i c does not prevent the formation of APBs [11], because the carbonization process is not a simple epitaxy but rather a strong chemical reaction, as Kim et al. [12] pointed out, which occurs at a depth of several atomic layers by interdiffusion of Si and carbon atoms resulting in a rough interface, as shown in Fig. 3. A limited carbon diffusion into Si was also predicted by molecular dynamic calculations [6]. The roughness of the /3-SiC/Si interface was also confirmed by X T E M observations in commercially available r - s i c films grown epitaxially on Si wafers by the CVD method [11]. The back side of
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Fig. 7. AFM micrograph from the back side of the t-SiC after dissolution of the Si substrate. The roughness of the t-SiC at the interface in atomic scale is evident.
the /3-SIC films, which actually characterizes the quality of the /3-SCI/Si interface as seen from the /3-SIC side, was studied by AFM in atomic scale. The back side was studied after the dissolution of the Si substrate. As solvent, a solution of 95% HNO 3 and 5% HF was used, which dissolves the Si leaving the /3-SIC intact. The distortion of the atomic plane as well as the roughness at the back side becomes evident by AFM, as shown in the 3D image in Fig. 7. The roughness of the surface must not be attributed to the roughness of the Si substrate that, even in the case of a vicinal wafer, is atomically smooth, but to the strong reaction of the carbonized zone with the silicon substrate to form /3-SIC.
misoriented, resulting in the formation of planar defects. Deposition at lower temperatures, around 725°C, gives the best SiC films. This result is in agreement with similar observations in SiC films deposited by CVD [13]. Thick SiC films formed by carbonization are partially polycrystalline due to the formation of multiple-first and higher-order twins.
4. Conclusions
References
SiC is formed at temperatures as low as 660°C. In the early stage of growth, the epitaxy proceeds solely by islands which develop {111} facets resulting in the formation of microtwins. The growth of the islands occurs by Si surface diffusion from the surroundings of the SiC islands. The islands are slightly
Acknowledgements This work was supported by EU BRITE-EURAM under contract No. BRE2-CT92-0211.
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[5] CREE Research Inc., 2810 Meridian Parkway, Durham, NC27713. [6] M. Kitabataka, M. Deguch and T. Hirao, J. Appl. Phys. 74 (1993) 4438. [7] G. Brocks, P.J. Kelly and R. Car, 20th Int. Conf. on the Physics of Semiconductors, Thessaloniki, Greece, Aug. 6-10, 1990, Eds. E.M. Anastassakis and J.D. Ioannopoulos (World Scientific, Singapore, 1990) p. 127. [8] S.M. Hu, Mater. Res. Soc. Proc. 59 (1985) 223. [9] J.P. Li, A.J. Steckl, I. Golecki, F. Reidinger, L. Wang, X.J. Ning and P. Pirows, Appl. Phys. Lett. 62 (1993) 3135.
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[10] M. Kawabe and T. Ueda, Jap. J. Appl. Phys. 25 (1986) L285. [11] J. Stoemenos, C. Dezauzier, G. Arnaud, S. Contreras, J. Camassel, J. Pascual and J.L. Robert, Mater. Sci. Eng. B 29 (1995) L285. [12] H.J. Kim, R.F. Davis, X.B. Cox and R.W. Linton, J. Electrochem. Soc. 134 (1987) 2269. [13] H. Kobashi, M. Hirai, M. Kusaka, M. lwami, Y. Yokota, H. Kado and T. Tohda, 21th Int. Conf. on the Physics of Semiconductors, Beijing, China, Aug. 10-14, 1992, Eds. P. Jiang and H.Z. Zheng (World Scientific, Singapore, 1992) p. 373.