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SiC interface
ARTICLE IN PRESS Journal of Crystal Growth 311 (2009) 1456–1459
Contents lists available at ScienceDirect
Journal of Crystal Growth journal homepage: www.elsevier.com/locate/jcrysgro
Atomic structure of the m-plane AlN/SiC interface ¨ zgu¨r b,c, H. Morkoc- b,c, R.P. Devaty d, W.J. Choyke d, David J. Smith a ¨. O Lin Zhou a,, X. Ni b,c, U a
Department of Physics, Arizona State University, Tempe, AZ 85287, USA Department of Electrical and Computer Engineering, Virginia Commonwealth University, Richmond, VA 23284, USA Department of Physics, Virginia Commonwealth University, Richmond, VA 23284, USA d Department of Physics and Astronomy, University of Pittsburgh, Pittsburgh, PA 15260, USA b c
a r t i c l e in f o
a b s t r a c t
Article history: Received 19 September 2008 Received in revised form 10 December 2008 Accepted 22 December 2008 Communicated by M. Skowronski Available online 6 January 2009
High-temperature m-plane AlN nucleation layers have been used for the growth of planar GaN films by metalorganic chemical vapor deposition on (1 0 1¯ 0) m-plane 6H-SiC substrates. Structural studies using transmission electron microscopy reveal the presence of a novel AlN intermediary layer preceding the remainder of the 2H-AlN buffer layer. High-resolution observations and image simulations indicate that this initial AlN layer has a faulted hexagonal structure with a six-layer repeating stacking sequence of yCBCACBCBCACBy along the transverse [0 0 0 1] direction, which does not replicate the underlying 6H-SiC stacking. Based on image analysis, the space group of this novel phase is tentatively identified as P3m1, which is also hexagonal. A structural model of the m-plane AlN/SiC interface with periodic misfit defects is also proposed. Charge neutrality analysis indicates that the interface has an equal mixture of C–Al and Si–N bonds. & 2008 Elsevier B.V. All rights reserved.
PACS: 61.72.Ff 68.37.Og 81.05.Ea Keywords: A1. Characterization A3. Metalorganic chemical vapor deposition B1. Nitrides
1. Introduction The growth of m-plane GaN-based heterostructures in the nonpolar [1 0 1¯ 0] direction has received considerable attention [1–3]. Nitride films having this particular orientation are free from the large spontaneous and piezoelectric polarization fields along the growth direction that are inherent to the usual c-plane orientation. Several groups have reported the growth of m-plane GaN on g-LiAlO2 (1 0 0) or m-plane SiC substrates using metalorganic vapor deposition (MOCVD), molecular beam epitaxy (MBE) or hydride vapor phase epitaxy [1–4]. For m-plane GaN films grown on m-plane SiC substrates, a thin AlN nucleation layer is commonly employed to improve the structural quality of the subsequent GaN layer [1]. Although the interplanar mismatch between AlN and 6H-SiC is less than 1% along both [0 0 0 1] and [1 0 1¯ 0] directions [5], most m-plane nitride films still have high densities of basal-plane stacking faults (BSFs) when no epitaxial lateral overgrowth technique is employed [1,3]. These BSFs usually originate from the substrate/film interface and have low formation energy, and they are likely to have an adverse impact on the optical properties of the material since they are electrically active Corresponding author. Tel.: +1 480 965 9815.
E-mail address: [email protected] (L. Zhou). 0022-0248/$ - see front matter & 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.jcrysgro.2008.12.047
[3,6]. Recently, the epitaxial growth of a 6H-AlN polytype on m-plane 6H-SiC by MBE was reported [7]. The presence of a 6HAlN polytype inclusion in GaN grown on GaAs (0 0 1) has also been reported [8]. However, no detailed description of this 6H-AlN structure was provided. In this present paper, we present highresolution observations of a six-layer intermediary AlN phase observed at the m-plane AlN/6H-SiC interface. Structural analysis combined with image simulations revealed that the stacking sequence of this AlN layer did not exactly replicate the 6H polytype structure of the substrate. The detailed atomic structure of the AlN/SiC interface was also determined.
2. Experimental details The m-plane GaN films investigated in this study were grown on (1 0 1¯ 0) m-plane 6H-SiC substrates (less than 0.11 miscut as measured by X-ray diffraction) using MOCVD. Prior to growth, the SiC substrate was annealed ex situ in H2 atmosphere for 10 min at 1500 1C to remove the polishing-induced surface damage. Following in situ annealing of the chemically cleaned SiC substrate, a 100-nm-thick high-temperature AlN nucleation layer was deposited at 1050 1C. The temperature was then lowered to 1030 1C for subsequent growth of the m-plane GaN epilayer. Trimethylgallium
ARTICLE IN PRESS L. Zhou et al. / Journal of Crystal Growth 311 (2009) 1456–1459
(TMGa) and ammonia (NH3) were used as the Ga and N sources, respectively. The flow rates of TMGa and NH3 were kept at 117 mmol/min and 550 sccm, respectively, with the growth pressure kept at 30 Torr. Transmission electron microscopy (TEM) observations were carried out on cross-sectional samples prepared by standard mechanical polishing and final Ar ion-beam-thinning at 4.5 keV. Electron micrographs were recorded with a JEM-4000EX highresolution electron microscope (HREM) equipped with a topentry, double-tilt specimen holder and operated at 400 keV. Relevant instrumental parameters used for image simulations include: spherical aberration coefficient, Cs=1.00 mm; beam divergence, div=0.50 mrad; and focal spread D=7 nm. Photographic negatives were scanned digitally to facilitate image processing, and the molecular simulation program CERIUS was used for image simulations.
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101 0 GaN 1120 0001
AlN
5nm
3. Results and discussion Cross-sectional TEM images of the m-plane GaN/AlN/SiC heterostructure revealed threading dislocations as well as high densities of BSFs in both AlN and GaN layers. Fig. 1 is a transmission electron micrograph taken under two-beam conditions with diffraction vector g=(0 0 0 2) showing threading dislocations inside the film. The density of the threading dislocation is highest near GaN/AlN interfaces, and then decreases toward the surface of the GaN layer. The BSF density in the GaN layer well away from the heavily faulted AlN/GaN interface was estimated to be on the order of 105 cm 1. Closer examination of the AlN/GaN interface showed that many BSFs in the AlN layer continued straight through into the GaN layer. However, a significant fraction terminated at the AlN/GaN interface, whilst others originated at the interface and propagated into the GaN layer. These features are more clearly visible in the enlarged image shown in Fig. 2. Fig. 3 shows a high-resolution electron micrograph from the area of the AlN/6H-SiC interface, taken with the incident electron beam oriented along the [11 2¯ 0] zone axis. Corresponding fast Fourier transform (FTT) diffraction patterns from the boxes labeled as A, B and C are shown in Fig. 4(a), (b), and (c), respectively. Fig. 4(a) is taken from an area which includes both 6H-SiC and the bottom part of the AlN layer. Only one set of spots
Fig. 2. Cross-sectional transmission electron micrograph taken along [11 2¯ 0] zone axis showing enlarged region of AlN/GaN interface. Note BSFs present in both AlN and GaN as well as discontinuities across the interface.
1010
1120 0001 AlN
6H-SiC 6H-SiC
Fig. 3. Cross-sectional transmission electron micrograph taken along [11 2¯ 0] zone axis showing the m-plane AlN/6H-SiC interface area. Areas labeled A, B and C are further analyzed in Fig. 4.
GaN
AlN
15nm
200nm
Fig. 1. Cross-sectional transmission electron micrograph of GaN/AlN/SiC heterostructure taken under two-beam conditions with diffraction vector g=(0 0 0 2) showing threading dislocations inside the GaN film.
is apparent confirming, as expected, that the lattice mismatch between AlN and 6H-SiC is small along [0 0 0 1] and [1 0 1¯ 0] directions. Fig. 4(a) also confirms the epitaxial growth of AlN on 6H-SiC. Fig. 4(b) is taken from the AlN layer (lateral dimensions 15 nm) immediately adjacent to the 6H-SiC substrate. This FFT pattern appears to be identical to that of 6H-SiC. The presence of five extra spots between (0 0 0 0) and (0 0 0 6) spots means that there are six layers in each AlN unit cell along the [0 0 0 1] horizontal direction, which is obviously different from the usual pattern for 2H-AlN (wurtzite) material. The AlN structure in this region might, therefore, be casually termed as ‘6H-AlN’ [7]. However, as shown below, structural analysis based on the image appearance indicates that the tetrahedral stacking across the AlN unit cell is different from the normal 6H hexagonal
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1010
1010
0000
0006
1010
0000
0006
0000
0002
Fig. 4. (a–c): Corresponding FFT diffraction patterns taken from areas in Fig. 3 labeled as A, B and C, respectively, showing that the AlN immediately adjacent to the m-plane 6H-SiC has a different crystalline structure from the remaining AlN layer.
stacking sequence. The streaking of diffraction spots visible along the [0 0 0 1] direction is due to the presence of some BSFs in this region. Fig. 4(c) is an FFT diffraction pattern taken from a region of the AlN that is located 30 nm away from the AlN/6H-SiC interface. The characteristic spot pattern expected for wurtzite 2H-AlN is visible, but no spots from the ‘6H-AlN’ phase are visible. Streaking due to BSFs in the analyzed area is also apparent. Identification of the structure of this thin ‘6H-AlN’ layer is an important step towards a full understanding of the initiation of AlN growth on m-plane 6H-SiC. Previous high-resolution studies of a GaN/SiC heterointerface showed that it was possible to identify the tetrahedral stacking sequence in both of these tetrahedrally based materials [9]. A high-resolution TEM image of the AlN/SiC interface taken along the [11 2¯ 0] zone axis is shown at high magnification in Fig. 5. The characteristic yABCACBABCACBy stacking sequence along the [0 0 0 1] direction of 6H-SiC is clearly visible. Close examination of the ‘6H-AlN’ layer indicates, however, that it has a six-layer yCBCACBCBCACBy stacking sequence. This stacking sequence clearly does not replicate the stacking sequence for a 6H hexagonal polytype along the [0 0 0 1] direction, as we elaborate further below. The circles highlighted the positions where the misfits in the stacking sequence between SiC and ‘6H-AlN’ occur. In order to confirm the crystal structure of the ‘6H-AlN’, through-thickness-through-focal image simulation tableaus were generated for both the 6H-SiC and the ‘6H-AlN’. The lattice parameters used for 6H-SiC were a: 3.081 A˚ and c: 15.092 A˚ [5]. Since the structure of ‘6H-AlN’ has not previously been reported, the lattice parameters used were based on those of the 2H-AlN structure, giving values of 3.114 and 14.9376 A˚ (3 4.9792 A˚) along the a and c directions, respectively. The corresponding image simulations (defocus of 450 A˚ and crystal thickness of 60 A˚) are shown superimposed on the experimental image. The close agreement between simulated and experimental images confirms both the microscope operating conditions and the proposed stacking sequence of the interfacial AlN layer. From these results, it appears that the ‘6H-AlN’ layer acts as a metastable buffer layer for transition from the 6H stacking of 6HSiC to the 2H stacking of AlN, presumably as a means to reduce strain caused by the small lattice mismatch and chemical differences. This situation is different from the growth of stable polytype-replicated 6H-AlN on m-plane SiC by MBE [7]. In effect, the ‘6H-AlN’ replicates closely the 6H-SiC crystal structure, as reported previously for 4H-AlN grown on 4H-SiC [10–12]. From a careful comparison with the stacking sequence of 6H-SiC, it seems that the ‘6H-AlN’ has a periodic stacking fault per unit cell which is introduced by shifting a layer of atoms from the A to C position (ABCACB to CBCACB). Alternatively, the ‘6H-AlN’ might be
Fig. 5. High-resolution transmission electron micrograph of the m-plane AlN/6HSiC interface taken along [11 2¯ 0] zone axis, with matching image simulations superimposed within rectangular boxes. Circles indicate positions of stacking defects between ‘6H-AlN’ and SiC.
described as a highly defective 2H structure with stacking disorder every third 2H unit cell (CBCBCB to CBCACB). Based on the determined stacking sequence of this ‘6H-AlN’, it is interesting to speculate on the likely space group. According to Verma and Krishna [13], there are only four possible space groups for SiC, and therefore AlN, namely, P3m1, R3m1, P63mc and F43m. The latter corresponds to the cubic 3C polytype, and R3m1 corresponds to the various rhombohedral polytypes. Since this new ‘6H-AlN’ does not have screw axis symmetry (rotation combined with half the c-axis period displacement) then P63mc is also eliminated. The remaining possibility for the space group is P3m1 which is also a hexagonal space group.
ARTICLE IN PRESS L. Zhou et al. / Journal of Crystal Growth 311 (2009) 1456–1459
B C
C
A
C
B
N N
N
N
N C*
Si* C
N
? ?
?
? ?* ?* ?
?
C C
N
N
N N*
Si* C
B A
C
A
C
B
Si
C
Al
N
1010 1120
1459
be a stable configuration alone. However, an equal mixture of these two possible bonding states would have overall charge neutrality. The Al atoms can be treated as acceptors residing on the Si sublattice of SiC. Alternatively, the C atoms can be treated as acceptors residing on the N sublattice of AlN [5]. By inserting these two bonding states into the structural model at the same time, charge neutrality could then be achieved for the structures drawn in Fig. 6(b) and (c). For example, in Fig. 6(b), there are nine Si–N bonds and one C–Al bond which would give 94 extra electron charge (14 per Si–N bond) and 14 deficit electron charge, respectively. Moreover, the atoms labeled C* and Si* are only bonded by three atoms, which would thus give another charge deficit of 1 electron. Thus, the overall total charge at the interface for this model is 9/4+( 1/4)+( 1 2)=0, and is therefore feasible. Further analyses of these interface structures are in process and will be reported at a later date.
0001
Fig. 6. (a) Schematic diagram showing the proposed atomic arrangements at the m-plane AlN/6H-SiC interface. (b) and (c) Possible atomic arrangements at the m-plane AlN/6H-SiC interface.
Because 2H stacking is the stable AlN crystalline structure, it might be anticipated that the intermediary ‘6H-AlN’ will transform to 2H-AlN with increasing film thickness away from the interface. This structural transition can be realized by introducing some low-energy type-I BSFs [14] into the 6H crystal, for example, from CBCACB to CBCACA or from CBCACB to CBCBCB. The first configuration will result in a type-I BSF in the 2H-AlN layer. Thus, there would be a high density of BSFs present in the subsequent 2H-AlN layer, which is consistent with the high-resolution images and the FFT pattern shown in Fig. 4(c). Further structural analysis of the m-plane ‘6H-AlN’/6H-SiC interface was carried out. First, an atomic-scale structural model was constructed based on the structures of 6H-SiC, ‘6H-AlN’ and the appearance of the high-resolution TEM images. A schematic diagram of the structural model of the interface projected along the [11 2¯ 0] direction is shown in Fig. 6(a). In this model, all atoms at the interface are bonded by four atoms, except for the two atoms labeled with asterisks, which are bonded only by three atoms. Next, we consider the possible chemical bonding at the ‘6H-AlN’/6H-SiC interface. There are four possible bonding configurations at the interface: Si–C–N–Al, Si–C–Al–N, C–Si– Al–N and C–Si–N–Al, as described previously [5]. Based on previous structural models of the c-plane AlN/SiC interface, the relationship between the corresponding bond lengths (d) can be expressed: dSi–Al4dC–AldSi–N4dC–N. Close examination of Fig. 5 shows that there are no abrupt changes in the interplanar separation between any two lattice planes along the m-direction. Thus, the –C–N–, and –Si–Al– bonds can be eliminated from consideration. While both –C–Al– and –Si–N– bonds are consistent with our observations, they cannot be differentiated by TEM examination. Moreover, charge neutrality is essential for the stability of the interface [5]. In a tetrahedral bonding configuration, both Si and C will contribute one electron per bond, whereas Al and N will contribute 34 and 54 electrons per bond, respectively [5]. It is obvious that neither the –C–Al– (charge deficit of 14 electron per bond), nor the –Si–N– (charge excess of 14 electron per bond) could
4. Conclusions A thin six-layered AlN intermediary layer, tentatively identified as having the P3m1 space group, has been observed at the m-plane AlN/6H-SiC interface using high-resolution electron microscopy. This layer facilitates the transition from the 6H-SiC substrate to the 2H-AlN hexagonal structure. Image analysis shows that the ‘6H-AlN’ has a faulted yCBCACBCBCACBy stacking sequence along the [0 0 0 1] direction, rather than the expected 6H hexagonal arrangement. The 2H-AlN layer has a high density of BSFs. Further analysis of the m-plane AlN/6H-SiC interface indicates that it must have an equal mixture of C–Al and Si–N bonds.
Acknowledgments This work was partially supported by ONR Grant N-00014-041-0020. The work at Virginia Commonwealth University was supported by the Air Force Office of Scientific Research. We acknowledge use of facilities in the John M. Cowley Center for High-Resolution Electron Microscopy at Arizona State University. References [1] T. Kawashima, T. Nagai, D. Iida, A. Miura, Y. Okadome, Y. Tsuchiya, M. Iwaya, S. Kamiyama, H. Amano, I. Akasaki, J. Crystal Growth 298 (2007) 261. [2] P. Waltereit, O. Brandt, A. Trampert, H.T. Grahn, J. Menniger, M. Ramsteiner, M. Reiche, K.H. Ploog, Nature 406 (2000) 865. [3] B.A. Haskell, T.J. Baker, M.B. McLaurin, F. Wu, M.P.T. Fini, S.P. DenBaars, J.S. Speck, S. Nakamura, Appl. Phys. Lett. 86 (2005) 111917. [4] M. McLaurin, T.E. Mates, J.S. Speck, Appl. Phys. Lett. 86 (2005) 262104. [5] F.A. Ponce, C.G. Van de Walle, J.E. Northrup, Phys. Rev. B 53 (1996) 7473. [6] R. Kro¨ger, T. Paskova, S. Figge, D. Hommel, A. Rosenauer, Appl. Phys. Lett. 90 (2007) 081918. [7] D.M. Schaadt, O. Brandt, A. Trampert, H.-P. Scho¨nherr, K.H. Ploog, J. Crystal Growth 300 (2007) 127. [8] M. Funato, T. Ishido, S. Fujita, S. Fujita, Appl. Phys. Lett. 76 (2000) 330. [9] J.N. Stirman, F.A. Ponce, A. Pavlovska, I.S.T. Tsong, D.J. Smith, Appl. Phys. Lett. 76 (2000) 822. [10] N. Onojima, J. Suda, T. Kimoto, H. Matsunami, Appl. Phys. Lett. 83 (2003) 5208. [11] M. Horita, J. Suda, T. Kimoto, Appl. Phys. Lett. 89 (2006) 112117. [12] M. Horita, T. Himoto, J. Suda, Appl. Phys. Lett. 93 (2008) 082106. [13] A.R. Verma, P. Krishna, Polymorphism and Polytypism in Crystals, Wiley, New York, 1966 p. 160. [14] J. H. Edgar, S. Strite, I. Akasaki, H. Amano, C. Wetzel, Properties, Processing and Applications of Gallium Nitride and Related Semiconductors, INSPEC, 1999.
Contents lists available at ScienceDirect
Journal of Crystal Growth journal homepage: www.elsevier.com/locate/jcrysgro
Atomic structure of the m-plane AlN/SiC interface ¨ zgu¨r b,c, H. Morkoc- b,c, R.P. Devaty d, W.J. Choyke d, David J. Smith a ¨. O Lin Zhou a,, X. Ni b,c, U a
Department of Physics, Arizona State University, Tempe, AZ 85287, USA Department of Electrical and Computer Engineering, Virginia Commonwealth University, Richmond, VA 23284, USA Department of Physics, Virginia Commonwealth University, Richmond, VA 23284, USA d Department of Physics and Astronomy, University of Pittsburgh, Pittsburgh, PA 15260, USA b c
a r t i c l e in f o
a b s t r a c t
Article history: Received 19 September 2008 Received in revised form 10 December 2008 Accepted 22 December 2008 Communicated by M. Skowronski Available online 6 January 2009
High-temperature m-plane AlN nucleation layers have been used for the growth of planar GaN films by metalorganic chemical vapor deposition on (1 0 1¯ 0) m-plane 6H-SiC substrates. Structural studies using transmission electron microscopy reveal the presence of a novel AlN intermediary layer preceding the remainder of the 2H-AlN buffer layer. High-resolution observations and image simulations indicate that this initial AlN layer has a faulted hexagonal structure with a six-layer repeating stacking sequence of yCBCACBCBCACBy along the transverse [0 0 0 1] direction, which does not replicate the underlying 6H-SiC stacking. Based on image analysis, the space group of this novel phase is tentatively identified as P3m1, which is also hexagonal. A structural model of the m-plane AlN/SiC interface with periodic misfit defects is also proposed. Charge neutrality analysis indicates that the interface has an equal mixture of C–Al and Si–N bonds. & 2008 Elsevier B.V. All rights reserved.
PACS: 61.72.Ff 68.37.Og 81.05.Ea Keywords: A1. Characterization A3. Metalorganic chemical vapor deposition B1. Nitrides
1. Introduction The growth of m-plane GaN-based heterostructures in the nonpolar [1 0 1¯ 0] direction has received considerable attention [1–3]. Nitride films having this particular orientation are free from the large spontaneous and piezoelectric polarization fields along the growth direction that are inherent to the usual c-plane orientation. Several groups have reported the growth of m-plane GaN on g-LiAlO2 (1 0 0) or m-plane SiC substrates using metalorganic vapor deposition (MOCVD), molecular beam epitaxy (MBE) or hydride vapor phase epitaxy [1–4]. For m-plane GaN films grown on m-plane SiC substrates, a thin AlN nucleation layer is commonly employed to improve the structural quality of the subsequent GaN layer [1]. Although the interplanar mismatch between AlN and 6H-SiC is less than 1% along both [0 0 0 1] and [1 0 1¯ 0] directions [5], most m-plane nitride films still have high densities of basal-plane stacking faults (BSFs) when no epitaxial lateral overgrowth technique is employed [1,3]. These BSFs usually originate from the substrate/film interface and have low formation energy, and they are likely to have an adverse impact on the optical properties of the material since they are electrically active Corresponding author. Tel.: +1 480 965 9815.
E-mail address: [email protected] (L. Zhou). 0022-0248/$ - see front matter & 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.jcrysgro.2008.12.047
[3,6]. Recently, the epitaxial growth of a 6H-AlN polytype on m-plane 6H-SiC by MBE was reported [7]. The presence of a 6HAlN polytype inclusion in GaN grown on GaAs (0 0 1) has also been reported [8]. However, no detailed description of this 6H-AlN structure was provided. In this present paper, we present highresolution observations of a six-layer intermediary AlN phase observed at the m-plane AlN/6H-SiC interface. Structural analysis combined with image simulations revealed that the stacking sequence of this AlN layer did not exactly replicate the 6H polytype structure of the substrate. The detailed atomic structure of the AlN/SiC interface was also determined.
2. Experimental details The m-plane GaN films investigated in this study were grown on (1 0 1¯ 0) m-plane 6H-SiC substrates (less than 0.11 miscut as measured by X-ray diffraction) using MOCVD. Prior to growth, the SiC substrate was annealed ex situ in H2 atmosphere for 10 min at 1500 1C to remove the polishing-induced surface damage. Following in situ annealing of the chemically cleaned SiC substrate, a 100-nm-thick high-temperature AlN nucleation layer was deposited at 1050 1C. The temperature was then lowered to 1030 1C for subsequent growth of the m-plane GaN epilayer. Trimethylgallium
ARTICLE IN PRESS L. Zhou et al. / Journal of Crystal Growth 311 (2009) 1456–1459
(TMGa) and ammonia (NH3) were used as the Ga and N sources, respectively. The flow rates of TMGa and NH3 were kept at 117 mmol/min and 550 sccm, respectively, with the growth pressure kept at 30 Torr. Transmission electron microscopy (TEM) observations were carried out on cross-sectional samples prepared by standard mechanical polishing and final Ar ion-beam-thinning at 4.5 keV. Electron micrographs were recorded with a JEM-4000EX highresolution electron microscope (HREM) equipped with a topentry, double-tilt specimen holder and operated at 400 keV. Relevant instrumental parameters used for image simulations include: spherical aberration coefficient, Cs=1.00 mm; beam divergence, div=0.50 mrad; and focal spread D=7 nm. Photographic negatives were scanned digitally to facilitate image processing, and the molecular simulation program CERIUS was used for image simulations.
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101 0 GaN 1120 0001
AlN
5nm
3. Results and discussion Cross-sectional TEM images of the m-plane GaN/AlN/SiC heterostructure revealed threading dislocations as well as high densities of BSFs in both AlN and GaN layers. Fig. 1 is a transmission electron micrograph taken under two-beam conditions with diffraction vector g=(0 0 0 2) showing threading dislocations inside the film. The density of the threading dislocation is highest near GaN/AlN interfaces, and then decreases toward the surface of the GaN layer. The BSF density in the GaN layer well away from the heavily faulted AlN/GaN interface was estimated to be on the order of 105 cm 1. Closer examination of the AlN/GaN interface showed that many BSFs in the AlN layer continued straight through into the GaN layer. However, a significant fraction terminated at the AlN/GaN interface, whilst others originated at the interface and propagated into the GaN layer. These features are more clearly visible in the enlarged image shown in Fig. 2. Fig. 3 shows a high-resolution electron micrograph from the area of the AlN/6H-SiC interface, taken with the incident electron beam oriented along the [11 2¯ 0] zone axis. Corresponding fast Fourier transform (FTT) diffraction patterns from the boxes labeled as A, B and C are shown in Fig. 4(a), (b), and (c), respectively. Fig. 4(a) is taken from an area which includes both 6H-SiC and the bottom part of the AlN layer. Only one set of spots
Fig. 2. Cross-sectional transmission electron micrograph taken along [11 2¯ 0] zone axis showing enlarged region of AlN/GaN interface. Note BSFs present in both AlN and GaN as well as discontinuities across the interface.
1010
1120 0001 AlN
6H-SiC 6H-SiC
Fig. 3. Cross-sectional transmission electron micrograph taken along [11 2¯ 0] zone axis showing the m-plane AlN/6H-SiC interface area. Areas labeled A, B and C are further analyzed in Fig. 4.
GaN
AlN
15nm
200nm
Fig. 1. Cross-sectional transmission electron micrograph of GaN/AlN/SiC heterostructure taken under two-beam conditions with diffraction vector g=(0 0 0 2) showing threading dislocations inside the GaN film.
is apparent confirming, as expected, that the lattice mismatch between AlN and 6H-SiC is small along [0 0 0 1] and [1 0 1¯ 0] directions. Fig. 4(a) also confirms the epitaxial growth of AlN on 6H-SiC. Fig. 4(b) is taken from the AlN layer (lateral dimensions 15 nm) immediately adjacent to the 6H-SiC substrate. This FFT pattern appears to be identical to that of 6H-SiC. The presence of five extra spots between (0 0 0 0) and (0 0 0 6) spots means that there are six layers in each AlN unit cell along the [0 0 0 1] horizontal direction, which is obviously different from the usual pattern for 2H-AlN (wurtzite) material. The AlN structure in this region might, therefore, be casually termed as ‘6H-AlN’ [7]. However, as shown below, structural analysis based on the image appearance indicates that the tetrahedral stacking across the AlN unit cell is different from the normal 6H hexagonal
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1010
1010
0000
0006
1010
0000
0006
0000
0002
Fig. 4. (a–c): Corresponding FFT diffraction patterns taken from areas in Fig. 3 labeled as A, B and C, respectively, showing that the AlN immediately adjacent to the m-plane 6H-SiC has a different crystalline structure from the remaining AlN layer.
stacking sequence. The streaking of diffraction spots visible along the [0 0 0 1] direction is due to the presence of some BSFs in this region. Fig. 4(c) is an FFT diffraction pattern taken from a region of the AlN that is located 30 nm away from the AlN/6H-SiC interface. The characteristic spot pattern expected for wurtzite 2H-AlN is visible, but no spots from the ‘6H-AlN’ phase are visible. Streaking due to BSFs in the analyzed area is also apparent. Identification of the structure of this thin ‘6H-AlN’ layer is an important step towards a full understanding of the initiation of AlN growth on m-plane 6H-SiC. Previous high-resolution studies of a GaN/SiC heterointerface showed that it was possible to identify the tetrahedral stacking sequence in both of these tetrahedrally based materials [9]. A high-resolution TEM image of the AlN/SiC interface taken along the [11 2¯ 0] zone axis is shown at high magnification in Fig. 5. The characteristic yABCACBABCACBy stacking sequence along the [0 0 0 1] direction of 6H-SiC is clearly visible. Close examination of the ‘6H-AlN’ layer indicates, however, that it has a six-layer yCBCACBCBCACBy stacking sequence. This stacking sequence clearly does not replicate the stacking sequence for a 6H hexagonal polytype along the [0 0 0 1] direction, as we elaborate further below. The circles highlighted the positions where the misfits in the stacking sequence between SiC and ‘6H-AlN’ occur. In order to confirm the crystal structure of the ‘6H-AlN’, through-thickness-through-focal image simulation tableaus were generated for both the 6H-SiC and the ‘6H-AlN’. The lattice parameters used for 6H-SiC were a: 3.081 A˚ and c: 15.092 A˚ [5]. Since the structure of ‘6H-AlN’ has not previously been reported, the lattice parameters used were based on those of the 2H-AlN structure, giving values of 3.114 and 14.9376 A˚ (3 4.9792 A˚) along the a and c directions, respectively. The corresponding image simulations (defocus of 450 A˚ and crystal thickness of 60 A˚) are shown superimposed on the experimental image. The close agreement between simulated and experimental images confirms both the microscope operating conditions and the proposed stacking sequence of the interfacial AlN layer. From these results, it appears that the ‘6H-AlN’ layer acts as a metastable buffer layer for transition from the 6H stacking of 6HSiC to the 2H stacking of AlN, presumably as a means to reduce strain caused by the small lattice mismatch and chemical differences. This situation is different from the growth of stable polytype-replicated 6H-AlN on m-plane SiC by MBE [7]. In effect, the ‘6H-AlN’ replicates closely the 6H-SiC crystal structure, as reported previously for 4H-AlN grown on 4H-SiC [10–12]. From a careful comparison with the stacking sequence of 6H-SiC, it seems that the ‘6H-AlN’ has a periodic stacking fault per unit cell which is introduced by shifting a layer of atoms from the A to C position (ABCACB to CBCACB). Alternatively, the ‘6H-AlN’ might be
Fig. 5. High-resolution transmission electron micrograph of the m-plane AlN/6HSiC interface taken along [11 2¯ 0] zone axis, with matching image simulations superimposed within rectangular boxes. Circles indicate positions of stacking defects between ‘6H-AlN’ and SiC.
described as a highly defective 2H structure with stacking disorder every third 2H unit cell (CBCBCB to CBCACB). Based on the determined stacking sequence of this ‘6H-AlN’, it is interesting to speculate on the likely space group. According to Verma and Krishna [13], there are only four possible space groups for SiC, and therefore AlN, namely, P3m1, R3m1, P63mc and F43m. The latter corresponds to the cubic 3C polytype, and R3m1 corresponds to the various rhombohedral polytypes. Since this new ‘6H-AlN’ does not have screw axis symmetry (rotation combined with half the c-axis period displacement) then P63mc is also eliminated. The remaining possibility for the space group is P3m1 which is also a hexagonal space group.
ARTICLE IN PRESS L. Zhou et al. / Journal of Crystal Growth 311 (2009) 1456–1459
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be a stable configuration alone. However, an equal mixture of these two possible bonding states would have overall charge neutrality. The Al atoms can be treated as acceptors residing on the Si sublattice of SiC. Alternatively, the C atoms can be treated as acceptors residing on the N sublattice of AlN [5]. By inserting these two bonding states into the structural model at the same time, charge neutrality could then be achieved for the structures drawn in Fig. 6(b) and (c). For example, in Fig. 6(b), there are nine Si–N bonds and one C–Al bond which would give 94 extra electron charge (14 per Si–N bond) and 14 deficit electron charge, respectively. Moreover, the atoms labeled C* and Si* are only bonded by three atoms, which would thus give another charge deficit of 1 electron. Thus, the overall total charge at the interface for this model is 9/4+( 1/4)+( 1 2)=0, and is therefore feasible. Further analyses of these interface structures are in process and will be reported at a later date.
0001
Fig. 6. (a) Schematic diagram showing the proposed atomic arrangements at the m-plane AlN/6H-SiC interface. (b) and (c) Possible atomic arrangements at the m-plane AlN/6H-SiC interface.
Because 2H stacking is the stable AlN crystalline structure, it might be anticipated that the intermediary ‘6H-AlN’ will transform to 2H-AlN with increasing film thickness away from the interface. This structural transition can be realized by introducing some low-energy type-I BSFs [14] into the 6H crystal, for example, from CBCACB to CBCACA or from CBCACB to CBCBCB. The first configuration will result in a type-I BSF in the 2H-AlN layer. Thus, there would be a high density of BSFs present in the subsequent 2H-AlN layer, which is consistent with the high-resolution images and the FFT pattern shown in Fig. 4(c). Further structural analysis of the m-plane ‘6H-AlN’/6H-SiC interface was carried out. First, an atomic-scale structural model was constructed based on the structures of 6H-SiC, ‘6H-AlN’ and the appearance of the high-resolution TEM images. A schematic diagram of the structural model of the interface projected along the [11 2¯ 0] direction is shown in Fig. 6(a). In this model, all atoms at the interface are bonded by four atoms, except for the two atoms labeled with asterisks, which are bonded only by three atoms. Next, we consider the possible chemical bonding at the ‘6H-AlN’/6H-SiC interface. There are four possible bonding configurations at the interface: Si–C–N–Al, Si–C–Al–N, C–Si– Al–N and C–Si–N–Al, as described previously [5]. Based on previous structural models of the c-plane AlN/SiC interface, the relationship between the corresponding bond lengths (d) can be expressed: dSi–Al4dC–AldSi–N4dC–N. Close examination of Fig. 5 shows that there are no abrupt changes in the interplanar separation between any two lattice planes along the m-direction. Thus, the –C–N–, and –Si–Al– bonds can be eliminated from consideration. While both –C–Al– and –Si–N– bonds are consistent with our observations, they cannot be differentiated by TEM examination. Moreover, charge neutrality is essential for the stability of the interface [5]. In a tetrahedral bonding configuration, both Si and C will contribute one electron per bond, whereas Al and N will contribute 34 and 54 electrons per bond, respectively [5]. It is obvious that neither the –C–Al– (charge deficit of 14 electron per bond), nor the –Si–N– (charge excess of 14 electron per bond) could
4. Conclusions A thin six-layered AlN intermediary layer, tentatively identified as having the P3m1 space group, has been observed at the m-plane AlN/6H-SiC interface using high-resolution electron microscopy. This layer facilitates the transition from the 6H-SiC substrate to the 2H-AlN hexagonal structure. Image analysis shows that the ‘6H-AlN’ has a faulted yCBCACBCBCACBy stacking sequence along the [0 0 0 1] direction, rather than the expected 6H hexagonal arrangement. The 2H-AlN layer has a high density of BSFs. Further analysis of the m-plane AlN/6H-SiC interface indicates that it must have an equal mixture of C–Al and Si–N bonds.
Acknowledgments This work was partially supported by ONR Grant N-00014-041-0020. The work at Virginia Commonwealth University was supported by the Air Force Office of Scientific Research. We acknowledge use of facilities in the John M. Cowley Center for High-Resolution Electron Microscopy at Arizona State University. References [1] T. Kawashima, T. Nagai, D. Iida, A. Miura, Y. Okadome, Y. Tsuchiya, M. Iwaya, S. Kamiyama, H. Amano, I. Akasaki, J. Crystal Growth 298 (2007) 261. [2] P. Waltereit, O. Brandt, A. Trampert, H.T. Grahn, J. Menniger, M. Ramsteiner, M. Reiche, K.H. Ploog, Nature 406 (2000) 865. [3] B.A. Haskell, T.J. Baker, M.B. McLaurin, F. Wu, M.P.T. Fini, S.P. DenBaars, J.S. Speck, S. Nakamura, Appl. Phys. Lett. 86 (2005) 111917. [4] M. McLaurin, T.E. Mates, J.S. Speck, Appl. Phys. Lett. 86 (2005) 262104. [5] F.A. Ponce, C.G. Van de Walle, J.E. Northrup, Phys. Rev. B 53 (1996) 7473. [6] R. Kro¨ger, T. Paskova, S. Figge, D. Hommel, A. Rosenauer, Appl. Phys. Lett. 90 (2007) 081918. [7] D.M. Schaadt, O. Brandt, A. Trampert, H.-P. Scho¨nherr, K.H. Ploog, J. Crystal Growth 300 (2007) 127. [8] M. Funato, T. Ishido, S. Fujita, S. Fujita, Appl. Phys. Lett. 76 (2000) 330. [9] J.N. Stirman, F.A. Ponce, A. Pavlovska, I.S.T. Tsong, D.J. Smith, Appl. Phys. Lett. 76 (2000) 822. [10] N. Onojima, J. Suda, T. Kimoto, H. Matsunami, Appl. Phys. Lett. 83 (2003) 5208. [11] M. Horita, J. Suda, T. Kimoto, Appl. Phys. Lett. 89 (2006) 112117. [12] M. Horita, T. Himoto, J. Suda, Appl. Phys. Lett. 93 (2008) 082106. [13] A.R. Verma, P. Krishna, Polymorphism and Polytypism in Crystals, Wiley, New York, 1966 p. 160. [14] J. H. Edgar, S. Strite, I. Akasaki, H. Amano, C. Wetzel, Properties, Processing and Applications of Gallium Nitride and Related Semiconductors, INSPEC, 1999.