J. Phys. Chem. C 2007, 111, 7889-7892
7889
Self-Organization of 3D Triangular GaN Nanoislands and the Shape Variation to Hexagonal Zhilai Fang* and Junyong Kang Semiconductor Photonics Research Center and Department of Physics, Xiamen UniVersity, Xiamen 361005, People’s Republic of China ReceiVed: March 5, 2007; In Final Form: April 10, 2007
We report on the self-organization of large-scale uniform aligned three-dimensional (3D) GaN islands with distinct triangular (0001) and smooth side facets and the shape variations of the (0001) facets from triangular to hexagonal during metalorganic vapor-phase epitaxy (MOVPE) growth of GaN films on Si-rich SiNx patterned sapphire substrates. The triangular island shaping during the recrystallization processes of GaN nucleation layers (NLs) can be attributed to the enhanced diffusion and regrowth anisotropy. The island shape transition from triangular to hexagonal in the early stages of high-temperature growth of GaN epilayers is due to the gas-phase transport dominating growth mechanism and the limited diffusion length of edge adatoms compared with the increased island size.
I. Introduction Heteroepitaxy often leads to the formation of 3D islands in early growth stages, which has significant influence on film qualities. The stringent requirements for high-quality films needed for developing high-performance and future miniature devices have motivated intense research interests on the island shaping and shape variations.1-4 However, the island shaping is critically dependent on materials systems, and previous studies are mostly on metal and SixGe1-x based materials. The formation of faceted 3D GaN islands of triangular (0001) facets and the shape variations from triangular to hexagonal are not yet observed and well understood. The exploration of the GaN 3D island shaping and shape variations and thus discovery of the growth mechanisms are of general scientific significance as well as technological application significance and enormous economic benefits.5-8 Recently, the topic of self-organization of controlled nanocrystalline 3D shaped islands has become an important aspect of research because of the technological significances as well as the potential applications in fabrication of ordered arrays of nanoscale devices.9,10 In this paper, by purposeful employment of Si-rich SiNx treatment to the sapphire substrates and control of the initial GaN growth conditions, we have successfully realized the formation of large-scale uniform aligned GaN nanocrystalline islands. By studying the influences of the Sirich SiNx treatment on surface morphology of the subsequently grown GaN films, we propose kinetic mechanisms for the triangular island shaping and the shape variation processes from triangular to hexagonal for GaN films on the Si-rich SiNx patterned sapphire substrates. II. Experimental Methods The epitaxial growth of GaN and in-situ deposition of SiNx layers were carried out in a close-coupled showerhead (CCS) 3 × 2′′ reactor of a metalorganic chemical vapor deposition * To whom correspondence should be addressed. Tel: +86-592-2184220. Fax: +86-592-2187737. E-mail:
[email protected].
(MOCVD) system. The (0001) sapphire substrate surfaces were prepared by thermal cleaning at 1060 °C for 8 min followed by nitridation at 550 °C for 4 min. Before the growth of GaN epitaxial layers, an incomplete SiNx layer was in situ predeposited, and then a conventional 25 nm low-temperature (LT) GaN nucleation layer (NL) was grown at 535 °C and 500 Torr on the nitridated sapphire substrates followed by a hightemperature (HT) recrystallization process. The SiNx deposition was performed by introducing SiH4 (100 ppm, 20-40 sccm) and NH3 (2000 sccm) into the reactor simultaneously with H2 as the carrier gas (5500 sccm). Growth at low temperature with low NH3/SiH4 ratio and H2 ambient was used for patterning Si-rich SiNx nanoislands on the substrate surface. Trimethylgallium (TMGa) and high-purity ammonia were used as the source precursors for deposition of LT GaN NLs under Garich growth conditions and with hydrogen as the carrier gas. After recrystallization of the LT GaN NLs, epitaxial growth of HT GaN films was performed at 1035 °C and 300 Torr with low V/III ratios in the early stages. To investigate the island shape and shape variation processes, different growth stages of GaN films were prepared by varying growth time, temperature, pressure, and V/III ratio. The surface morphologies of SiNx layers, LT GaN NLs, and HT GaN films at different growth stages were investigated by atomic force microscopy (AFM, PicoSPM and PSI XE-100) and scanning electron microscopy (SEM, LEO1530) equipped with an energy-dispersive X-ray spectrometer (EDX). The surface elemental information was characterized by X-ray photoelectron spectroscopy (XPS, PHI Quantum2000) and SEM-EDX. The crystal structure was characterized by high-resolution X-ray diffraction (XRD, PHILIPS X’Pert PRO). III. Results and Discussion The in-situ predeposited SiNx layers exhibited islandlike surface structure with island size of ∼100 nm, height of ∼2 nm, and density of ∼1.6 × 109 cm-2 as shown in Figure 1a. The formation of islandlike structure is due to the application of low NH3/SiH4 ratio and high flow rate of H2 carrier gas (antisurfactant effect). By doing SEM-EDX measurements, we found that at the island sites Si was detectable whereas at the
10.1021/jp071803c CCC: $37.00 © 2007 American Chemical Society Published on Web 05/15/2007
7890 J. Phys. Chem. C, Vol. 111, No. 22, 2007
Figure 1. (a) The surface morphology of thin SiNx layers predeposited on sapphire substrate. (b) The Si2p X-ray photoelectron peak of the SiNx layers coated on sapphire.
Figure 2. The surface morphology of the ∼25 nm LT GaN NLs grown on the SiNx layers before annealing. The height of the surface was scanned over 2.5 µm to show a surface roughness of ∼2 nm.
other sites Si was not detected.11 Therefore, we conclude that the sapphire substrate surface was partly covered by the SiNx islands. The surface chemistry of the in-situ patterned SiNx layers on the nitridated sapphire substrates was also examined by XPS studies. Figure 1b shows the XPS spectrum of the Si2p photoelectron peak. The peak can be fitted into three peaks “P1”, “P2”, and “P3” corresponding to SiAl, Si0, and SiNx, respectively. By comparing the fitted peak areas, we find that the SiNx layers were Si-rich with a Si/N ratio of about 3. Subsequent growth of GaN NLs was carried out at low temperature. The very thin SiNx layers (2 nm) were continuously covered by the comparatively thick GaN NLs (25 nm) with a roughness of 2 nm as observed from the line scan in Figure 2. After annealing (recrystallization), the surface morphology greatly changed and showed structures of large-scale aligned triangular 3D nanocrystalline islands of distinct and smooth sidewall faceting as shown in Figure 3. We estimate the island size of ∼250 nm, island height of ∼20 nm, and island density of ∼1.4 × 109 cm-2. The island density is close to that of the predeposited SiNx nanoislands, and thus we infer that GaN nucleated on the bare sapphire substrate surface surrounding the Si-rich SiNx islands during the recrystallization process (after recrystallization, the very thin SiNx islands were fully covered
Fang and Kang
Figure 3. (a) The surface morphology of the ∼25 nm LT GaN NLs grown on the SiNx layers after annealing. (b) The 3D image of the GaN islands shows the distinct and smooth sidewall faceting. In the figure, the crystallographic alignment of the triangular base is derived from the substrate.
by the relatively thick GaN islands). This judgment has been further supported by SEM-EDX experiments in which silicon was detectable at the island sites, whereas it was out of detection limit at the other sites.11 The incorporation of Si into GaN bulk is expected to reduce the compressive strain for GaN on sapphire,12,13 and thus GaN preferentially nucleates around the Si-rich SiNx islands. Further studies will be carried out to in situ control the island density and the dispersion of island sites via controlling the Si-rich SiNx predeposition. In the 2D island shaping, formation of triangular islands is possible on films of either fcc(111)2,3 or hcp(0001) phase.14 XRD studies show that the annealed LT GaN NLs are of hexagonal nature.11 Different from the formation of small nonuniform triangular 2D islands during molecular beam epitaxy (MBE) of GaN films,14 the self-organization of triangular 3D islands occurred in the annealing processes of GaN NLs prepared on in-situ Si-rich SiNx patterned sapphire by MOVPE. The distinct triangular islands are uniform and are aligned with sharp triangle corners (Figure 3). During the annealing processes of GaN NLs, GaN decomposition is prevalent and the desorbed Ga atoms may reincorporate to nuclei through gas-phase transport to form newly shaped islands.15,16 The adatoms attached to the island edges may also effectively migrate along the island peripheries especially when the island size is small (less than hundreds of nanometers) and the Ga adatom diffusion is enhanced.17,18 Therefore, the island shaping is determined by the competition between the growth rate of island edge facets (via decomposition and reincorporation) and the Ga adatom diffusion around the island edges. To further understand the kinetic mechanism for the triangular island shaping, we refer to the atomic model of wurtzite GaN (Figure 4). In the (2h110) cross-sectional view (Figure 4a) and (0001) plan view (Figure 4b) of wurtzite GaN, we denote two types of bilayer edges as I and II characterized by single- and double-dangling bonding per edge atom, respectively. It has been found that local atomic-bonding geometry has a significant effect on the growth rate of island edges and the diffusion rate along island edges.4,18,19 According to the atomic model and considering the binding energy between adatoms and island edges, it is obvious that the type II edges have higher rate to capture adatoms whereas the type I edges have higher rate to decompose, that is, there exists anisotropy of decomposition and reincorporation of edge adatoms. However, triangular islands are seldom observed in MOVPE growth of GaN films.11 This indicates that the decomposition and reincorporation anisotropy is generally not very effective for formation of triangular islands. Considering another competition mechanism for island shaping, similarly because of different binding energies, there exists
Self-Organization of 3D Triangular GaN Nanoislands
J. Phys. Chem. C, Vol. 111, No. 22, 2007 7891
Figure 5. (a) SEM image of the surface morphology of GaN films showing different intermediate forms and coalescent states of GaN islands. (b) Sketch of the intermediate forms for the GaN epitaxial growth. In the figure, “T” denotes the triangular shape, “M1” quasitriangular, “M2” quasi-hexagonal, “H” hexagonal, and LI (LII) the length of the type I (II) edges. (c) Enlarged view of the island M showing quasi-triangular shape at the small top (0001) facet and quasi-hexagonal shape at the large bottom (0001) facet. (d) Schematic representation of the 3D island M.
Figure 4. (a) Cross-sectional view of wurtzite GaN showing two types of bilayer edges I and II characterized by edge atoms of single- and double-dangling bonding, respectively. (b) Plan view of wurtzite GaN showing the geometry of the type I and II edges. The solid lines and dash lines surrounding the hexagonal cell denote type I and II edges, respectively. In the figure, gI denotes the growth rate of the type I edges and gII denotes that of the type II edges.
diffusion anisotropy along the island edges.18 Under appropriate conditions, for example, enhanced diffusion and small island size, adatoms at type I edges will be able to effectively migrate along the edges and be trapped by type II edges. As a result, type I edges will extend and type II edges will fade out, which will eventually form triangular (0001) facets. The reduction of the adatom diffusion barrier can be caused by an increase of the “surface Ga” coverage and can be realized by the Ga-rich growth conditions and Si-Ga exchange processes.17 Furthermore, Ga metals can catalyze H2 and form Ga-Hx species, which are more mobile. Therefore, in our studies, we intentionally employed Ga-rich growth conditions and Si-rich SiNx pretreatment to sapphire for enhancement of adatom diffusion. Under appropriate growth and annealing conditions, selforganization of distinct uniform aligned triangular GaN islands was successfully realized (Figure 3). The observation of smooth sidewall facets (see Figure 3b) further supports the importance of enhanced adatom diffusion for island shaping at small length scales. Actually, diffusion anisotropy along island periphery modifies the growth rates of island edges and thus enhances the growth anisotropy of island edges. Therefore, we attribute the kinetic mechanism for the formation of triangular islands to (Ga adatom) diffusion-enhanced regrowth anisotropy. Subsequent growth of GaN films on the annealed NLs was carried out at HT with initial low V/III ratios. The isolated islands coarsened with lateral expansions followed by coalescence of the isolated islands. To investigate the processes of island coarsening, shape variations, and coalescence, different
growth stages of GaN films were prepared by variation of growth time from 0 to 400 s. When islands coarsened, for example, island size increased to 1 µm, diffusion along island edges would become much less effective in island shaping because of the limited diffusion length. Accordingly, for large island size and HT growth conditions where the gas dissociation kinetics is very fast, the growth rates determined by the gasphase transport rates of island edges become dominant in island shaping and shape variations. For a CCS planetary reactor used in our MOCVD system, there are the same mass transport rate and thus the same growth rate (i.e., there is gI ) gII) for the type I and II island edges. Consequently, as triangular island coarsening proceeds, the type II edges are extended and the shape variation processes from triangular to hexagonal gradually is initiated. The above predictions have been supported by the observation of both islands H and B in Figure 5a. In the figure, the islands marked as B can be expressed by quasi-triangular M1 or quasi-hexagonal M2 and the island H by hexagonal H as schematically drawn in Figure 5b. The observation of forms M1, M2, and H definitely confirms the conversion processes of island shape from triangular to hexagonal. Experimentally, by measuring the lengths LI (LII) of the type I (II) edges for the (0001) facets of different sizes, we can quantitatively estimate the relationship between the ratio of LII to LI and the size of the facets. According to the experimental data, we schematically sketch in Figure 5b the intermediate forms for GaN growth as triangular T, quasitriangular M1, quasi-hexagonal M2, and hexagonal H. The growth rate of the type I (II) edges gI (gII) is proportional to the distance between the origin (center) and the type I (II) edge faces. Apparently, when island coarsened at the HT growth stages, growth anisotropy between gI and gII was weakened, and thus the difference between LI and LII diminished. As a result, with the increase of island size the shape of the (0001) facets varied from the form T to M1 or M2 and eventually to H. In Figure 5c, we present a close-up view of the 3D island M.
7892 J. Phys. Chem. C, Vol. 111, No. 22, 2007 The 3D form of the island can be schematically expressed as in Figure 5d. As we have discussed above for the (0001) facets, the island shape variations from triangular T to the intermediate forms and eventually hexagonal H occurred with the increase of island size. Accordingly, under proper conditions, quasitriangular shape (M1) was observed for the small (0001) top facet of the island whereas quasi-hexagonal shape (M2) was observed for the large bottom facet as shown in Figure 5c. This has further supported that diffusion (anisotropy) becomes less effective for changes of island shapes when the size of island facets become larger. In addition to the shape variation processes and island reshaping, both partly coalesced island C1 and fully coalesced islands C2/C3 were also observed in Figure 5a, which had manifested the coalescence processes. IV. Conclusion In summary, we have reported the self-organization of largescale aligned nanocrystalline 3D GaN islands of triangular (0001) base and the shape variations of the (0001) facets from triangular to hexagonal during MOVPE growth of GaN on Sirich SiNx patterned sapphire substrates in the early stages. We found that during annealing of nucleation layers, it was the enhanced diffusion and regrowth anisotropy responsible for the triangular island shaping. Further growth at high temperature resulted in island coarsening with gas-phase transport dominating growth mechanism. The equal gas-phase transport rate for both types of island edges and the limited diffusion length compared with the large island size resulted in the shape transition of the (0001) facet from triangular to hexagonal. Acknowledgment. This work was supported by the “863” program (2006AA03A110), National basic research program (A1420060155), and Xiamen science & technology program (3502Z20063001) of China. We thank W.J. Huang and H.T. Sun for AFM measurements, S.J. Wang and H.Z. Sun for XPS
Fang and Kang measurements, and C. Li, M. Lu, Q.H. Wu, B.P. Zhang, and J.C. Li for helpful discussions. Supporting Information Available: SEM-EDX data of thin SiNx layers and ∼25 nm LT GaN NLs and XRD 2θ-scan data of the annealed LT GaN NLs. This material is available free of charge via the Internet at http://pubs.acs.org. References and Notes (1) Zhang, Z. Y.; Lagally, M. G. Science 1997, 276, 377. (2) (a) Kodambaka, S.; Khare, S. V.; Petrov, I.; Greene, J. E. Surf. Sci. Rep. 2006, 60, 55. (b) Evans, J. W.; Thiel, P. A.; Bartelt, M. C. Surf. Sci. Rep. 2006, 61, 1, and references therein. (3) Michely, T.; Hohage, M.; Bott, M.; Comsa, G. Phys. ReV. Lett. 1993, 70, 3943. (4) Zhang, Z. Y.; Chen, X.; Lagally, M. G. Phys. ReV. Lett. 1994, 73, 1829. (5) Ponce, F. A.; Bour, D. P. Nature 1997, 386, 351. (6) Nakamura, S. Science 1998, 281, 956. (7) Jain, S. C.; Willander, M.; Narayan, J.; Overstraeten, R. Van J. Appl. Phys. 2000, 87, 965. (8) Akasaki, I. J. Cryst. Growth 2002, 237, 905. (9) Deng, X.; Weil, J. D.; Krishnamurthy, M. Phys. ReV. Lett. 1998, 80, 4721. (10) Wang, L. G.; Kratzer, P.; Scheffler, M. Phys. ReV. Lett. 1999, 82, 4042. (11) See Supporting Information. (12) Rosa, A. L.; Neugebauer, J.; Northrup, J. E.; Lee, C. D.; Feenstra, R. M. Appl. Phys. Lett. 2002, 80, 2008. (13) Chine, Z.; Rebey, A.; Touati, H.; Goovaerts, E.; Oueslati, M.; Jani, B. El; Laugt, S. Phys. Status Solidi A 2006, 203, 1954. (14) Xie, M. H.; Seutter, S. M.; Zhu, W. K.; Zheng, L. X.; Wu, H. S.; Tong, S. Y. Phys. ReV. Lett. 1999, 82, 2749. (15) Koleske, D. D.; Coltrin, M. E.; Cross, K. C.; Mitchell, C. C.; Allerman, A. A. J. Cryst. Growth 2004, 273, 86. (16) Mitchell, C. C.; Coltrin, M. E.; Han, J. J. Cryst. Growth 2001, 222, 144. (17) Rosa, A. L.; Neugebauer, J. Phys. ReV. B 2006, 73, 205314. (18) Xie, M. H.; Gong, M.; Pang, E. K. Y.; Wu, H. S.; Tong, S. Y. Phys. ReV. B 2006, 74, 085314. (19) Du, D.; Srolovitz, D. J.; Coltrin, M. E.; Mitchell, C. C. Phys. ReV. Lett. 2005, 95, 155503.
