Journal of ELECTRONIC MATERIALS, Vol. 41, No. 5, 2012
DOI: 10.1007/s11664-012-1967-z Ó 2012 TMS
Effects of Growth Temperature on Indium Incorporation in InAlN Alloys Grown by GSMBE on Si(111) MD RAKIB UDDIN,1,2,5 MAHESH PANDIKUNTA,1,2 VLADIMIR MANSUROV,1,2 SANDEEP SOHAL,1,3 DENIS MYASISHCHEV,1,3 GEORGIY M. GURYANOV,4 VLADIMIR KURYATKOV,1,2 MARK HOLTZ,1,3 and SERGEY NIKISHIN1,2 1.—Nano Tech Center, Texas Tech University, Lubbock, TX 79409, USA. 2.—Department of Electrical and Computer Engineering, Texas Tech University, Lubbock, TX 79409, USA. 3.—Department of Physics, Texas Tech University, Lubbock, TX, USA. 4.—Corning Inc., Corning, NY 14831, USA. 5.—e-mail:
[email protected]
InxAl1xN alloys with low indium content (0.025 < x < 0.080) were grown on Si(111) substrates, with an AlN buffer layer, using gas source molecular beam epitaxy with ammonia under nitrogen-rich conditions. Composition was varied by changing the growth temperature from 580°C to 660°C. Growth temperature in excess of 580°C was found to be necessary to obtain compositional uniformity. As temperature was varied from 590°C to 660°C, both the growth rate and indium incorporation decreased substantially. Rising In content observed near the surface of each sample was attributed to native indium oxide formation. Key words: GSMBE with ammonia, InAlN, crystalline quality, sixfold symmetry, optical bandgap, In incorporation
INTRODUCTION III-Nitride semiconductors have attracted attention for applications in areas such as optoelectronics, high-power and high-frequency electronics, and solar cells. The InAlN alloys are receiving significant attention due to potential applications for heterostructure field-effect transistors,1,2 bottom cladding layer in GaN-based edge-emitting lasers,3 electron-blocking layers in III-nitride visible lightemitting diodes,4 low-refractive-index layer in IIInitride distributed Bragg reflectors,5 and efficient materials for solar cell applications.6 These diverse applications have motivated investigations of different growth methods of these alloys. Approaches have included plasma-assisted molecular beam epitaxy (PAMBE),7,8 metalorganic chemical vapor deposition (MOCVD),9–11 and radio frequency (RF) magnetron sputtering.6,12 Much less is known about the growth of InAlN by gas source molecular beam epitaxy (GSMBE) with ammonia.13 Alloys have only (Received August 2, 2011; accepted January 27, 2012; published online February 24, 2012)
824
been demonstrated with InN concentrations to 2 mol.% using this method, primarily due to the low growth temperatures required and the resulting low NH3 decomposition. We report the compositional dependence of InxAl1xN with variation in growth temperature using GSMBE with NH3 on silicon substrates. Postgrowth characterization was carried out using high-resolution x-ray diffraction (XRD), secondaryion mass spectroscopy (SIMS),14 scanning electron microscopy (SEM), atomic force microscopy (AFM), optical reflectance, and spectroscopic ellipsometry (SE). Based on the growth and characterization results, we calculate the indium incorporation dependence on growth temperature and discuss a mechanism of In incorporation. EXPERIMENTAL PROCEDURES For the GSMBE growth, 50-mm-diameter Sbdoped n-type Si(111) wafers were used as substrates. Prior to loading in the vacuum chamber, substrates were prepared by wet chemical etching, yielding the formation of hydrogen-saturated Si
InAlN on Si(111)
825
Fig. 1. Schematic diagram showing the different steps of the epitaxial growth: (I) thermal precleaning of substrate from room temperature (TRT) to T1 = 630°C at a ramp of 10°C/min, (II) thermal cleaning including rapid heating (100°C/min) to T2 = 830°C and cooling (100°C/min) to T1, (III) nitridation of Si substrate and nucleation of AlN buffer layer at T1, (IV) heating (10°C/min) to T3 = 920°C for AlN growth, (V) AlN growth for 1 h, (VI) cooling (10°C/min) to InAlN growth temperatures ranging from the lowest TGL = 580°C to the highest TGH = 660°C, (VII) InAlN growth for 3 h, and (VIII) cooling (10°C/min) to RT for unloading.
surfaces.15 Epitaxial growth of InAlN consisted of eight steps with the temperature profile depicted in Fig. 1. Two steps were used for thermal cleaning in the growth chamber. Step I produced bare Si surface at around 630°C. The rapid temperature spike to 830°C in step II was used to drive off any Si–O which may remain while avoiding partial surface nitridation by background ammonia. To form Si-NAl interface, nitridation of Si and nucleation of AlN were carried out for several minutes at 630°C (step III). The details of this step are described elsewhere.16 Following the temperature ramp in step IV, the epitaxial AlN buffer layer was grown in step V for 1 h at 920°C with Al and NH3 precursors. The substrate was cooled in step VI to the alloy growth temperature. The InAlN alloys were grown for 3 h (step VII) at temperatures ranging from TGL = 580°C to TGH = 660°C with Al, In, and NH3 precursors. Five samples of InAlN, denoted A, B, C, D, and E, were grown at temperatures of 580°C, 590°C, 610°C, 630°C, and 660°C, respectively. All samples were grown at the same ammonia and metal fluxes under nitrogen-rich conditions. Beam-equivalent pressure (BEP) of Al and In sources were 1.17 9 107 Torr and 2.70 9 107 Torr, respectively. Ammonia flux was held at 70 standard cubic centimeters per minute (sccm). The substrate temperature was monitored using a backside noncontact thermocouple calibrated using three established temperatures in separate experiments: (1) melting temperature of indium on silicon substrate at 150°C, (2) temperature-induced reconstruction of hydrogen-terminated Si(111) surfaces from (1 9 1) to (7 9 7) at 650°C,15 and (3) temperature-induced reconstruction of bare Si(111) surface from (7 9 7) to (1 9 1) at 830°C.15 In each case, the transition was observed using reflection high-energy electron diffraction (RHEED). Film thicknesses were 100 nm for AlN and 500 nm to 600 nm for InAlN, as measured by
Fig. 2. Symmetric (0002) 2h–x rocking curves from XRD show peak positions of InAlN (left peaks) and AlN (right peaks) for each sample grown at different temperatures (a) and long-range asymmetric phi-scan XRD for sample C (b). (Color figure online). ð1011Þ
cross-section SEM. Growth rates ranged from 113 nm/h at 660°C to 193 nm/h at 590°C. RESULTS AND DISCUSSION Figure 2a shows 2h–x XRD rocking curves of the (0002) diffraction for each sample. The peak for each InxAl1xN sample shifts toward the AlN Bragg position with increasing growth temperature from sample A to E, confirming better In incorporation into the epitaxial layer at lower temperatures. Peak positions at 35.58°, 35.63°, 35.74°, 35.82°, and 35.88° correspond to InN concentrations of 8.0 mol.%, 7.1 mol.%, 5.3 mol.%, 3.7 mol.%, and 2.6 mol.% for samples A, B, C, D, and E, respectively. The (0002) diffraction from the AlN buffer layer is also present. Note that the AlN peaks in Fig. 2a, each measured at the center of the 50-mm wafer, exhibit slight deviation in position. This can be attributed to the different transition times from three-dimensional (3D) to two-dimensional (2D) growth mode of AlN arising from different degrees of initial nitridation of the Si substrate. The latter depends on the background pressure of ammonia in
826
Uddin, Pandikunta, Mansurov, Sohal, Myasishchev, Guryanov, Kuryatkov, Holtz, and Nikishin
the growth chamber, especially when the samples are grown one after another without removal of frozen ammonia from the cryopanels. We also found that the AlN peak position varied slightly from the edge to the center of each sample, corresponding to slightly higher degree of stress relaxation at the center of each wafer. One can expect that different residual stresses described above for the AlN buffer layers could influence both the residual stress in the InAlN films and the In incorporation. However, the angular peak position of InAlN films grown at 610°C and above did not vary across the wafers. We observed that the InAlN peak position followed that of the AlN peak only in the sample grown at 580°C, whereas the InAlN diffraction angle across the wafer for the sample grown at 590°C did not correlate with the AlN peak. For this sample, the InAlN peak position at the wafer center was shifted from the peaks obtained from the edge and at a position 12 mm from the center, which agreed well with each other. This shift can be attributed to higher relaxation of the InAlN film toward the wafer center, although decreasing In concentration may also contribute to this observed shift. We believe that variations in the angular position of diffraction from InAlN grown at 590°C and below are mostly due to decreasing Al surface diffusion efficiency at low temperatures. More detailed investigations are required to evaluate the influence of residual stress in AlN layers on composition and stress distribution in InAlN grown over these buffer layers. Figure 2b shows long-range asymmetric ð1011Þ phi-scan XRD for sample C. Peaks separated by 60° demonstrate that the InAlN films have high crystalline quality and the sixfold symmetry representative of the hexagonal wurtzite structure. Similar results were obtained for all samples. Narrow, intense (0002) peaks in Fig. 2a correspond to good crystalline quality of InAlN alloys grown at 660°C, 630°C, and 610°C. The asymmetric shape observed for the InAlN peaks of the samples grown at 580°C and 590°C are discussed below. The indium concentrations along the growth direction in samples C, D, and E were uniform. The SIMS profile obtained for sample D grown at 630°C is shown in Fig. 3a. AFM surface morphology of this sample is shown in Fig. 3b. The relatively high root-meansquare (RMS) surface roughness of 5.7 nm is attributed to low growth temperature, at which the Al surface diffusion is weak yielding a 3D growth mode for InAlN films. This is consistent with the 3D growth mode observed by RHEED, shown in Fig. 3c. The average InN concentration 4.3 ± 0.3 mol.% is in good agreement with the value from XRD measurements. The same good correlation between XRD and SIMS data was observed for samples C and E grown at 610°C and 660°C, respectively. The weak increase of InN concentration near the surface of sample D, seen in Fig. 3a, was also observed in other samples. The greatest rise was observed in
Fig. 3. SIMS profiles of Al and In for sample D grown at 630°C (a), 5 lm 9 5 lm ex situ AFM image (b) and in situ RHEED image for this sample (c).
sample A grown at 580°C, as shown in Fig. 4a, as discussed further below. The shapes of the XRD peaks in Fig. 2a for the samples grown at 580°C and 590°C are slightly asymmetric. This shape is most pronounced for sample A, as shown in Fig. 5a. To understand this asymmetry, we carried out simulations of the (0002) XRD peak shape assuming an increasing concentration of indium towards the surface. This assumption was based on previously published results.17,18 Simulation results shown in Fig. 5a exhibit good agreement with the XRD data. We tentatively attribute the absence of fringes in the data to the presence of structural defects in the thin InAlN films. Based on our simulations, the observed shape corresponds to 3 mol.% rise of InN concentration from the AlN/InAlN interface to the epilayer surface. This is consistent with the SIMS depth profile which shows increasing InN concentration from 5.1 mol.% at interface with AlN buffer to 8.7 mol.% near the surface, as shown in Fig. 5b. As mentioned above, the SIMS data in Fig. 4a show a small rise in Al concentration near the sample surface, along with a much larger rise in the In concentration. To check whether this accumulation was the result of our growth conditions, a bare Si(111) substrate was exposed to the same BEP of In used for the InAlN alloy growth for 2 h at substrate
InAlN on Si(111)
827
Fig. 4. SIMS profiles of Al and In near the surface of sample A grown at 580°C (a) and normalized oxide distributions (b).
temperature of 580°C. The ammonia injector and Al cell were off during this period. Background pressure of ammonia in the growth chamber was less than 109 Torr. No evidence of indium accumulation on the Si surface was observed using SEM. We conclude that the incoming In atoms were desorbed from the substrate surface at this temperature. As mentioned above, near-surface segregation of In was observed in all SIMS profiles, but it was most pronounced for sample A. To check the composition of the near-surface material, additional SIMS investigations were used to check for the presence of oxides of indium and aluminum based on the InO/O and AlO/O intensity ratios. As shown in Fig. 4b, the InO/O ratio confirms the enrichment of In at the film surface with thickness 1 nm. We believe that this is due to formation of the native oxide during the 2-month ambient storage of this sample. No clear evidence was obtained for the formation of the corresponding aluminum oxide. It is possible that the presence of the indium native oxide serves as a getter to enrich the In composition near the surface. Studies on freshly grown samples and following exposure to atmosphere for different durations are planned to check this hypothesis. Optical bandgap energies of the alloys with uniform In distribution, samples B, C, D, and E, were estimated from optical reflectance and spectroscopic ellipsometry measurements. Fabry–Perot interference fringes are present for light having subbandgap photon energy, where the InAlN is transparent, and vanish close to the optical bandgap energy as the material becomes opaque. The disappearance of
Fig. 5. Experimental and simulated symmetric (0002) 2h–x rocking curves for sample A grown at 580°C (a) and SIMS profiles of Al and In for this sample (b). (Color figure online).
the fringes is useful for obtaining the optical bandgap value.19 Figure 6 shows the bandgap energies versus InN mole fraction. Also included is our result for pure AlN. The observed dependence is in agreement with previously published experimental work.20 Recent theoretical calculations, based on local density approximations, consistently provide estimates below our measured results, although the theory describes reasonably well the observed trend for InAlN alloys when their curve is shifted by +0.28 eV to match our AlN bandgap.21 The narrow range of compositions investigated is insufficient to determine the bowing parameter; a linear fit is included to guide the eye. The InN concentration and growth rate are shown in Fig. 7a and b, respectively, versus growth temperature for samples B, C, D, and E. It is clear that both the growth rate and incorporation of indium decrease when the growth temperature increases. Analyzing the data using an Arrheniustype dependence does not yield a common activation energy for the growth rate and In incorporation. Several models have been developed to describe incorporation of group III elements during MBE of III-nitrides.17,22–24 However, these models do not adequately describe our experimental results. One model developed for plasma-assisted MBE
828
Uddin, Pandikunta, Mansurov, Sohal, Myasishchev, Guryanov, Kuryatkov, Holtz, and Nikishin
Fig. 6. Optical bandgap versus composition of InAlN alloy. The line is a guide to the eye.
(PAMBE) of InAlN alloys assumes that the epitaxial growth rate can be represented as a sum of the growth rates of AlN and InN.17 Using this approach they determine the best conditions for growth of InAlN using PAMBE. This model takes into account In desorption, a factor which appears to be critical in the growth of InAlN. We have applied the same approach for analysis of our experimental results and found that that model17 qualitatively follows our observations. We conclude from this discussion that desorption of In from the growing InAlN surface is the main factor in both decreasing growth rate and indium incorporation with increasing growth temperature. Evidently, the competition between indium desorption17 and ammonia decomposition efficiency,25 which both increase with temperature, favors the former under our conditions, resulting in diminishing InAlN growth rate at higher growth temperature. We apply a straightforward approach to estimate the indium incorporation coefficient. Using the SIMS data and the growth rate for each sample we calculate the areal density of In incorporated in each alloy per unit time. In separate experiments, the In flux was determined using low-temperature deposition on a bare Si substrate, for which desorption is negligible. The total thickness of the resulting In film was measured ex situ using SEM cross-sections. From the deposition time we obtain the incident In flux. The ratio of these two quantities allows us to estimate the In incorporation coefficient. Results are illustrated in Fig. 7c. The degree of indium incorporation decreases from 9.5% at growth temperature of 590°C to only 3% at 660°C. These are considerably lower than what we calculate using the model in Ref. 17. This discrepancy may be because their model does not take into account the NH3 decomposition efficiency, which is a strong function of temperature25 in the range investigated by us. Previous work has reported the dependences of the growth rate and In incorporation into InGaN at
Fig. 7. InN concentration (a), growth rate (b), and In incorporation (c) in InAlN alloys versus growth temperature. Growth rates were determined from SEM cross-sections and InAlN growth time. (Color figure online).
different temperatures during GSMBE with ammonia.23 For these materials, In incorporation and growth rate are both found to decrease with increasing growth temperature, although the dependences are significantly different. They report the indium incorporation to decrease dramatically, by a factor of 3, while the growth rate drops by 10% to 15% in the 590°C to 660°C range. This was attributed to differences in the volatilities of In and Ga for different growth temperatures and alloy compositions. Further studies of the In evaporation rate from different InAlN alloy surfaces are planned to model the growth of InAlN. CONCLUSIONS GSMBE with ammonia has been used to grow InAlN alloys on Si(111) substrates. The alloys were grown at temperatures ranging from 580°C to 660°C immediately following the growth of an AlN buffer layer at 920°C. Compositions were determined to be in the range of 2.5 mol.% to 8.0 mol.% InN using XRD. SIMS depth profiles confirmed these concentrations, and showed an interesting increase in In content near the surface for all samples studied. By checking the relationship with oxygen content, we believe the observed rise to be related to formation of In oxide. For the sample grown at temperature
InAlN on Si(111)
829
580°C, the SIMS profiles showed a pronounced gradient. We conclude from this that temperatures above this value are needed in GSMBE to obtain compositional uniformity. Bandgaps are in agreement with previously published work. We find that both the growth rate and In incorporation are strong functions of growth temperature, under our conditions, both decreasing substantially from 590°C to 660°C. This dependence can be qualitatively described using the model of Ref. 17. However, our In incorporation coefficients are significantly smaller than what we predict when applying their model. We attribute this discrepancy to the temperature sensitivity of NH3 decomposition in our range of growth temperatures. ACKNOWLEDGEMENTS This work was supported by the U.S. National Science Foundation (ECS-0609416) and the State of Texas NHARP (Contract No. 003644-0042-2009). The authors acknowledge V. A. Elyukhin and S. Yu. Karpov for useful discussions.
REFERENCES 1. M. Higashiwaki and T. Matsui, Jpn. J. Appl. Phys. 43, L768 (2004). 2. J. Xue, Y. Hao, J. Zhang, X. Zhou, Z. Liu, J. Ma, and Z. Lin, Appl. Phys. Lett. 98, 113504 (2011). 3. A. Castiglia, E. Feltin, G. Cosendey, A. Altoukhov, J.F. Carlin, R. Butte´, and N. Grandjean, Appl. Phys. Lett. 94, 193506 (2009). 4. S. Choi, H.J. Kim, S.S. Kim, J. Liu, J. Kim, J.H. Ryou, R.D. Dupuis, A.M. Fischer, and F.A. Ponce, Appl. Phys. Lett. 96, 221105 (2010). 5. J.F. Carlin, J. Dorsaz, E. Feltin, R. Butte´, N. Grandjean, M. Ilegems, and M. Lau¨gt, Appl. Phys. Lett. 86, 031107 (2005). 6. H. He, Y. Cao, R. Fu, H. Wang, J. Huang, C. Huang, M. Wang, and Z. Deng, J. Mater. Sci. Mater. Electron. 21, 676 (2010).
7. W. Terashima, S.B. Che, Y. Ishitani, and A. Yoshikawa, Jpn. J. Appl. Phys. 45, L539 (2006). 8. S.L. Sahonta, G.P. Dimitrakopulos, T. Kehagias, J. Kioseoglou, A. Adikimenakis, E. Ilipoulos, A. Georgakilas, H. Kirmse, W. Neumann, and P. Komninou, Appl. Phys. Lett. 95, 021913 (2009). 9. T.T. Kang, M. Yamamoto, M. Tanaka, A. Hashimoto, and A. Yamamoto, J. Appl. Phys. 106, 053525 (2009). 10. V. Darakchieva, M. Beckers, M.Y. Xie, L. Hultman, B. Monemar, J.F. Carlin, E. Feltin, M. Gonschorek, and N. Grandjean, J. Appl. Phys. 103, 103513 (2008). 11. C. Hums, J. Bla¨sing, A. Dadgar, A. Diez, T. Hempel, J. Christen, A. Krost, K. Lorenz, and E. Alves, Appl. Phys. Lett. 90, 022105 (2007). 12. Q. Han, C. Duan, G. Du, W. Shi, and L. Ji, J. Electron. Mater. 39, 5 (2010). 13. S. Iwata, Y. Nanjo, T. Okuno, S. Kurai, and T. Taguchi, Jpn. J. Appl. Phys. 46, 3394 (2007). 14. G. Guryanov, T.P.S. Clair, R. Bhat, C. Caneau, S. Nikishin, B. Borisov, and A. Budrevich, Appl. Surf. Sci. 252, 7208 (2006). 15. S.A. Nikishin, V.G. Antipov, S. Francoeur, N.N. Faleev, G.A. Seryogin, V.A. Elyukhin, H. Temkin, T.I. Prokofyeva, M. Holtz, A. Konkar, and S. Zollner, Appl. Phys. Lett. 75, 4 (1999). 16. V. Mansurov, X. Xu, M. Pandikunta, R. Uddin, and S. Nikishin, J. Vac. Sci. Technol. B 29, 03C129 (2011). 17. Z.Y. Wang, B.M. Shi, Y. Cai, N. Wang, and M.H. Xie, J. Appl. Phys. 108, 033503 (2010). 18. S. Zhang, B. Liu, J.Y. Yin, H.H. Sun, Z.H. Feng, and L.C. Zhao, J. Phys. D Appl. Phys. 44, 075405 (2011). 19. M. Holtz, T. Prokofyeva, M. Seon, K. Copeland, J. Vanbuskirk, S. Williams, S.A. Nikishin, V. Tretyakov, and H. Temkin, J. Appl. Phys. 89, 7977 (2001). 20. K.S. Kim, A. Saxler, P. Kung, M. Razeghi, and K.Y. Lim, Appl. Phys. Lett. 71, 800 (1997). 21. R.R. Pela´, C. Caetano, M. Marques, L.G. Ferreira, J. Furthmu¨ller, and L.K. Tales, Appl. Phys. Lett. 98, 151907 (2011). 22. V.A. Elyukhin, G.G. Salgado, R.P. Sierra, and S.A. Nikishin, J. Appl. Phys. 93, 5185 (2003). 23. S.Yu. Karpov, Yu.N. Makarov, and M.S. Ramm, MRS Internet J. Nitride Semicond. Res. 2, article 45 (1997). 24. S.T. Liu, X.Q. Wang, G. Chen, Y.W. Zhang, L. Feng, C.C. Huang, F.J. Xu, N. Tang, L.W. Sang, M. Sumiya, and B. Shen, J. Appl. Phys. 110, 113514 (2011). 25. M. Mesrine, N. Grandjean, and J. Massies, Appl. Phys. Lett. 72, 350 (1998).
