- Email: [email protected]
Growth of GaNAs films by molecular beam epitaxy
Journal of Crystal Growth 227–228 (2001) 486–490
Growth of GaNAs films by molecular beam epitaxy C.T. Foxona,*, S.V. Novikova, R.P. Campiona, C.S. Davisa, T.S. Chenga, A.J. Winsera,b, I. Harrisonb b
a School of Physics and Astronomy, University of Nottingham, Nottingham NG7 2RD, UK School of Electrical and Electronic Engineering, University of Nottingham, Nottingham NG7 2RD, UK
Abstract We have investigated arsenic doping of GaN films grown at high temperature (8008C) by molecular beam epitaxy. Growth in an arsenic environment changes the growth mode and promotes the growth of the cubic polytype. Growth under arsenic also produces films which show strong blue emission at room temperature, whose intensity is dependent on both the As flux and the Ga : N ratio. We propose a growth model based on As incorporation onto the Ga sublattice, which can explain all of our observed data. # 2001 Elsevier Science B.V. All rights reserved. PACS: 81.15.Hi; 81.05.Ea Keywords: A1. Growth models; A3. Molecular beam epitaxy; B1. Gallium compounds; B1. Nitrides; B2. Semiconducting III–V materials
1. Introduction Group III-Nitrides (AlN, GaN, InN and their mixed group III alloys) are used increasingly for amber, green and blue light emitting diodes; as the basis for white light sources; as laser diodes for the blue/UV and also for high-power, high-frequency electronic devices. At present, the group IIINitride layers are grown predominantly by metal-organic vapour phase epitaxy (MOVPE), but molecular beam epitaxy (MBE) offers the potential advantages of both in situ monitoring using e.g. reflection high energy electron diffraction (RHEED) and of lower temperature growth. The *Corresponding author. Tel.: +44-115-951-5164; fax: +44115-951-5180. E-mail address: [email protected] (C.T. Foxon).
latter may help to avoid the problem of phase separation in the alloys, which are the active part of any laser diode. Recently, the mixed group V alloys GaNxAs1 x have been grown both by MBE and MOVPE. Because of the unusually large bowing parameter, significant changes in band gap can be obtained with only small changes in composition. Due to a miscibility gap, it has not been possible to obtain alloys over the whole composition range. GaNxAs1 x alloys (solid solutions) have been obtained with an As concentration of 1 x50.01 on the GaN rich side [1,2] and with a N concentration x of 50.15 on the GaAs rich side [3,4]. For GaAs rich alloys, the maximum N concentration increases with decreasing growth temperatures [4,5], but for the GaN rich alloys, the temperature dependence of As incorporation has not been determined.
0022-0248/01/$ - see front matter # 2001 Elsevier Science B.V. All rights reserved. PII: S 0 0 2 2 - 0 2 4 8 ( 0 1 ) 0 0 7 5 3 - 9
C.T. Foxon et al. / Journal of Crystal Growth 227–228 (2001) 486–490
Some reports show that doping GaN with As improves both the surface morphology [6] and the electronic properties [7,8]; improving the electron mobility and decreasing the yellow luminescence. For As-doped GaN films grown by MOVPE at high temperature, a characteristic blue emission at 2.6 eV has been observed [9]. We have recently demonstrated that strong blue emission can be obtained in MBE grown As-doped GaN [10–12]. GaNAs alloys may therefore be a possible alternative to InGaN in optical device structures, however this has not so far been demonstrated. GaN layers usually grow as a hexagonal (wurtzite) polytype, but under some conditions the cubic (zinc-blende) phase can be obtained. In our previous investigations of GaNAs, we observed that at high MBE growth temperatures with a high As flux present, pure zinc-blende GaN can be grown with no hexagonal GaN or GaAs inclusions [13,14]. More recently, it has been shown that As acts as a surfactant giving rise to specific surface reconstructions during the growth by MBE [15–17], which may be due to a high As surface concentration, but this has not been demonstrated so far. In this paper, we report on the surface, structural and optical properties of GaN rich GaNAs layers grown by MBE. We have studied the influence of the As flux on the properties of GaNAs layers grown with both arsenic dimers and tetramers. Our results suggest a model for the growth and we propose that the blue emission arises as a result of As present on the Ga sub-lattice.
2. Experimental details The GaNAs films were grown by MBE on Mo coated sapphire substrates at approximately 8008C in a reactor, which has been described elsewhere [18]. Active nitrogen for the growth was obtained from an Oxford Applied Research CARS25 RF plasma source. A conventional effusion source was used for Ga and arsenic dimers or tetramers were generated using a purpose made two-zone effusion source. Before film growth, the sapphire substrates were nitrided at the growth temperature for 30 min. During nitridation and growth, the plasma
487
source was operated at 450 W with nitrogen flow rate of a few standard cubic centimetres per minute (sccm) resulting in a system pressure of 1 2 10 5 Torr. The Ga and N fluxes were initially set to give stoichiometric conditions for the GaN and are measured using an ion gauge placed in the sample position. The temperature of the substrate was measured both by a thermocouple and a pyrometer, for the latter an emissivity of 0.14 is used. In situ RHEED patterns prior to, during and after growth were monitored using a VG LEG110 electron gun operating at 12 keV. After growth the films were studied using Auger electron spectroscopy to determine the chemical composition, ex situ electron and X-ray diffraction to study their structural properties, atomic force microscopy to show surface morphology and photoluminescence to determine their optical properties.
3. Results and discussion At a growth temperature of 8008C, with the Ga and N fluxes adjusted to give stoichiometric growth, we achieve a growth rate of approximately 0.3 mm/h, limited by the arrival rate of active nitrogen. At this temperature, except for the region at the very edge of the samples where the temperature is lower, any excess Ga is desorbed. X-ray studies of such films show the familiar [0 0 0 2] growth direction for wurtzite GaN, with the usual rotation of GaN with respect to sapphire of 308. Reciprocal lattice maps of such films show the usual mosaic spread, with clear evidence for some cubic material, probably at the sapphire/ GaN interface. In situ RHEED studies of the films grown at about 800oC show significant differences with increasing arsenic flux. In the absence of As, during growth a (1 1) pattern is observed, which changes to a (4 4) pattern on cooling. For all As fluxes, a clear (2 2) pattern is observed about 30 s after termination of growth. However, in this situation a background pressure of As is still present during this time. At small As fluxes, mixed 2 and 3 patterns are observed on cooling indicating mixed domains on the surface. How-
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ever, at high As fluxes only a (2 2) pattern is observed. This reconstruction is stable over a wide range of temperatures and can be maintained for several days unlike the situation for the growth of GaN without As. Auger electron spectroscopy (AES) studies of the films after growth show the presence of a small amount of As, which is almost independent of the As flux. Atomic force microscopy (AFM) studies show clear evidence for improved surface morphology for a range of As fluxes. The low temperature PL spectra from samples grown without As show the usual donor bound exciton peak. At room temperature, however, free exciton emission is observed as shown in Fig. 1. With an As2 flux, a quite different PL spectrum is measured as also shown in Fig. 1. Firstly, very strong blue emission is observed at room temperature, the integrated intensity being more than a decade stronger than the band edge emission without As, secondly the appearance of additional band edge emission from cubic GaN. Very similar spectra are observed with As4, however, the dependence of peak intensity on the dimer and tetramer fluxes is not identical, as
shown in Fig. 2 and as we have previously reported in more detail [10–12]. The blue emission consists of several overlapping transitions whose energies increase slightly with increasing As2 flux [10–12]. The transition energies correspond to 2.3–2.45 eV, 2.55–2.7 eV and 2.75–2.9 eV, the second peak being generally the most intense. The results suggest that there may be at least two defect transitions involving As. The intensity of the cubic band edge emission shows different power law dependence for As2 and As4, as shown in Fig. 2 on a log–log plot. We have previously proposed a simple model [14], which will predict the behaviour we observe for the formation of cubic GaN. In this model, cubic GaAs is first formed by the reaction of Ga and As on the surface and then the As is replaced by N, due to the higher bond strength of GaN compared to GaAs. For this simple model, we might expect the intensity of cubic emission to increase as As0.5 2 and as As0.25 4 . From Fig. 2 we see that this is approximately true for the dimer, but the slope for the tetramer is greater than predicted. Nevertheless this gives us reasonable confidence that
Fig. 1. Room temperature PL from GaN grown at 8008C under stoichiometric conditions and from GaN grown at 8008C, with an additional As flux of 1.3 10 5 mbar (beam equivalent pressure). The peaks at 3.41 and 3.22 correspond to free exciton emission from wurtzite and zinc-blende GaN respectively, the broad emission at 2.6 eV corresponds to As-doped GaN, which may result from As on the Ga sublattice.
C.T. Foxon et al. / Journal of Crystal Growth 227–228 (2001) 486–490
489
Fig. 2. The dependence of PL intensity on As flux for cubic GaN band edge and blue emission.
Fig. 3. The dependence of blue emission intensity on the OED for fixed Ga and As fluxes.
the basic model is correct. In our previous studies of the growth of cubic GaN induced by As, (0 0 1) oriented GaAs and GaP substrates were used with the resulting cubic GaN also growing in the [0 0 1] direction. In the present study, the cubic phase GaN has the [1 1 1] direction parallel to the [0 0 0 1] directions of sapphire and wurtzite GaN. As a result, from
simple y–2y X-ray studies we cannot distinguish cubic from hexagonal material, however, from off-axis measurements using appropriate reflections, clear evidence for the cubic phase can be obtained. The data will be discussed elsewhere in detail. More recently, we have been studying the dependence of blue emission intensity on
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C.T. Foxon et al. / Journal of Crystal Growth 227–228 (2001) 486–490
the N : Ga ratio for a fixed flux of As [12]. Since excess Ga is present, the growth rate and hence film thickness increase slightly with increasing N flow rate, the growth rate being linearly dependent on the optical emission from the plasma as we have previously observed [19]. The intensity of the blue emission clearly increases with increasing N, as shown in Fig. 3. We interpret this as strong evidence for the blue emission arising from As incorporated onto a Ga site. With increasing N : Ga ratio, we expect the density of Ga vacancies to increase, thus increasing the probability of incorporation of As onto the Ga sub-lattice. Such behaviour has been predicted theoretically [20] and our data provides strong support for that model.
4. Conclusions Arsenic doping of GaN films grown at high temperature produces two independent effects. First there is a change in the growth mode, part of the film grows as the cubic polytype. This can be explained using our previously proposed model for the growth of cubic GaN induced by As [13]. More significantly, strong blue emission is observed at room temperature whose intensity is dependent on both the As flux and the Ga : N ratio. This observation can be explained by the presence of As on the Ga sublattice in agreement with previous theoretical predictions [20]. The above data suggests that As-doped GaN may be a suitable alternative to (InGa)N for the growth of lasers by MBE.
Acknowledgements This work was supported by grants from the EPSRC (GR/M67438) and the Royal Society of UK (UK–China Joint Project Q744).
References [1] K. Iwata, H. Asahi, K. Asami, R. Kuroiwa, S. Gonda, Jpn. J. Appl. Phys. 37 (1998) 1436. [2] R. Kuroiwa, H. Asahi, K. Asami, S.J. Kim, K. Iwata, S. Gonda, Appl. Phys. Lett. 73 (1998) 2630. [3] M. Kondow, K. Uomi, A. Niwa, T. Kitatani, S. Watahiki, Y. Yazawa, K. Hosomi, T. Mozume, Solid-State Electron. 41 (1997) 209. [4] W.G. Bi, C.W. Tu, Appl. Phys. Lett. 70 (1997) 1608. [5] K. Uesugi, I. Suemune, Jpn. J. Appl. Phys. 36 (1997) L1572. [6] G. Mendoza Diaz, K.S. Stevens, A.F. Schwartzman, R. Beresford, J. Crystal Growth 178 (1997) 45. [7] L.J. Guido, P. Mitev, M. Gherasimova, B. Gaffey, Appl. Phys. Lett. 72 (1998) 2005. [8] S.R. Jin, M. Ramsteiner, H.T. Grahn, K.H. Ploog, Z.Q. Zhu, D.X. Shen, A.Z. Li, P. Metev, L.J. Guido, J. Crystal Growth 212 (2000) 56. [9] X. Li, S. Kim, E.E. Reuter, S.G. Bishop, J.J. Coleman, Appl. Phys. Lett. 72 (1998) 1990. [10] A.J. Winser, S.V. Novikov, C.S. Davis, T.S. Cheng, C.T. Foxon, I. Harrison, Appl. Phys. Lett. 77 (2000) 2506. [11] C.T. Foxon, S.V. Novikov, T.S. Cheng, C.S. Davis, R.P. Campion, A.J. Winser, I. Harrison, J. Crystal Growth 219 (2000) 327. [12] S.V. Novikov, A.J. Winser, I. Harrison, C.S. Davis, C.T. Foxon, Semicond. Sci. Technol. 16 (2001) 103. [13] T.S. Cheng, L.C. Jenkins, S.E. Hooper, C.T. Foxon, J.W. Orton, D.E. Lacklison, Appl. Phys. Lett. 66 (1995) 1509. [14] L.C. Jenkins, T.S. Cheng, C.T. Foxon, J.W. Orton, S.E. Hooper, S.V. Novikov, V.V. Tret’yakov, J. Vac. Sci. Technol. B 13 (1995) 1585. [15] G. Feuillet, H. Hamaguchi, K. Ohta, P. Hacke, H. Okumura, S. Yoshida, Appl. Phys. Lett. 70 (1997) 1025. [16] J. Neugebauer, T. Zywietz, M. Scheffler, J.E. Northrup, C.G. Van de Walle, Phys. Rev. Lett. 80 (1998) 3097. [17] V. Ramachandran, C.D. Lee, R.M. Feenstra, A.R. Smith, J.E. Northrup, D.W. Greve, J. Crystal Growth 209 (2000) 355. [18] C.T. Foxon, T.S. Cheng, S.V. Novikov, N.J. Jeffs, O.H. Hughes, Yu.V. Melnik, A.E. Nikolaev, V.A. Dmitriev, Surf. Sci. 421 (1999) 377. [19] A.V. Blant, O.H. Hughes, T.S. Cheng, S.V. Novikov, C.T. Foxon, Plasma Source Sci. Technol. 9 (2000) 12. [20] C.G. Van de Walle, J. Neugebauer, Appl. Phys. Lett. 76 (2000) 1009.
Growth of GaNAs films by molecular beam epitaxy C.T. Foxona,*, S.V. Novikova, R.P. Campiona, C.S. Davisa, T.S. Chenga, A.J. Winsera,b, I. Harrisonb b
a School of Physics and Astronomy, University of Nottingham, Nottingham NG7 2RD, UK School of Electrical and Electronic Engineering, University of Nottingham, Nottingham NG7 2RD, UK
Abstract We have investigated arsenic doping of GaN films grown at high temperature (8008C) by molecular beam epitaxy. Growth in an arsenic environment changes the growth mode and promotes the growth of the cubic polytype. Growth under arsenic also produces films which show strong blue emission at room temperature, whose intensity is dependent on both the As flux and the Ga : N ratio. We propose a growth model based on As incorporation onto the Ga sublattice, which can explain all of our observed data. # 2001 Elsevier Science B.V. All rights reserved. PACS: 81.15.Hi; 81.05.Ea Keywords: A1. Growth models; A3. Molecular beam epitaxy; B1. Gallium compounds; B1. Nitrides; B2. Semiconducting III–V materials
1. Introduction Group III-Nitrides (AlN, GaN, InN and their mixed group III alloys) are used increasingly for amber, green and blue light emitting diodes; as the basis for white light sources; as laser diodes for the blue/UV and also for high-power, high-frequency electronic devices. At present, the group IIINitride layers are grown predominantly by metal-organic vapour phase epitaxy (MOVPE), but molecular beam epitaxy (MBE) offers the potential advantages of both in situ monitoring using e.g. reflection high energy electron diffraction (RHEED) and of lower temperature growth. The *Corresponding author. Tel.: +44-115-951-5164; fax: +44115-951-5180. E-mail address: [email protected] (C.T. Foxon).
latter may help to avoid the problem of phase separation in the alloys, which are the active part of any laser diode. Recently, the mixed group V alloys GaNxAs1 x have been grown both by MBE and MOVPE. Because of the unusually large bowing parameter, significant changes in band gap can be obtained with only small changes in composition. Due to a miscibility gap, it has not been possible to obtain alloys over the whole composition range. GaNxAs1 x alloys (solid solutions) have been obtained with an As concentration of 1 x50.01 on the GaN rich side [1,2] and with a N concentration x of 50.15 on the GaAs rich side [3,4]. For GaAs rich alloys, the maximum N concentration increases with decreasing growth temperatures [4,5], but for the GaN rich alloys, the temperature dependence of As incorporation has not been determined.
0022-0248/01/$ - see front matter # 2001 Elsevier Science B.V. All rights reserved. PII: S 0 0 2 2 - 0 2 4 8 ( 0 1 ) 0 0 7 5 3 - 9
C.T. Foxon et al. / Journal of Crystal Growth 227–228 (2001) 486–490
Some reports show that doping GaN with As improves both the surface morphology [6] and the electronic properties [7,8]; improving the electron mobility and decreasing the yellow luminescence. For As-doped GaN films grown by MOVPE at high temperature, a characteristic blue emission at 2.6 eV has been observed [9]. We have recently demonstrated that strong blue emission can be obtained in MBE grown As-doped GaN [10–12]. GaNAs alloys may therefore be a possible alternative to InGaN in optical device structures, however this has not so far been demonstrated. GaN layers usually grow as a hexagonal (wurtzite) polytype, but under some conditions the cubic (zinc-blende) phase can be obtained. In our previous investigations of GaNAs, we observed that at high MBE growth temperatures with a high As flux present, pure zinc-blende GaN can be grown with no hexagonal GaN or GaAs inclusions [13,14]. More recently, it has been shown that As acts as a surfactant giving rise to specific surface reconstructions during the growth by MBE [15–17], which may be due to a high As surface concentration, but this has not been demonstrated so far. In this paper, we report on the surface, structural and optical properties of GaN rich GaNAs layers grown by MBE. We have studied the influence of the As flux on the properties of GaNAs layers grown with both arsenic dimers and tetramers. Our results suggest a model for the growth and we propose that the blue emission arises as a result of As present on the Ga sub-lattice.
2. Experimental details The GaNAs films were grown by MBE on Mo coated sapphire substrates at approximately 8008C in a reactor, which has been described elsewhere [18]. Active nitrogen for the growth was obtained from an Oxford Applied Research CARS25 RF plasma source. A conventional effusion source was used for Ga and arsenic dimers or tetramers were generated using a purpose made two-zone effusion source. Before film growth, the sapphire substrates were nitrided at the growth temperature for 30 min. During nitridation and growth, the plasma
487
source was operated at 450 W with nitrogen flow rate of a few standard cubic centimetres per minute (sccm) resulting in a system pressure of 1 2 10 5 Torr. The Ga and N fluxes were initially set to give stoichiometric conditions for the GaN and are measured using an ion gauge placed in the sample position. The temperature of the substrate was measured both by a thermocouple and a pyrometer, for the latter an emissivity of 0.14 is used. In situ RHEED patterns prior to, during and after growth were monitored using a VG LEG110 electron gun operating at 12 keV. After growth the films were studied using Auger electron spectroscopy to determine the chemical composition, ex situ electron and X-ray diffraction to study their structural properties, atomic force microscopy to show surface morphology and photoluminescence to determine their optical properties.
3. Results and discussion At a growth temperature of 8008C, with the Ga and N fluxes adjusted to give stoichiometric growth, we achieve a growth rate of approximately 0.3 mm/h, limited by the arrival rate of active nitrogen. At this temperature, except for the region at the very edge of the samples where the temperature is lower, any excess Ga is desorbed. X-ray studies of such films show the familiar [0 0 0 2] growth direction for wurtzite GaN, with the usual rotation of GaN with respect to sapphire of 308. Reciprocal lattice maps of such films show the usual mosaic spread, with clear evidence for some cubic material, probably at the sapphire/ GaN interface. In situ RHEED studies of the films grown at about 800oC show significant differences with increasing arsenic flux. In the absence of As, during growth a (1 1) pattern is observed, which changes to a (4 4) pattern on cooling. For all As fluxes, a clear (2 2) pattern is observed about 30 s after termination of growth. However, in this situation a background pressure of As is still present during this time. At small As fluxes, mixed 2 and 3 patterns are observed on cooling indicating mixed domains on the surface. How-
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C.T. Foxon et al. / Journal of Crystal Growth 227–228 (2001) 486–490
ever, at high As fluxes only a (2 2) pattern is observed. This reconstruction is stable over a wide range of temperatures and can be maintained for several days unlike the situation for the growth of GaN without As. Auger electron spectroscopy (AES) studies of the films after growth show the presence of a small amount of As, which is almost independent of the As flux. Atomic force microscopy (AFM) studies show clear evidence for improved surface morphology for a range of As fluxes. The low temperature PL spectra from samples grown without As show the usual donor bound exciton peak. At room temperature, however, free exciton emission is observed as shown in Fig. 1. With an As2 flux, a quite different PL spectrum is measured as also shown in Fig. 1. Firstly, very strong blue emission is observed at room temperature, the integrated intensity being more than a decade stronger than the band edge emission without As, secondly the appearance of additional band edge emission from cubic GaN. Very similar spectra are observed with As4, however, the dependence of peak intensity on the dimer and tetramer fluxes is not identical, as
shown in Fig. 2 and as we have previously reported in more detail [10–12]. The blue emission consists of several overlapping transitions whose energies increase slightly with increasing As2 flux [10–12]. The transition energies correspond to 2.3–2.45 eV, 2.55–2.7 eV and 2.75–2.9 eV, the second peak being generally the most intense. The results suggest that there may be at least two defect transitions involving As. The intensity of the cubic band edge emission shows different power law dependence for As2 and As4, as shown in Fig. 2 on a log–log plot. We have previously proposed a simple model [14], which will predict the behaviour we observe for the formation of cubic GaN. In this model, cubic GaAs is first formed by the reaction of Ga and As on the surface and then the As is replaced by N, due to the higher bond strength of GaN compared to GaAs. For this simple model, we might expect the intensity of cubic emission to increase as As0.5 2 and as As0.25 4 . From Fig. 2 we see that this is approximately true for the dimer, but the slope for the tetramer is greater than predicted. Nevertheless this gives us reasonable confidence that
Fig. 1. Room temperature PL from GaN grown at 8008C under stoichiometric conditions and from GaN grown at 8008C, with an additional As flux of 1.3 10 5 mbar (beam equivalent pressure). The peaks at 3.41 and 3.22 correspond to free exciton emission from wurtzite and zinc-blende GaN respectively, the broad emission at 2.6 eV corresponds to As-doped GaN, which may result from As on the Ga sublattice.
C.T. Foxon et al. / Journal of Crystal Growth 227–228 (2001) 486–490
489
Fig. 2. The dependence of PL intensity on As flux for cubic GaN band edge and blue emission.
Fig. 3. The dependence of blue emission intensity on the OED for fixed Ga and As fluxes.
the basic model is correct. In our previous studies of the growth of cubic GaN induced by As, (0 0 1) oriented GaAs and GaP substrates were used with the resulting cubic GaN also growing in the [0 0 1] direction. In the present study, the cubic phase GaN has the [1 1 1] direction parallel to the [0 0 0 1] directions of sapphire and wurtzite GaN. As a result, from
simple y–2y X-ray studies we cannot distinguish cubic from hexagonal material, however, from off-axis measurements using appropriate reflections, clear evidence for the cubic phase can be obtained. The data will be discussed elsewhere in detail. More recently, we have been studying the dependence of blue emission intensity on
490
C.T. Foxon et al. / Journal of Crystal Growth 227–228 (2001) 486–490
the N : Ga ratio for a fixed flux of As [12]. Since excess Ga is present, the growth rate and hence film thickness increase slightly with increasing N flow rate, the growth rate being linearly dependent on the optical emission from the plasma as we have previously observed [19]. The intensity of the blue emission clearly increases with increasing N, as shown in Fig. 3. We interpret this as strong evidence for the blue emission arising from As incorporated onto a Ga site. With increasing N : Ga ratio, we expect the density of Ga vacancies to increase, thus increasing the probability of incorporation of As onto the Ga sub-lattice. Such behaviour has been predicted theoretically [20] and our data provides strong support for that model.
4. Conclusions Arsenic doping of GaN films grown at high temperature produces two independent effects. First there is a change in the growth mode, part of the film grows as the cubic polytype. This can be explained using our previously proposed model for the growth of cubic GaN induced by As [13]. More significantly, strong blue emission is observed at room temperature whose intensity is dependent on both the As flux and the Ga : N ratio. This observation can be explained by the presence of As on the Ga sublattice in agreement with previous theoretical predictions [20]. The above data suggests that As-doped GaN may be a suitable alternative to (InGa)N for the growth of lasers by MBE.
Acknowledgements This work was supported by grants from the EPSRC (GR/M67438) and the Royal Society of UK (UK–China Joint Project Q744).
References [1] K. Iwata, H. Asahi, K. Asami, R. Kuroiwa, S. Gonda, Jpn. J. Appl. Phys. 37 (1998) 1436. [2] R. Kuroiwa, H. Asahi, K. Asami, S.J. Kim, K. Iwata, S. Gonda, Appl. Phys. Lett. 73 (1998) 2630. [3] M. Kondow, K. Uomi, A. Niwa, T. Kitatani, S. Watahiki, Y. Yazawa, K. Hosomi, T. Mozume, Solid-State Electron. 41 (1997) 209. [4] W.G. Bi, C.W. Tu, Appl. Phys. Lett. 70 (1997) 1608. [5] K. Uesugi, I. Suemune, Jpn. J. Appl. Phys. 36 (1997) L1572. [6] G. Mendoza Diaz, K.S. Stevens, A.F. Schwartzman, R. Beresford, J. Crystal Growth 178 (1997) 45. [7] L.J. Guido, P. Mitev, M. Gherasimova, B. Gaffey, Appl. Phys. Lett. 72 (1998) 2005. [8] S.R. Jin, M. Ramsteiner, H.T. Grahn, K.H. Ploog, Z.Q. Zhu, D.X. Shen, A.Z. Li, P. Metev, L.J. Guido, J. Crystal Growth 212 (2000) 56. [9] X. Li, S. Kim, E.E. Reuter, S.G. Bishop, J.J. Coleman, Appl. Phys. Lett. 72 (1998) 1990. [10] A.J. Winser, S.V. Novikov, C.S. Davis, T.S. Cheng, C.T. Foxon, I. Harrison, Appl. Phys. Lett. 77 (2000) 2506. [11] C.T. Foxon, S.V. Novikov, T.S. Cheng, C.S. Davis, R.P. Campion, A.J. Winser, I. Harrison, J. Crystal Growth 219 (2000) 327. [12] S.V. Novikov, A.J. Winser, I. Harrison, C.S. Davis, C.T. Foxon, Semicond. Sci. Technol. 16 (2001) 103. [13] T.S. Cheng, L.C. Jenkins, S.E. Hooper, C.T. Foxon, J.W. Orton, D.E. Lacklison, Appl. Phys. Lett. 66 (1995) 1509. [14] L.C. Jenkins, T.S. Cheng, C.T. Foxon, J.W. Orton, S.E. Hooper, S.V. Novikov, V.V. Tret’yakov, J. Vac. Sci. Technol. B 13 (1995) 1585. [15] G. Feuillet, H. Hamaguchi, K. Ohta, P. Hacke, H. Okumura, S. Yoshida, Appl. Phys. Lett. 70 (1997) 1025. [16] J. Neugebauer, T. Zywietz, M. Scheffler, J.E. Northrup, C.G. Van de Walle, Phys. Rev. Lett. 80 (1998) 3097. [17] V. Ramachandran, C.D. Lee, R.M. Feenstra, A.R. Smith, J.E. Northrup, D.W. Greve, J. Crystal Growth 209 (2000) 355. [18] C.T. Foxon, T.S. Cheng, S.V. Novikov, N.J. Jeffs, O.H. Hughes, Yu.V. Melnik, A.E. Nikolaev, V.A. Dmitriev, Surf. Sci. 421 (1999) 377. [19] A.V. Blant, O.H. Hughes, T.S. Cheng, S.V. Novikov, C.T. Foxon, Plasma Source Sci. Technol. 9 (2000) 12. [20] C.G. Van de Walle, J. Neugebauer, Appl. Phys. Lett. 76 (2000) 1009.